Thomas Carlier

Marie-Hélène Chambrier

Antonio Da Costa

Florent Blanchard

Thibaud Denneulin

Manon Létiche

Pascal Roussel

Rachel Desfeux

Anthony Ferri

Ferroelectric state in a α-Nd 2 WO 6 polymorph stabilized in thin film

)-oriented crystallites with respect to the plane of the substrate mainly at the interface film/substrate and dispersed in the (001)-NdWO matrix. In addition, robust piezoelectricity and ferroelectricity are revealed at room temperature through both local hysteresis loops and domain manipulation experiments using piezoresponse force microscopy technique. Typical polarization retention behaviour associated to specific nanoscale conduction are in good agreement with classical ferroelectric phenomenon in oxide materials. The successful observation of piezo-/ferroelectricity at room temperature in innovative strain-stabilized α-NdWO thin films paves the way for new lead-free functional materials devoted to numerous applications, including actuators, sensors, or non-volatile memory devices.

INTRODUCTION

Since World War II, advanced electronic devices based on piezoelectric and ferroelectric materials attracted a lot of attention and represent an important part of industrial-business and innovations, taking in mind the fact that piezo-/ferroelectric applications areas are wide ranging from consumer devices to high technology industries: laptops, sonars, frequency filters, gas sensors, data storage systems (Ferroelectric Random Access Memories (Fe-RAMs)), etc. Moreover, promising and impending technological developments such as energy harvesting, geothermal power plant sensors, aerospace, miniaturization still keep piezo-/ferroelectric materials highly attractive. Since the 40's, lead zirconate-titanate Pb(Zr,Ti)O3 (PZT), [START_REF] Nagarajan | Dynamics of Ferroelastic Domains in Ferroelectric Thin Films[END_REF][START_REF] Vrejoiu | Intrinsic Ferroelectric Properties of Strained Tetragonal PbZr0.2Ti0.8O3 Obtained on Layer-by-Layer Grown, Defect-Free Single Crystalline Films[END_REF] (Pb,La)(Zr,Ti)O3 (PLZT) [START_REF] Leclerc | Elaboration of (111)-oriented La-doped PZT Thin Films on Platinized Silicon Substrates[END_REF] or xPb(MgNb)O3-(1-x)PbTiO3 (PMN-PT) [START_REF] Detalle | Nanoscale Study by Piezoresponse Force Microscopy of Relaxor 0.7Pb(Mg1/3Nb2/3)O3-0.3PbTiO3 and 0.9Pb(Mg1/3Nb2/3)O3-0.1PbTiO3 Thin Films Grown on Platinum and LaNiO3 Electrodes[END_REF] to cite the main ones of the well-known lead-based perovskite type family have been in the first place on the market of such electromechanical materials due to exceptional piezoelectric coefficient and polarization at room temperature. [START_REF] Saito | Lead-free piezoceramics[END_REF][START_REF] Shvartsman | Nonlinear local piezoelectric deformation in ferroelectric thin films studied by scanning force microscopy[END_REF][START_REF] Zhang | Lead-free piezoelectric ceramics vs. PZT?[END_REF] It is interesting to note that since the 70's, plenty of patents for Pb-containing relaxor based on multi-layer piezo-/ferroelectric ceramic capacitor devices have been filed. [START_REF] Uchino | Ferroelectric devices[END_REF] However, in reference to the European Union legislations, REACH (Registration, Evaluation, Authorization and Restriction of Chemicals) classification, connected to the environmental concerns, has decreed that even if leadbased substances are still allowed for the moment, they have to be progressively replaced by less toxic substances. [9, 10] In this context, new investigations have been conducted, mainly based on two strategies: either optimize wellknown piezoelectric materials such as BaTiO3 (BTO) [START_REF] Rödel | Perspective on the development of Lead-Free Piezoceramics[END_REF] and KxNa1-xNbO3 (KNN) [START_REF] Wu | Potassium-Sodium Niobate Lead-Free Piezoelectric Materials: Past, Present, and Future of Phase Boundaries[END_REF] , or explore new systems. Our approach belongs to this last option with the aim to discover new lead-free piezo-/ferroelectric compounds.

In the present work, we focus on lanthanide tungstate Ln2WO6 systems for different well-targeted reasons: i) some polymorphs are noncentro-symmetric and therefore susceptible to present ferroelectricity, ii) some tungsten bronze KxW2O6 are already known to form a ferroelectric class and also iii) dozen of papers deals with Bi2WO6 compound and its ferroelectric properties.

Even if studies of A2O3-MO3 (A = Ln 3+ , Bi 3+ , Y 3+ , Fe 3+ , Sc 3+ and M = W 6+ , Mo 6+ ) phase diagram were intensive in 70's in order to determine their structure, nowadays some of these phase diagrams still remain incomplete in terms of characterization properties. Concerning piezo-/ferroelectric properties, only a few works and even patents have been written for A2M3O12 (A = Gd, Tb and Eu , Y; M = W and Mo). [START_REF] Borchardt | Ferroelectric Rare Earth Molybdates[END_REF][START_REF] Flippen | Mechanical transducers using coupled ferroelectric-ferroelastic crystals[END_REF] On the other hand, Ln2O3-WO3 systems are interesting because of the low cost of both lanthanide and tungsten elements, combined to the low toxicity of these elements, which anticipates both European decrees as previously reported and industrial constraints. [9; 10] In terms of piezoelectric properties, lanthanide elements are stimulating owing their high polarizability (5.01 Å [START_REF] Leclerc | Elaboration of (111)-oriented La-doped PZT Thin Films on Platinized Silicon Substrates[END_REF] for neodymium for instance). [START_REF] Lu | Tunneluing Hot Spots in Ferroelectrics SrTiO3[END_REF] Concerning Ln2WO6 compounds, depending on the lanthanide radius, three different tungstate structures are observed according to the temperature and the exact method of synthetic preparation. [START_REF] Brixner | Refined Cell parameters of the Ln2WO6-Type Rare Earth Tunsgates[END_REF][START_REF] Efremov | Polymorphism of Oxytungstates Tr2WO6. Mechanism of Structural Changes of Er2WO6[END_REF][START_REF] Efremov | Polymorphism of Oxytungstates Tr2WO6. Analysis of Structural Type II (Gd2WO6 and Gd2MoO6). Mechanism of Structural Change in Gd2WO6 in the Phase Transition II <-> V[END_REF][START_REF] Polyanskaya | A New Form of the Scheelite Structural Type: Crystal Structure of Nd2WO6[END_REF][START_REF] Tyulin | Polymorphism of Oxytungstates Tr2WO6. Mechanisms of Structural Changes in Y2WO6[END_REF][START_REF] Borchardt | Rare-Earth Tungstates and 1:1 Oxytungstates[END_REF][START_REF] Efremov | The Crystal Structure of Monoclinic Y2WO6 and Yb2WO6[END_REF][START_REF] Efremov | Characteristic features of the crystal chemistry of lanthanide molybdates and tungstates[END_REF] La2WO6 presents two polymorphic structures depending on the temperature: i) the low temperature form, β-La2WO6, crystallizes in the space group P212121 (n°19) with cell parameters a = 7.5196(1) Å, b = 10.3476(1) Å, c = 12.7944(2) Å, [START_REF] Chambrier | Ab-initio Structure Determination of β-La2WO6[END_REF] ii) the high temperature form, [START_REF] Allix | Synthesis and Structure Determination of the High Temperature Form of La2WO6[END_REF] α-La2WO6, crystallizes in the space group Pm21n (n°31) with cell parameters a = 16.5531(1) Å, b = 5.5200(1) Å, c = 8.8833(1) Å. A first order phase transition between beta and alpha occurs at 1440°C. The structure for the α-La2WO6 phase is closely related to the Scheelite one. Recently, piezoelectricity and ferroelectricity within this compound were evidenced by our group when stabilized in thin film form on SrTiO3 and LaAlO3 substrates. [START_REF] Carlier | Lead-Free α-La2WO6 Ferroelectric Thin Films[END_REF][START_REF] Carlier | Microstucture and local electrical investigation of lead-free α-La2WO6 ferroelectric thin films by piezoresponse force microscopy[END_REF] On the one hand, Ln2WO6 compounds with large cations (Ln = Ce to Dy) usually crystallize in a monoclinic fluorine type-superstructure (C 2 𝑐 ; aM ≈ 3aF; bM ≈ 2bF; cM ≈ cF and β ≈ 107.7°). [START_REF] Apostolov | Thermal Expansion of Ln6WO12 and Ln2WO6 -An In Situ Synchrotron Xray Diffraction Study[END_REF] On the other hand, for the smallest cations (Ln = Ho to Lu) it results in a P [START_REF] Zhijun | Structure refinement Lu2WO6 and luminescent properties of Eu 3+ , Pr 3+ doped Lu2WO6[END_REF][START_REF] Beaury | Investigation of a new structural type for Y2WO6[END_REF] At high temperature, all the Ln2WO6 compounds with Ln = Ce to Lu undergo a phase transition (>1500°C) to a tetragonal compound.

Concerning Nd2WO6, Yoshimura et al. have reported another metastable phase between the low and high temperature prototypes in addition to the phase transition close to 1430°C. [START_REF] Yoshimura | Polymorphism of R2O3-WO3 (R = Rare Earth) compounds at high temperatures[END_REF] The low temperature form crystallizes in the space group C 2 𝑐 (n°15) with cell parameters a = 5.55(1) Å, b = 11.38(1) Å, c = 15.86(1) Å, β = 107.7°. The structure consists of two strongly distorted polyhedra: i) tungsten atoms are surrounded by five oxygen atoms, which build up an isolated trigonal bipyramid [WO5] -4 ; ii) neodymium atoms are forming [LnO8] -10 units (Figure 1a). The "intermediate temperature" polymorph observed between 1430°C and 1450°C is metastable and has been described in the orthorhombic system with the space group P212121 (n°19) and cell parameters a = 5.536(1) Å, b = 9.231(1) Å, c = 10.170(1) Å. [START_REF] Brixner | Refined Cell parameters of the Ln2WO6-Type Rare Earth Tunsgates[END_REF][START_REF] Efremov | The structure of a new modification of Nd2SO6[END_REF] This allotropic variety can be described with distorted isolated [WO6] and two different types of polyhedra for the neodymium: a square-triangle polyhedron [Nd(1)O7] and a monocapped trigonal prism [Nd(2)O6] (Figure 1b). [START_REF] Meier | Th. Crystal structure of dineodymium(III) hexaoxotellurate(VI), Nd2TeO6[END_REF] Finally, the third polymorph at high temperature (T>1450°C), has been first indexed by Yoshimura et al. in a tetragonal cell closely related to an ordered fluorite (aT ≈ aF = 5.426Å; cT ≈ cF = 8.695Å) (Figure 1c). [START_REF] Yoshimura | Polymorphism of R2O3-WO3 (R = Rare Earth) compounds at high temperatures[END_REF][START_REF] Yoshimura | Phases de haute température des tungstates de lanthanides du type Ln2O3-WO3 (Ln = La[END_REF] Yoshimura et al. and Efremov have noticed that all Ln2WO6 compounds, [31; 23] have the same high temperature structural type even if low temperature polymorphs are different. Later, Tyulin et al. solved the structure of Nd1.2Lu0.8WO6 which is isostructural to the Ln2WO6 high temperature form; [START_REF] Tyulin | Polymorphism of Ln2WO6 oxitungstates[END_REF] it crystallizes in the space group P4 ̅ 21m (n°113) with cell parameters a = 5.276(1) Å; c = 8.658(1) Å. The structure consists of large isolated [WO6] octahedra (figure 1c), while the neodymium atoms are in cubic coordination.

In our work, we focus on the growth of Nd2WO6 thin film by means of Pulsed Laser Deposition (PLD) on (001)oriented SrTiO3 (STO) substrate. Advanced microstructural characterizations carried out by X-Ray Diffraction (XRD) and Transmission Electron Microscopy (TEM) reveal the stabilization of a new Nd2WO6 polymorph, hereafter labelled alpha. Nanoscale ferroelectricity is highlighted in this new oxide thin film by carefully investigating electrical properties using Atomic Force Microscopy tools and, more specifically, the Piezoresponse Force Microscopy (PFM) and Conductive-AFM (c-AFM) modes.

EXPERIMENTAL METHODS

A Nd2WO6 ceramic pellet, 1 inch in diameter and 5 mm-thick, was used as target for laser ablation. The ceramic was prepared by solid-state reaction route method, mixing dried oxides Nd2O3 and WO3 (Sigma Adrich, purities of 99.9%) weighted in stoichiometric proportions. Twice, with an intermediate grinding, the mixture was heated at 1350°C in air for one night. The resulting powder was then ground, pressed in the form of a disc using an uni-axial press, sintered up to 1450°C for 12h in air and slowly cooled down to room temperature. The XRD pattern confirms the synthesis of a pure powder of the low temperature polymorph of Nd2WO6 that was afterward used as target. [START_REF] Brixner | Refined Cell parameters of the Ln2WO6-Type Rare Earth Tunsgates[END_REF] The depositions as thin film form by PLD were carried out with a Compex Pro 102 KrF excimer laser ( = 248 nm) on (001)-oriented STO single-crystal substrates (Crystal GmbH, Germany). The fluence was adjusted to 2 J/cm² and the deposition rate was 2 Hz. The target-substrate distance was adjusted to d = 4.5 cm. All substrates were first ultrasonically cleaned in ethanol for 5 minutes, and then dried. The pressure in the pulsed-laser chamber was lowered down to 10 -4 Pa. Then, the substrate temperature was raised through a 10°C/min step elevation to the growth temperature, i.e. 900°C under a dynamical O2 pressure of 10 -1 mbar, while the deposition corresponds to 12000 laser-pulse shots corresponding to a film thickness around 200 nm. Finally, O2 gas was introduced into the chamber (200 mBar) and the film was slowly cooled down to room temperature.

The bulk-powder X-ray diffractograms were obtained using a Rigaku Ultima IV equipped with a Cu-anode Xray tube (λKα = 1.5418 Å). The thin films crystalline nature (indexation, rocking curves and Reciprocal Space Maps (RSMs)) were characterized using a Rigaku SmartLab X-ray diffractometer equipped with a 9 kW rotating anode X-ray generator and a copper anode. For 2θ-ω and rocking curves analysis, -scan and RSMs, a high resolution configuration was used: the X-Ray beam was made parallel with a cross-beam optics and monochromatized with a double Ge (220) monochromator. In this study all the (2θ-ω) scans were performed in the range of 10-90°, with a step size of 0.02° and a 1°/min speed whereas the Φ-scan was acquired in the range of -10-360°, with a step size of 0.4° and a 40°/min speed. Pole figure measurements were made in medium resolution configuration, i.e. without the monochromator (λKα = 1.5418 Å) using a 0.5° grazing incidence and an in-plane geometry.

The TEM images were obtained using a Hitachi HF-3300 equipped with a cold field emission gun operated at 300 kV, an image aberration corrector (CEOS B-COR) and a 4k×4k camera (Gatan Oneview). Cross-section electron transparent thin foils were prepared using a dual beam machine (scanning electron microscope and focused gallium ion beam) FEI Helios. High resolution TEM images were processed using the Digital Micrograph software (Gatan). A color-coded map was created to show the distribution of two domain orientations in the film. A Fourier transform of the HRTEM image was first calculated. Circular masks were applied to isolate spots corresponding to a single domain orientation and the inverse Fourier transform was then calculated. The same operation was repeated twice for the second domain orientation and for the substrate. The three complex images obtained after inverse Fourier transform were used to create a RGB composite image where the modulus of each image is affected to a different color. Lattice deformation maps were calculated from the HRTEM image using geometrical phase analysis. [START_REF] Hÿtch | Quantitative measurement of displacement and strain fields from HREM micrographs[END_REF] Fourier transform operations and masks in Fourier space were used to reconstruct two phase images corresponding to two sets of non-colinear reciprocal lattice vectors. Horizontal (in-plane), vertical (out-of-plane) and shear deformation maps were calculated from the gradient of the phase images. The local deformation E(x,y) at a position (x,y) in the map is defined E(x,y)=(d(x,y)-dSTO)⁄dSTO where dSTO is the average interplanar distance in the SrTiO3 substrate and d(x,y) is the local interplanar distance. The spatial resolution in the deformation maps is 5 nm and the precision is 2%.

PFM and c-AFM experiments were performed with an AFM microscope (MFP-3D, Asylum Research/Oxford Instruments) at room temperature and under ambient conditions to get access to the local electrical properties. In this case, semi-conductive (001)-oriented STO:Nb (1.4 at%) substrates were used. Ti/Ir-coated tip and cantilever with a nominal ~2.8 N.m -1 spring constant (ASYELEC-01, Asylum Research) were used. PFM analysis were carried out by using dual AC resonance tracking (DART) mode to enhance the detected electromechanical signal. [START_REF] Rodriguez | Dual-frequency resonance tracking atomic force microscopy[END_REF] PFM in spectroscopic mode was conducted by applying a driving voltage of 2V to the local probe. The piezoloops were recorded in remnant mode, i.e. voltage pulses are applied to the active layer and the PFM signal is measured at zero field. Such measurements promote electromechanical activity against electrostatic interaction. [START_REF] Qiao | Electrostatic contribution to hysteresis loop in piezoresponse force microscopy[END_REF] Poling experiments were performed by applying DC voltages of ±10 V to the probe tip. Further details can be found elsewhere. [START_REF] Desfeux | Nanoscale investigations of switching properties and piezoelectric activity in ferroelectric thin films using piezoresponse force microscopy[END_REF] For c-AFM measurements, the probe tip was grounded and the bias voltage was applied to the bottom electrode.

RESULTS AND DISCUSSION

Structural characterizations

In order to characterize the crystalline quality of the films, 2θ-ω diffractograms were collected (see Figure 2a). Several well-defined peaks are observed, corresponding to both substrate (in black) and film (in red). As already reported in the literature and summarized in the introduction part of this paper, depending on the temperature, three allotropic varieties of Nd2WO6 can be stabilized as long as powders are concerned (reminder: low temperature form is monoclinic C 2) Å for the high temperature polymorph). By comparison with ICDD-PDF cards, we identified a preferential growth direction of the film (6 harmonic diffraction peaks series) with a parameter around 8.68 Å, i.e. either the 𝑐 direction of the high temperature NdWO tetragonal variety, or the 𝑐 direction of the high temperature polymorph α-LaWO, orthorhombic. [START_REF] Carlier | Lead-Free α-La2WO6 Ferroelectric Thin Films[END_REF] Indeed, the two high temperature varieties (NdWO-tetragonal and LaWO orthorhombic) are strongly related, and only the a parameter differs between them: aLWO-ortho~3×aNdWO-tetra; b LWO- ortho~b NWO-tetra; c LWO-ortho~c NWO-tetra. However, at this step, and considering only the 00l peaks, it is not possible to discriminate them.

In the case of La2WO6, [START_REF] Carlier | Lead-Free α-La2WO6 Ferroelectric Thin Films[END_REF] we have already shown that starting from a target containing the low temperature polymorph, PLD technique stabilizes the high temperature polymorph. It is important, in the present unknown case of Nd2WO6, to conclude on the true symmetry of the structure and know if the film is either tetragonal or orthorhombic. In addition, if the majority of the diffraction peaks can be attributed to only one c-axis growth direction, the two peaks at 2θ = 32.47° and 68.13°, attributed to (5 1 0) and (10 2 0) orientations respectively, indicate that crystallites are also growing in another direction with respect to the plane of the substrate. It is possibly due to the increasing refractory behavior of Nd compared to La, and the fact that the maximum of temperature that our heater can reach is Tmax = 900°C. To answer this open complex question, we decided to make a complete XRD study.

Considering aSTO = 3.905 Å for commercial STO substrates (Crystal GmbH), and aNdWO = bNdWO = 5.426 Å (by referencing to the Yoshimura results on the tetragonal NdWO form), [START_REF] Yoshimura | Polymorphism of R2O3-WO3 (R = Rare Earth) compounds at high temperatures[END_REF] the lattice mismatch between the film and the substrate can been calculated as follows:

𝛿 = 𝑎𝑆𝑇𝑂 × √2 -𝑎𝑁𝑑𝑊𝑂 𝑎𝑆𝑇𝑂 × √2 × 100 Equation [1]
The quite low value of the lattice mismatch, =1.7%, i.e. meaning low tensile strains induced by the substrate in the plane of the film, indicates a possible matching of aNdWO on a (001)-STO allowing an expected good crystalline quality of the thin film growth.

Epitaxy relationship

In order to further investigate the thin film quality, rocking curves on both 001 peak of the STO substrate and 002 peak of the NdWO thin film were recorded (Figure 2b). The 0.17° full width at half-maximum (FWHM) value measured for NdWO, with respect to the 0.07° one obtained for the substrate, indicates a relatively good crystalline quality with a weak disorientation of the coherent domains observed along the c-axis thin film growth direction.

To get better insight in the characterization of the microstructural properties of our NdWO thin films, φ scan of the 101 and 011 reflections in the tetragonal hypothesis (in green on the Figure 3) indexed as 301 and 011 with orthorhombic hypothesis (in black on the Figure 3), have been recorded. Those specific reflections were chosen because of their relatively high intensities and their position quite isolated from substrate's ones. Whatever the symmetry, these planes diffract at 2ΘBragg = 19.19° and the χ angle between (001) and (301)/(011) is 59.25° (calculated from Carine software [40] ). As expected for an epitaxial growth, four peaks are observed on these scans, with a 90° angular-gap between two consecutive ones, in agreement with simulation made by Stereopole software in the case of tetragonal or orthorhombic symmetry. [START_REF] Salzmann | Stereopole version 1.2 X-Ray Pole Figure Analysis with IDL : software CaRIne[END_REF] Unfortunately, at this step, it is still not sure if the NdWO thin film crystallizes either in the tetragonal cell or in the 3×a superstructure orthorhombic. In order to verify this point, a RSM was recorded on the area of the assumed 521 reflection, peak that exists only if the cell is orthorhombic (it would lead to a non-integer 2.5 2 1 peak in the tetragonal hypothesis). Since, we observed some non-negligible intensity in this area, it proves that the NdWO thin film is orthorhombic (Figure 4a).

In a second time, we located the exact positions (after optimization of the  and  angles) of the 301 (2θ = 19.22°) and 011 (2θ = 19.20°) reflections in order to extract the a and b cell parameters of the film. The c-lattice parameter value was indeed previously calculated from the 2θ- diffractogram given in Figure 2a, i.e. c = 8.68 Å. The a and b film parameters were deduced from the 2θ values, leading to 16.34(5) Å and 5.46(5) Å respectively. By comparison with the ones of α-LaWO (a = 16.56(1) Å, b = 5.72(1) Å; c = 8.87(1) Å), an 8% decrease of the cell volume is observed. This is in accordance with the evolution observed for the cell volume of the Ln2WO6 series with the rare earth radii.

Finally, to investigate the "clamping" state of α-NdWO thin film on the substrate, a last RSM was recorded around the [001]* node of the STO, close to the [002]* node of the film: a perfect vertical alignment between the thin film and the substrate is observed (Figure 4b). It means that the film is perfectly clamped on the substrate, in other words that substrate imposes its in-plane parameters to the in-plane ones of the film.

Epitaxial relationships were deduced from examination of the pole-figure measurements on the 315 film 2θ position, position that allowed to observe on the same figure both 315 peaks from the film and 112 reflection from the substrate. The comparison between the observed and calculated density of poles, using the Stereopole software, [START_REF] Salzmann | Stereopole version 1.2 X-Ray Pole Figure Analysis with IDL : software CaRIne[END_REF] shows that NdWO thin film adopts the following epitaxial relationships: Concerning the (510)-orientation of crystallites with respect to the plane of the substrate, the lattice mismatch was also calculated using the following formula with cNdWO = 8.695 Å:

𝛿′ = 𝑎𝑆𝑇𝑂 × √5 -𝑐𝑁𝑑𝑊𝑂 𝑎𝑆𝑇𝑂 × √5 × 100 Equation [2]
As a result, the value of the lattice mismatch is very low, ' = 0.4% meaning [001]NdWO ‖ [210]STO. However, no clear relationship in the perpendicular direction in the plane of the substrate is evidenced between [51 ̅ 0]NdWO and [12 ̅ 0]STO. This could explain the (00l)-preferential growth of the film when optimized deposition conditions are performed as reported for LWO films. [START_REF] Carlier | Lead-Free α-La2WO6 Ferroelectric Thin Films[END_REF] As previously reported, we believe that the temperature of deposition, too low with our deposition system, is the cause of the coexistence of these two orientations. Nevertheless, as observed on Figure 5, the rocking curve measured on the 510 peak is on the same order of value as the one measured on the 003 peak meaning a similar crystallization of these (510)-oriented crystallites compared to the (00l)-oriented ones.

Figure 6a presents a TEM image of the thin film cross-section, while inset picture is a high-resolution image of the lattice at the interface, confirming that the film is epitaxially grown. Figure 6b shows Fourier transforms calculated in the three regions indicated by squares. The region number 1 corresponds to the STO substrate. The regions number 2 and 3, in the film, correspond to domains with two different orientations. As already pointed out on the 2θ-ω XRD scan (Figure 3a), beside 00l peaks, 510 and 10 2 0 peaks are observed, allowing easy identification of the spots in the vertical direction. Thus, regions number 2 and 3 correspond to a (510) and a (001) orientation, respectively. Figure 6c presents a color-coded map showing the distribution of the two types of domains in the layer, identifiable by triangular or trapezoidal shape on the image for the (510)-oriented domains. It appears that these (510)-domains (red color) are randomly distributed in the (001)-oriented matrix (green color), except at the interface with the substrate, where the (510) orientation is preferred.

In order to obtain more information about the residual strain distribution, lattice deformation maps were calculated from the image in Figure 6a using geometrical phase analysis. [START_REF] Hÿtch | Quantitative measurement of displacement and strain fields from HREM micrographs[END_REF] Figure 6d-f show maps of the horizontal deformation (Exx), vertical deformation (Eyy) and shear deformation (Exy+Eyx), where the deformation is defined with respect to the substrate lattice. Figure 6g shows profiles extracted from the maps along the vertical dashed arrows. The horizontal deformation Exx shows no strong variation at the bottom interface and at the domain walls.

It confirms that the lattice of the layer is epitaxially aligned onto the substrate one and that domain walls are also coherent. On the other hand, the Exx profile in Figure 6g shows that deformation becomes slightly negative away from the interface (in average -2% between 50 and 150 nm). It indicates that mismatch strain between NdWO and STO (estimated to =1.7%, see equation 1) is gradually relaxed along the growth direction. The vertical deformation map Eyy measures the (004) and (510) interplanar distances of the NdWO film with respect to the (002) interplanar distance of the STO substrate. Based on the lattice parameters of the film found using XRD (a = 16.34Å, b = 5.46 Å, c = 8.68 Å) and the lattice parameter of STO (a = 3.905 Å), the corresponding deformations can be calculated theoretically as follows ENdWO(004) = (dNdWO(004)-dSTO(002))/dSTO(002) = 11% for the (001)-oriented domains and ENdWO(510) = (dNdWO(510)-dSTO(002))/dSTO(002) = 43% for the (510)-oriented domains. The deformations measured experimentally are 39%±2% and 13%±2% in the (510) and (001) domains respectively, which is in quite good agreement with the theoretical values. The shear deformation map Exy+Eyx is uniform and close to zero over the whole image. This is another indication that lattices of both layer and substrate are well aligned.

On the other hand, during the cooling process of the thin film synthesis, some thermal mismatches could overcome due to differences between the substrate and film thermal expansion coefficients (TEC). As evocated, this NdWO polymorph has never been stabilized, consequently no α-NdWO TEC is known. As demonstrated by the present structural thin film characterization, α-NdWO and α-LWO are isostructural. Previously α-LWO bulk TEC has been determinated to be around 9.8×10 -6 Å.°C -1 , [START_REF] Allix | Synthesis and Structure Determination of the High Temperature Form of La2WO6[END_REF] and by extrapolation, it is reasonable to consider an α-NdWO TEC in the same order than the α-LWO one. As reminder the STO TEC is around 9.4×10 -6 Å.°C -1 which is near to the film one leading to a parallel thermal dilatation between the thin film and the substrate limiting the texturation (see Figure SI1) and the cracks of the thin film.

As a partial conclusion of this detailed structural and microstructural study, we showed that, by leveraging the PLD method, it is possible to stabilize a new Nd2WO6 polymorph, isostructural to α-La2WO6, which we will hereafter name α-Nd2WO6.

Nanoscale piezo-/ferroelectric properties AFM surface morphology of the synthesized film is presented on Figure 7a. The 8×8 µm² area shows distinct grains and grains boundaries. The grains present irregular shapes with lateral sizes between 100 and 900 nm and a root mean square roughness about 3.6 nm. Nanoscale piezo-/ferroelectric properties have been investigated using PFM tool. Both phase and amplitude loops, recorded in remnant mode for minimizing the electrostatic contribution, are depicted in the inset of the Figure 7a. Clear hysteresis with two opposite ferroelectric states are observed from the phase signal with almost 180° phase difference, while typical butterfly-shape is recorded for amplitude response, suggesting local ferroelectricity under the probing tip. However, in the last years, much attention was paid to curious PFM responses recorded on oxide films, i.e. classical PFM loops obtained on non-ferroelectric material, which can be attributed to several phenomena such as Joule heating, electrostriction, chemical dipole, charge injection, field effect, Vegard strain, deposition, subsurface damaging, surface damaging or vacancy ordering. [START_REF] Balke | Differentiating Ferroelectric and Nonferroelectric Electronmechanical Effects with Scanning Probe Microscopy[END_REF][START_REF] Vasudevan | Ferroelectric or non-Ferroelectric: Why So Many Materials Exhibit "Ferroecltricity" On The Nanoscale[END_REF][START_REF] Seol | Non-Piezoelectric Effects in Piezoresponse Force Microscopy[END_REF][START_REF] Martin | Interpretation of Multiscale Characterization Techniques to Assess Ferroelectricity: The Case of GaFeO3[END_REF][START_REF] Bark | Switchable Induced Polarization in LaAlO3/SrTiO3 Heterostructures[END_REF][START_REF] Borowiak | Electromechamical Response of Amorphous LaALO3 Thin Film probed by Scanning Probe Microscopies[END_REF][START_REF] Li | Electromechanical Response from LaAlO3/SrTiO3 Heterostructures[END_REF][START_REF] Seol | Determination of Ferroelectric Contributions to Electromechanical Response by Frequency Dependent Piezoresponse Force Microscopy[END_REF][START_REF] Carlier | Microstructure and local electrical behavior in [(Nd2Ti2O7)4/(SrTiO3)n]10 (n = 4-8) superlattices[END_REF] Thus, considering the complex nature of the physicochemical phenomena that can exist in such oxide thin films exposed to high electric fields and to accurately verify the ferroelectric-like behavior measured, further PFM analyses were conducted to really confirm that the present response is indeed due to ferroelectricity and not an artifact. More precisely, two strategies currently proposed in the literature were adopted: they consist of probing the biasand time-dependent hysteresis measurements, as reported by Strelcov et al. and Chen et al., respectively. [START_REF] Strelcov | Role of Measurement Voltage on Hysteresis Loop Shape in Piezoresponse Force Microscopy[END_REF], [START_REF] Chen | Mechanisms of Electromechanical coupling in Strain Based Scanning Probe Microscopy[END_REF] In this way, PFM voltage spectroscopy experiments under different AC driving were firstly carried out. Figure 7d displays the evolution of the phase signal by increasing the driving voltage. As observed, the piezoloops are insensitive to the VAC probing voltage as long as the magnitude is lower than the coercive voltage (~5 V), while we note both a narrowing of the hysteresis for applied VAC higher than the coercive voltage and the collapse of the loops. This is due to the flipping of the polarization during the experiments which disturbs the PFM signal at such high driving voltages, which agrees with a ferroelectric switching phenomenon, as already shown. [START_REF] Balke | Differentiating Ferroelectric and Nonferroelectric Electronmechanical Effects with Scanning Probe Microscopy[END_REF][START_REF] Strelcov | Role of Measurement Voltage on Hysteresis Loop Shape in Piezoresponse Force Microscopy[END_REF] A time-dependent hysteresis study was also undertaken and the results are depicted on the Figure 7e. No dependence of the amplitude PFM signal with respect to the cycling period is observed, consistent with a ferroelectric polarization switching process. This behavior enables to exclude an ionic movement under the probing tip which could occur for longer periods, and thus focus on the dynamics of intrinsic polarization as main factor for the amplitude signals recorded, as previously published. [START_REF] Chen | Mechanisms of Electromechanical coupling in Strain Based Scanning Probe Microscopy[END_REF][START_REF] Guan | Identifying Intrinsic Ferroelectricity of Thin Film with Piezoresponse Force Microscopy[END_REF] These bias-and time-dependent hysteresis measurements represent an essential step to ascribe the hysteresis loops measured by spectroscopic PFM to a pure ferroelectric polarization.

From spectroscopic PFM measurements, the surface deformation amplitude was plotted as a function of various driving AC voltages in order to further confirm the electromechanical nature of the detected signal. As seen on the Figure 7f, a linearity is obtained until about 2 V then a saturation is observed, attesting of the piezo-activity of the layer. In the PFM mode, the local deformation amplitude A measured by the probing tip is proportional to the amplitude of the AC voltage VAC, the quality factor Q and the effective piezoelectric coefficient d33 eff , i.e. A = d33×VAC×Q. For contact resonance, the amplitude signal is enhanced and Q is usually ranging from 30 to 100, [START_REF] Mclachlan | Probing Local and Global Ferroelectric Phase Stability and Polarization Switching in Ordered Macroporous PZT[END_REF] but for such inorganic NdWO compound the loss is considered as rather low which leads to a high value for Q, typically of about 70-100. Considering the slope of the positive linear part of the plot in Figure 7f is 185 pm/V (the linearly fitted line leads to y = 185x with r² = 0.9762), we can estimate the effective piezoelectric coefficient between 1.9 and 2.6 pm/V. Note that it is only an estimation and that it is important to consider that the amplitude response is strongly dependent of the measurement location over the free surface of the film. Besides, the clamping of the film on the substrate can lead to an underestimation of the determined value. In addition, it is also wellknown that the electric field beneath the AFM probing tip is inhomogeneous in such PFM experiments, leading to quantitative measurements of signal particularly difficult. [START_REF] Gruverman | Piezoresponse force microscopy and nanoferroic phenomena[END_REF] However, the essential point to keep in mind is that even if it is not possible to give an accurate value for the local d33 eff , the estimated value we measured for deformations (1.9-2.6 pm/V) are promising in view of the development of future lead-free piezoelectric systems based on Ln2WO6 family.

Switching behavior was then studied by means of poling experiments using PFM lithography. Artificial domains were locally induced by applying a DC voltage of +10 and -10 V to the AFM tip over rectangular area at the free surface of the film, then out-of-plane PFM signals were measured. As seen on Figure 7b, strong bright and dark contrasts are obtained for the phase PFM image corresponding to domains with 'upward' (away from the substrate) and 'downward' (toward the substrate) polarization, respectively. Amplitude vibration is revealed by the high contrasts seen on the Figure 7c, while domain boundaries are highlighted by the absence of amplitude signal due to the switch of polarization. The AFM morphology (Figure 7a) simultaneously recorded with PFM images does not reveal any change on the surface, excluding surface electrochemical processes during writing step. [START_REF] Ievlev | Chemical Sate Evolution in Ferroelectric Films During Tip-Induced Polarization and Electroresistive Switching[END_REF] The retention properties of the locally manipulated ferroelectric domains were then investigated. For that, two 'artificial' domains with reverse polarization were first formed by applying ±10 V between the AFM tip and the grounded substrate, as revealed by the contrasted out-of-plane PFM phase signal on the Figure 8b. The tip-induced polarizations were then probed during one day after poling and the obtained signals are presented on the Figures 8b-d (see Fig. SI2 for detailed signal evolution). First, no correlation is observed between the surface morphology (Figure 8a) and the corresponding PFM images (Figure 8b), excluding surface reaction to mainly explain such contrasts, while clear blue and yellow contrasts are observed even after 24h. Also, we can observe some regions with lower contrast from 8h after poling, particularly for positively polarized areas (blue regions), and this contrast decreases with time until 24h, suggesting relaxation phenomena, although contrasted areas remain. The study of the time-dependent back-switched domains was performed in view of determining the characteristic time τ which is a strong indication on the parameter governing the PFM contrasts, i.e. either intrinsic polarization or mediated by the field-induced ion migration in the bulk of the film. The Figure 8e depicts the fraction percentage of switched domains as a function of time. As already reported, we have fitted the relaxation evolution of manipulated domains by means of the following relation: [START_REF] Ganpule | Domain Nucleation and Relaxation Kinetics in Ferroelectric Thin Films[END_REF][START_REF] Guan | Identifying Intrinsic Ferroelectricity of Thin Film with Piezoresponse Force Microscopy[END_REF] 𝑓(𝑡) = 𝑓𝑜 + 𝐴𝑒𝑥𝑝(-𝑡 𝜏 ) Equation (2) with f(t) the fraction of reversed domain correlated to time t; f0 the final percentage of written domains and A the final percentage of back-switched domains. The dashed line in Figure 8e represents the fitting curve according to Equation 2 for values of f0 and A of 90.2 and 9.8%, respectively. We can extract a value for τ of about 30 min which is in good agreement with values already published for other 'true' ferroelectric oxides. [START_REF] Ganpule | Domain Nucleation and Relaxation Kinetics in Ferroelectric Thin Films[END_REF][START_REF] Guan | Identifying Intrinsic Ferroelectricity of Thin Film with Piezoresponse Force Microscopy[END_REF] Such value for the time required to nucleate a reverse domain demonstrates the nonvolatile nature and represents an another proof for ferroelectricity.

Considering both topographic AFM and phase PFM images, relaxed regions (highlighted by the white dotted circles in Figure 8a) correspond to surface heterogeneities and especially to the grain boundaries. However, the back-switching does not occur to the created domain walls, in contrast to previous reported studied on other ferroelectrics thin films, [START_REF] Ganpule | Domain Nucleation and Relaxation Kinetics in Ferroelectric Thin Films[END_REF][START_REF] Ganpule | Polarization Relaxation Kinetics and 180° domain Wall dynamics in Ferroelectric Thin Films[END_REF] except for a relaxation time of about 24h. As a consequence, the common starting point for manipulated domain relaxation seems to be a surface heterogeneity like grain boundaries which act as nucleation sites. This phenomenon was already observed in Pb(Zr0.2Ti0.8)O3 thin films. [START_REF] Gruverman | Nanoscale Imaging of Domain Dynamics and Retention in Ferroelectric Thin Film[END_REF] Moreover, it is well-known the grains boundaries present very high default density, such as oxygen vacancy which particularly exist in transition-metal oxides. Numbers of works demonstrated the ferroelectric domain retention was directly impacted by the oxygen defects into the active layer. [START_REF] Kitanaka | Oxygen-Vacancy-Induced 90°-Domain Clamping in Ferroelectric Bi4Ti3O12 single crystals[END_REF][START_REF] Yamamoto | High-Oxygen-Pressure Crystal Growth of Ferroelectric Bi4Ti3O12 single crystals[END_REF][START_REF] Lee | Polarization Switching and Relaxation Dynamics of Bismuth layered Ferroelectric Thin Films: Role of Oxygen Defect sites and Crystallinity[END_REF] Indeed, these vacancies play a major role on the domain wall mobility, this mobility being directly related to the relaxation dynamic. Besides, mainly positively manipulated domains (blue contrasts) show relaxation. This can be explained by the presence of surface heterogeneities mentioned above on these poled regions. However, this behavior is mainly supported by the fact that these domains were just positively polarized, while two consecutively reverse polarizations were applied to the other domains. Indeed, some studies have reported a better retention behavior for poled domains when the polarization was reverse switched. [START_REF] Ganpule | Polarization Relaxation Kinetics and 180° domain Wall dynamics in Ferroelectric Thin Films[END_REF][START_REF] Kim | Charge Retention Behavior of Preferentially Oriented and Textured Bi3.25La0.75Ti3O12 Thin Films by Electrostatic Force Microscopy[END_REF][START_REF] Lee | Local Retention Behaviors of Epitaxial and Polycrystalline PbMg1/3Nb2/3O3-PbTiO3 Thin Films by Scanning Force[END_REF] Also, the decrease of contrast can be attributed to the depolarization field or owing the domains switched under the tip which might not extend throughout the whole film thickness. [START_REF] Mazet | Structural Study and Ferroelectricity of Epitaxial BaTiO3 Films on Silicon Grown by Molecular Beam Epitaxy[END_REF] As a consequence, these poling studies both confirm the room temperature stability of the ferroelectricity and reveal a typical behavior of ferroelectric domains relaxation.

Then, from both the specific relaxation behavior for manipulated domains and the inability to perform classical P-E hysteresis measurements on these thin films (mainly due to high current leakages), we undertook to locally probe eventual leakage currents into the NdWO layer. The nanoscale conductivity was measured by using c-AFM technique. Current maps were recorded by scanning the same region with the grounded tip and by varying the bias voltage between -10 and +10 V to the bottom electrode. Images measured at -7 V, -8 V, -9 V and -10 V are presented on Figure 9b-e. Differences in conduction properties are observed, as evidenced by the conducting paths in blue contrast. For relatively high voltages (ranging from +10 down to -7 V), the detected current is insignificant, as revealed on Figure 9b. From -8 V, local leakage begins to be detected at several locations, corresponding to 'hot spots' (blue regions in Figure 9c). By increasing bias (absolute value), most of the existing conducting regions extend and new conducting paths appear as seen on Figures 9d and8e, indicating the local conductivity is biasdependent. Black and white dashed circles on the Figures 9c and 8e highlight areas where current clearly extends or where new conducting channels appear, respectively. In addition, obvious correlation between c-AFM contrasts and grain boundaries seen on the associated topographic image (Figure 9a) is revealed. Indeed, most of the conductive areas are located between the grains, as particularly observed on Figure 9e, while the grain regions do not exhibit significant leakage. A careful interpretation of such current signal located at grain boundaries must be done since a convolution effect with the tip can occur inducing such contrasted areas. Considering the contrasts take only place from a high negative threshold bias, we can reasonably conclude that potential convolution effect between the probe tip and the topography is not the main factor leading to highly conductive zones. Therefore, the local conduction in our NdWO film is found to be mainly dominated by the grain boundaries. Such phenomenon was already observed in ferroelectric oxide thin film as Pb(ZrTi)O3, La-doped BiFeO3 or BiFeO3. [START_REF] Masuduzzaman | The Origine of Broad Distribution of Breakdown Times in Polycrystalline Thin Film Dielectrics[END_REF][START_REF] Zhou | Local Electrical Conduction in Polycrystalline La-Doped BiFeO3 thin Films[END_REF][START_REF] Shen | 2D Nanovaristors at Grain Boundaries Account for Memristive Switching in Polycrystalline BiFeO3[END_REF] The profile of leakage currents with the corresponding topography signal along the same line (dashed yellow line in Figures 9a ande) is displayed in Figure 9f, where an unambiguous correlation between the local conduction and the grain boundaries is demonstrated. In addition, we notice that current intensity is not directly proportional to the depth of the grain boundaries, which excludes the tip-sample convolution as being responsible of the obtained conductivity map. Strong current leakage in grain boundaries involves higher defect density in these zones, such as positively charged oxygen vacancies. Indeed, grain boundaries can act as reservoirs for the accumulation of oxygen vacancies and, as previously mentioned, such transition-metal oxides commonly present high oxygen vacancy rate. [START_REF] Shen | 2D Nanovaristors at Grain Boundaries Account for Memristive Switching in Polycrystalline BiFeO3[END_REF][START_REF] Yan | Roles of Grain Boundary and Oxygen Vacancies in Ba0.6Sr0.4TiO3 Films for Resistive Switching Device Application[END_REF] Under external bias voltages, a rearrangement of these vacancies occurs which increases their density in the grain boundaries. Consequently, conducting paths appear at these specific locations. Then, by increasing the bias, the charged oxygen vacancies migrate into the NdWO layer in order to reach a new thermodynamic equilibrium, associated to continuous variation of oxygen vacancy distribution. [START_REF] Choi | Insulating Interlocked Ferroelectric and Structural Antiphase Domain Walls in Multiferroic YMnO3[END_REF] Since the rate of oxygen vacancy diffusion along the grain boundaries is almost five orders of magnitude superior to one of the inner parts of the grains, [START_REF] Zhou | Local Electrical Conduction in Polycrystalline La-Doped BiFeO3 thin Films[END_REF] the current in the grain boundaries increases more rapidly, leading to the specific local current pattern obtained. These observations are very interesting by considering the retention behavior of the polarization discussed above (see Fig. 8). Indeed, in both cases the charged defects located at the grain boundaries, and more precisely the oxygen vacan-cies, can be responsible for the physical phenomena occurring at the nanoscale. As a remark, the observed conduction at grain boundaries is not uniform, i.e. each grain boundary does not reveal conductivity. It can be explained by the different degrees of angle mismatch of each grain boundaries owing the two different crystallographic orientations seen on the XRD pattern, namely (001) and (510) (Figure 2). Indeed, in this case the strain and interfacial energies should differ for different grain boundaries inducing higher defect density, highlighted by electrical hot spots. [START_REF] Masuduzzaman | The Origine of Broad Distribution of Breakdown Times in Polycrystalline Thin Film Dielectrics[END_REF][START_REF] Read | Dislocation Models of Crystal Grain Boundaries[END_REF] Besides, since our NdWO film presents local ferroelectric properties, the application of the high external bias voltage during c-AFM analyses modifies the electric polarization and gives rise to the growth of ferroelectric domains beneath the tip through the extension of the existing domain walls. By increasing the electric field, the domain walls extend towards the grains center. Since no significant current is detected into the grains, as shown on Figure 9e, no conductivity is observed at the typical location of ferroelectric domain walls at the contrary of works focused on BiFeO3 or PbZrTiO3 thin films for example. [START_REF] Seidel | Conduction at Domain Walls in Oxide Multiferroics[END_REF][START_REF] Farokhipoor | Conduction Through 71° Domain Walls in BiFeO3 Thin Films[END_REF][START_REF] Farokhipoor | Local Conductivity and The Role of Vacancies Around Twin Walls of (001)-BiFeO3 Thin Films[END_REF][START_REF] Guyonnet | Conduction at Domain Walls in Insulating Pb(Zr0.2Ti0.8)O3 Thin Films[END_REF][START_REF] Rojac | Domain-Wall Conduction in Ferroelectric BiFeO3 Controlled by Accumulation of Charged Defects[END_REF] This is reinforced by the current map measured after the manipulation of ferroelectric domains (see Figure 7b), where no change in the current pattern was obtained. Consequently, the local conduction is further confirmed to be mainly grain boundaries dependent.

CONCLUSION

A new polymorph of α-Nd2WO6 has successfully been isolated and fully characterized from a structural and ferroelectric point of view. This compound was stabilized in thin film form by taking advantage of the substrateinduced strain during the PLD growth. A particular attention has been paid on the structural characterization of the film by carrying out various experiments by high-resolution XRD including -2, rocking curves,  scan, pole figure and RSMs on well-targeted nodes. By carefully choosing a discriminant reciprocal space node, we demonstrated that this new α-Nd2WO6 polymorph is orthorhombic, isostructural to the high-temperature of bulk α-La2WO6. [START_REF] Carlier | Lead-Free α-La2WO6 Ferroelectric Thin Films[END_REF] By performing bias-and time-dependent PFM hysteresis experiments, and by achieving poling measurements, the local piezo-/ferroelectricity behavior was carefully demonstrated. The room temperature stability of the ferroelectric polarization was also shown. The nanoscale mapping of the conductivity has revealed to be biasdependent and dominated by the grain boundaries. Higher default density such as the accumulation of charged oxygen vacancies in combination with their good mobility in the grain boundaries under external electric field mainly explained the specific local conduction, as well as the polarization relaxation phenomena observed.

Our results extracted from PFM imaging and spectroscopic mode evidence robust intrinsic ferroelectricity into the NdWO layer at the nanoscale level. This finding extends the piezo-/ferroelectric behavior to a second compound belonging to the A2WO6 family (A = lanthanide), the first one being LaWO. [START_REF] Carlier | Lead-Free α-La2WO6 Ferroelectric Thin Films[END_REF][START_REF] Carlier | Microstucture and local electrical investigation of lead-free α-La2WO6 ferroelectric thin films by piezoresponse force microscopy[END_REF] As a consequence, our findings show the strain-engineering as a useful strategy for inducing novel functionalities within oxide thin films, and they particularly allow for extending the class of lead-free ferroelectric materials. [001]* STO

[002]* NdWO

2 𝑐

 2 with a = 5.55(1) Å, b = 11.38(1) Å, c = 15.86(1) Å, β = 107.7, it is orthorhombic P212121 a = 5.536(1) Å, b = 9.231(1) Å, c = 10.170(1) Å at intermediate temperature and finally tetragonal P4 ̅ 21m, a = 5.426(2) Å; c = 8.695(

  [100]NdWO ‖ [110]STO, [010]NdWO ‖ [11 ̅ 0]STO and [001]NdWO ‖ [001]STO (Figure SI1 ) i.e. the same 3-D heteroepitaxy as α-LWO.

Figure 1 .Figure 2 .

 12 Figure 1. Temperature-dependence of Nd2WO6 bulk crystal-structures. Structure projection of (a) low temperature monoclinic Nd2WO6 with [WO6] blocs; (b) orthorhombic high temperature Nd2WO6 allotropic variety with isolated [WO6] polyhedral and(c) high temperature tetragonal Nd2WO6 form which is isostructural to Nd1.2Lu0.8WO6 from Yoshimura results.[START_REF] Yoshimura | Polymorphism of R2O3-WO3 (R = Rare Earth) compounds at high temperatures[END_REF]

)Figure 3 .Figure 4 .

 34 Figure3. Phi-scan for Nd2WO6 on 101/011 peak in case of tetragonal cell (in green) or 301/011 peak on orthorhombic hypothesis (in black). Four peaks are observed; peak repetition rate is 180°.

Figure 5 .

 5 Figure 5. Normalized rocking curves on 003 and 510 NdWO thin film (FWHM are indicated in inset).

Figure 6 .

 6 Figure 6. Transmission electron microscopy images of the Nd2WO6 film. (a) TEM image of the film cross-section. Inset shows a high resolution image of the epitaxial interface. (b) Fourier transforms obtained in the three regions indicated by squares. (c) Color-coded map that shows the spatial distribution of the two types of domains in the image (a). 510-oriented domains are red and 001-oriented domains are green. The map was created by applying filters to the corresponding spots in Fourier space. (d,e,f) Exx horizontal deformation, Eyy vertical deformation and Exy+Eyx shear deformation maps obtained using geometrical phase analysis. (g) Deformation profiles extracted from the maps along the dashed arrows.

Figure 7 .Figure 8 .

 78 Figure 7. Nanoscale piezo-/ferroelectric behavior and bias-and time-dependent hysteresis measurements properties performed on Nd2WO6 thin film. (a) AFM morphology simultaneously measured with out-of-plane (b) phase and (c) amplitude PFM images after poling experiments. The inset of the (a) shows the local remnant phase and amplitude piezoresponse loops recorded on the free surface of the NdWO thin film. (d) Phase PFM loops under different AC driving voltage, (e) and amplitude PFM loops for different period time. (f) AC voltage dependence of the local surface displacement of the NdWO layer.The regression plot between 0 and 2 V corresponding to the linear part (highlighted by the blue rectangular) leads to an effective piezoelectric coefficient between 1.9 and 2.6 pm/V.

Figure 9 .

 9 Figure 9. Local conduction in Nd2WO6 thin film. (a) AFM topography, and current maps of the corresponding NdWO surface recorded with bias voltages of (b) -7 V, (c) -8 V, (d) -9 V and (e) -10 V. Higher conductivity is detected when larger bias voltage is applied. (f) Current signal and associated topography profile along the yellow dotted lines in (a) and (e), indicating the correlation between the local leakages and the grain boundaries. Scale bars for the current intensity are similar for all c-AFM images.

AUTHOR INFORMATION

Corresponding Author * MH Chambrier : mhelene.chambrier@univ-artois.fr

Author Contributions

The manuscript was written through contributions of all authors. / All authors have given approval to the final version of the manuscript. / ‡These authors contributed equally.

ACKNOWLEDGMENT

The authors would like to acknowledge L. Maës for his technical support and A. Verbeke for her help on powder synthesis. Chevreul Institute (FR 2638), Ministère de l'Enseignement Supérieur, de la Recherche et de l'Innovation, Hautsde-France Region, Fonds Européen de Développement Régional (FEDER) and Major Domain of Interest (DIM) "Eco-Energy Efficiency" of Artois University are acknowledged for supporting and funding partially this work. FEDER, CNRS, Hauts-de-France Region and Ministère de l'Education Nationale de l'Enseignement Supérieur, de la Recherche et de l'Innovation are acknowledged for fundings of X-ray diffractometers. Région Hauts-de-France and FEDER are gratefully acknowledged for funding the MFP-3D microscope under Program "Chemistry and Materials for a Sustainable Growth".


L Jeantet

A Regazzi

D Perrin

M F Pucci

S Corn

J.-C Quantin

P Ienny

TWENTY-THIRD INTERNATIONAL CONFERENCE ON COMPOSITE MATERIALS (ICCM23) SURFACE PROPERTIES ASSESSMENT OF RECLAIMED CARBON FIBRES FOR RECYCLING IN PA6/CF COMPOSITES

Keywords: Recycling, Carbon Fibre, Interfacial shear strength

Carbon fibre reclaiming and reuse is both an environmentally and economically fruitful endeavour. The reuse of pyrolysis reclaimed carbon fibres (RCF) in meaningful quantities for high performance parts is promising. However, the harmfulness of the reclaiming process to the mechanical properties of fibres and their adhesion capabilities with polymeric matrix must be investigated. This paper studies the mechanical and surface properties of pyrolysis RCF in comparison with sized and de-sized fibres to assess their potential for reuse. Tensile mechanical properties of these three fibre types were evaluated via micromechanical tensile testing of single fibre samples. RCF show a minor diminution in tensile strength and unchanged elastic modulus which makes these intrinsic mechanical properties suitable for reuse in composite parts. The morphology of the fibre surfaces were analysed using scanning electron microscopy and atomic force microscopy. The scans show clearly different morphologies. RCF present polymeric residues as well as smooth areas very similar to de-sized fibres. Atomic force microscopy was also used to measure the roughness of fibre surfaces. The mean roughness of RCF surfaces was higher than that of de-sized fibres but lower than that of sized fibres. In that context, the residual polymer might be a contributor to better interfacial properties between the fibre and the matrix. Tensiometric tests were carried out to determine fibre surface energy. The RCF showed higher surface energies than de-sized fibres suggesting better adhesion. However, this finding did not transfer to better cohesion between RCF and the polymer matrix. Indeed, in IFSS testing both RCF and de-sized fibres showed similar interfacial shear strength values. The outcomes of this work contribute to the optimisation of the reclaimed carbon fibre re-use in composites. In particular, they highlight the necessity for a surface treatment on RCF to enhance fibre/matrix adhesion and thus composite mechanical properties.

INTRODUCTION

Carbon fibre reinforced polymer (CFRP) composites have become in recent years a staple material in high performance industries such as aircraft manufacturing and automotive [START_REF] Vieira | Strategy and management for the recycling of carbon fiber-reinforced polymers (CFRPs) in the aircraft industry: a critical review[END_REF][START_REF] Das | Preparation, development, outcomes, and application versatility of carbon fiber-based polymer composites: a review[END_REF]. Their high tensile strength and lightweight properties are highly useful in those endeavours. The newfound abundance of these materials has created a need for a more sustainable production and a controlled life cycle. Reclaiming carbon fibre from composite waste has been established as economically and environmentally beneficial.

The carbon balance of the reclaiming process is positive, even for energy demanding processes such as pyrolysis, as the energy demands for the carbon fibre (CF) manufacturing process are tremendous [START_REF] Meng | Environmental aspects of use of recycled carbon fiber composites in automotive applications[END_REF][START_REF] Meng | Comparing life cycle energy and global warming potential of carbon fiber composite recycling technologies and waste management options[END_REF]. As a result, reclaiming processes that can be scaled to industrial proportions, such as pyrolysis reclaiming, have been developed and are used to process composite waste [START_REF] Khalil | Sustainability assessment of solvolysis using supercritical fluids for carbon fiber reinforced polymers waste management[END_REF]. With the increasing use of CFRP in airplane carriers, there is an endless supply of composite waste to repurpose and, at the same time, an increased demand for high performance composite parts.

However, the nature of the recycling process does not allow for reuse of the reclaimed fibres (RCF) in woven CF fabric, for the time being. To reuse the RCF in high performance application, the recycled composite formulation will need to be optimized for cut random fibres. Moreover, previous studies highlight that, although the reclaiming process does not affect the fibre intrinsic mechanical properties significantly, it does negatively affect the polymer/fibre interfacial properties [START_REF] López | Recovery of carbon fibres by the thermolysis and gasification of waste prepreg[END_REF][START_REF] Han | Evaluation of fiber surface treatment on the interfacial behavior of carbon fiber-reinforced polypropylene composites[END_REF]. This phenomenon then translates to lower composite mechanical properties as the interface between the components of the composite materials is impaired in comparison with a composite produced using sized virgin CF. This effect could be due to a degraded fibre surface chemical state, less mechanical interlocking via surface roughness modification, or both.

The aim of this study is to determine and quantify the factors that influence the interfacial properties of the pyrolysis RCF. This will then allow the selection of a specific restorative treatment to grant optimal cohesion between the RCF and the matrix. As a result, competitive mechanical properties of the recycled composite material will be ensured. The intrinsic mechanical properties of the RCF were evaluated via micromechanical testing to ensure no significant change in mechanical properties compared to virgin fibres (VF) would interfere with further experiments. The modifications of the wettability of fibres surface by the reclaiming process were characterized via single fibre contact angle determination. The RCF surface morphology and roughness were evaluated using scanning electron microscopy and atomic force microscopy. The fibre/matrix interfacial shear strength (IFSS) was then assessed from microdroplet tests.

MATERIALS

The thermoplastic matrix selected for the preparation of composites was Technyl C206F Natural, a commercially available PA6 produced by Solvay (Brussels, Belgium). Sized virgin fibres of different origins obtained from spools of pure carbon fibres with a PA6 sizing and cut to a length of 20 mm were supplied by Procotex (Mouscron, Belgium). Reclaimed carbon fibres were purchased from ELG Carbon Fibre (Bilston, United Kingdom), the fibres were cut to a 20 mm mean length after a two-step reclaiming process. According to previous studies using these fibres [START_REF] Holmes | Recycled carbon fiber composites become a reality[END_REF][START_REF] Akonda | Recycled carbon fibre-reinforced polypropylene thermoplastic composites[END_REF], the two-step treatment consists of a pyrolysis treatment of 30 min at 500 °C under nitrogen followed by a second step of thermo-oxidation of 10 min at 500 °C under air to remove the char produced during pyrolysis. A third type of fibre was prepared by making the virgin fibres undergo a thermo-oxidative treatment, like the char-removing treatment used by ELG (i.e. 10 min at 500 °C under air), to eliminate the coating. For clarity purposes virgin fibres will be noted as VF, commercial reclaimed fibres will be referred to as RCF and virgin fibres after thermal treatment as VFT.

METHODS

Fibres tensile testing

Fibre transverse dimensions were determined using a FDAS770 fibre dimensional analysis system from Dia-Stron (Andover, UK). 200 projections of the rotating fibre were acquired at 20 locations along its length to determine local minimum and maximum diameters 𝐷 𝑚𝑖𝑛 and 𝐷 𝑚𝑎𝑥 respectively. The corresponding cross-sectional areas were computed assuming that each cross-section is elliptical.

Fibre mechanical properties were evaluated using a LEX820 linear extensometer from Dia-Stron. Tensile stress at break was calculated using the following equation:

𝜎 = 𝐹 𝑆 𝑚𝑒𝑎𝑛 (1)
Where 𝐹 is the force at break, and 𝑆 𝑚𝑒𝑎𝑛 being the mean value of the elliptical cross-sectional areas of the fibre. Displacement values measured during tensile testing were adjusted to consider the inherent compliance of the experimental set-up according to the ISO 11566:1996(EN) standard. The compliance value was established at 0.28 mm.N -1 for our device (from a preliminary assessment performed on similar samples with various lengths). For each fibre type, the repeatability was ensured by testing 25 samples. The nominal fibre length of these samples was 12 mm. Tensile modulus was determined by linear regression of the stress vs strain plots, in the elastic deformation domain between 0.05 % and 0.25 % deformation according to the ISO 11566:1996(EN) standard..

Scanning Electron Microscopy and fibre morphology assessment

A Quanta 200 FEG scanning electron microscope (SEM) from FEI Company (Hillsboro, USA) operating at 4 kV was used to observe the surface morphology of the fibres.

Atomic Force Microscopy fibre topography and fibre roughness determination

An MFP-3D Infinity atomic force microscope (AFM) from Asylum Research (Santa Barbara, USA) was used in a tapping mode to assess the topography of fibres. From this topographic data, roughness values were computed. Single fibres were fixed to a glass slide for tests. A silicon probe (AC160TS-R3) with a resonant frequency of 300 kHz and a spring constant of 25 N/m was used. The scan rate was set to 1 Hz. Topography images with a region of interest (ROI) of 3 × 3 µm 2 were obtained at a resolution of 256 × 256 pixels. A three-dimensional topography of the fibre's surface was acquired using the IGOR Pro 6.27 software. These scans were then processed using a Matlab program to deduce the roughness values by subtracting the natural curvature of the fibre surfaces. According to the ISO 21920-2:2021 standard, the reported roughness values use the common statistical descriptor 𝑆 𝑞 (height root mean square) expressed as:

𝑆 𝑞 = √ 1 𝐴 ∬ d𝑧 2 (𝑥, 𝑦)d𝑥d𝑦 𝐴 (2)
Where 𝐴 is the area of the ROI and 𝑧(𝑥, 𝑦) a function giving the height value for a given data point defined by its abscissa 𝑥 and its ordinate 𝑦.

Fibres wettability evaluation

3.1 Contact angle tests on single fibres were performed using a Krüss K100 tensiometer at ambient temperature using the Wilhelmy principle [START_REF] Volpe | The Wilhelmy method: a critical and practical review[END_REF]. Test liquids used were n-hexane, water, diiodomethane and ethylene glycol. Water was used for its high polarity and surface tension, n-hexane was used as a totally wetting liquid, due to its low surface tension, diiodomethane was used for its low polarity while not totally dispersive, and the ethylene glycol was used for its good balance between the two components. The vessel speed was 1 mm.min -1 and the immersion length 3 mm. The Wilhelmy relationship was used to determine the contact angles:

𝐹 𝑐 = 𝑃 𝛾 𝑙 cos 𝜃 (3)
Where 𝐹 𝑐 is the capillary force measured by the tensiometer, 𝛾 𝑙 is the surface energy of the test liquid, 𝜃 is the measured contact angle, and 𝑃 is the perimeter of the fibre. Perimeter of fibre was obtained thanks to tests described in section 3.1. The advancing and receding contact angles were measured as well, only the static angle was considered for the surface energy assessment. Contact angles presented here were determined from a mean of at least 7 contact angles measurements. Surface energies of fibres were deduced from these values using the Owens-Wendt relation in a linear representation:

𝛾 𝑙 (1+𝑐𝑜𝑠𝜃) 2√𝛾 𝑙 𝑑 = √𝛾 𝑠 𝑝 √𝛾 𝑙 𝑝 √𝛾 𝑙 𝑑 + √𝛾 𝑠 𝑑 (4)
With the polar and dispersive components of test liquids (𝛾 𝑙 𝑑 and 𝛾 𝑙 𝑝 ) and equilibrium known contact angles, this equation allows the determination of the polar and dispersive components of solid surface energy of the fibre surfaces (𝛾 𝑠 𝑑 and 𝛾 𝑠 𝑝 ).

Microdroplet IFSS testing

The samples were prepared in a similar manner as described by Ma et al. [START_REF] Ma | Enhanced interfacial properties of carbon fiber reinforced polyamide 6 composites by grafting graphene oxide onto fiber surface[END_REF]. Finely ground PA6 was obtained by cryogenic grinding of the pellets. This powder was then applied to the fibre using electrostatic adhesion and heated to 240 °C using a copper wire. An example of a sample is presented in Figure 1.

Figure 1 -Microdroplet sample observed via optical microscopy

The melted polymer then naturally formed microdroplets around the fibre. Adequately sized droplets were selected using an optical microscope equipped with a camera to measure their diameter and embedded length. The droplets were selected to have an embedded length between 60 and 120 µm. Mechanical characterization of the fibre/matrix system was conducted using a LEX820 linear extensometer from Dia-Stron equipped with the microdroplet pull-out module and a 20 µm vise. The droplet was positioned at a 500 µm distance from the vise and then pulled over 800 µm through it using a 5 µm.s -1 testing speed. The contact of the droplet with the vise created an increase in the measured charge that ended with the detachment of the droplet. The peak force value at debonding 𝐹 𝑝 , is then used to determine the IFSS value, using the equation:

𝐼𝐹𝑆𝑆 = 𝐹 𝑝 𝜋𝐿(3(𝐷 𝑚𝑖𝑛 +𝐷 𝑚𝑎𝑥 )-√(3𝐷 𝑚𝑖𝑛 +𝐷 𝑚𝑎𝑥 )(𝐷 𝑚𝑖𝑛 +3𝐷 𝑚𝑎𝑥 ) ) (5)
Where 𝐿 is the measured embedded length for each fibre, and 𝐷 𝑚𝑖𝑛 and 𝐷 𝑚𝑎𝑥 are mean values determined for each fibre type using the FDAS770 fibre dimensional analysis system from Dia-Stron.

Cross-sections of the fibres were considered elliptical, and their perimeter was approximated using the Ramanujan formula [START_REF] Villarino | Ramanujan's inverse elliptic arc approximation[END_REF].

RESULTS AND DISCUSSION

Fibres mechanical properties

The mean tensile strength of the single fibre CF samples is presented for each fibre type in Figure 2. As expected, RCF shows a tensile strength 15 % inferior compared to the other two fibre types. These values agree with the mechanical performance reported by the supplier as well as the literature on pyrolysis RCF [START_REF] He | The effect of sizing and surface oxidation on the surface properties and tensile behaviour of recycled carbon fibre: An end-of-life perspective[END_REF][START_REF] Jeong | Fast recovery process of carbon fibers from waste carbon fibers-reinforced thermoset plastics[END_REF].

Figure 2 -Tensile strength mean values for each fibre type

The mean tensile strength values of the two other types of fibre, VFT and VF, are not significatively different. This suggests that the short (10 min) thermo-oxidative treatment used to remove the char from the surface of the VF does not negatively affect their mechanical properties. Long exposure to high temperatures (around 500 °C) however seems to cause a mild diminution of the tensile strength of the fibres. The diminution of the fibre mass and diameter, leading to a higher concentration of fracture inducing defects in a given volume, is the preferred hypothesis to explain this phenomenon in the literature [START_REF] Abdou | Recycling of polymeric composites from industrial waste by pyrolysis: Deep evaluation for carbon fibers reuse[END_REF].

Figure 3 -Tensile modulus values calculated using raw displacement output and corrected for system compliance

In contrast, the three fibre types have very similar tensile moduli, as shown in Figure 3. The main difference is the dispersion of the data, with the VFT and RCF having a larger standard deviation compared to VF. Combined to the diminution in tensile strength, these results indicate that the recycling process increases the probability of rupture of the fibre by increasing the incidence of defects. On the other hand, elastic properties are unaffected by these changes.

Fibre surface morphology

SEM observations of the three fibre types are illustrated in Figure 4. The textures of the different samples show clear differences. RCF have a smooth surface with localized aggregations of residual polymer that resisted the reclaiming process. VFT have a very similar smooth surface all around without the chunks of residual polymer, the sizing seems to have been entirely removed revealing the bare fibre surface. VF have a more textured appearance showing the presence of sizing, dark spots are dispersed on the surface demonstrating that the sizing is heterogeneously distributed on the fibre.

Figure 4 -SEM scan of a) bundle of VF, b) VFT, c) RCF

These SEM images show clear morphological differences between the surfaces of the three fibre types. In order to quantify the differences of fibre morphology, fibre surface roughness would need to be determined. That is why topography must be assessed by AFM to quantify the roughness values of the fibres surfaces.

Fibre surfaces topography

AFM scans of RCF, VFT and VF are illustrated in Figure 5. An additional scan of RCF showing a smooth portion of the fibre surface, without the presence of residual polymer, was also acquired. The surface of RCF shows a very heterogeneous surface with chunks of residual polymer as well as smoother areas. In the case of the sample without residual polymer, the fibre bare surface is exposed with a very smooth and homogenous surface aspect. In a similar fashion, VFT surface is very smooth and homogeneous without residual polymer nor sizing. In contrast, VF has a visually grainy appearance which suggests a high roughness due to the sizing heterogeneous repartition. The overall appearance of the samples is consistent with previous SEM observations. The results are shown in Table 1.

Table 1 -Roughness measurement for each fibre type

The VF has the highest roughness, that should facilitate the mechanical interlocking with the matrix. VFT and smooth RCF have low roughness, probably due to the removal of sizing. RCF with polymeric residues has higher mean roughness because of localized polymeric residues. Higher roughness could play a favourable role on mechanical interlocking and then on fibre/matrix cohesion. Cohesion will be characterised thanks to IFSS microdroplet testing.

Fibre surface wettability

Surface energies and components were determined to evaluate the potential adhesion of the fibre with the matrix. Values obtained for each fibre type are displayed in Figure 6. In previous studies [START_REF] Garat | Surface energy determination of fibres for Liquid Composite Moulding processes: Method to estimate equilibrium contact angles from static and quasi-static data[END_REF][START_REF] Chen | Effect of hybrid polysphosphazene coating treatment on carbon fibers on the interfacial properties of CF/PA6 composites[END_REF], the main predictors of the adhesion of fibres to polymer matrix were the total surface energy and the ratio of polar component over said total energy. These two factors are, respectively, 0.63 and 39.9 mN/m for VF; 0.52 and 36.2 mN/m for VFT, 0.58 and 38.9 mN/m for RCF. Despite the absence of sizing, RCF still has a relatively high surface energy ratio thanks to the presence of polymeric matrix residues, which affects the wettability of fibres surface. Based on surface energies, RCF exhibits a better potential for adhesion to a polymeric matrix than VFT. This is due to the residual polymer not eliminated by the recycling process allowing for better chemical bonding.

Figure 6 -Surface energy values for each fibre type

As expected, the sizing present on VF surfaces generates the highest values of surface energies that improve the adhesion with the thermoplastic matrix. The improvement in adhesion properties is correlated in the literature with increased interfacial shear strength. This has to be confirmed experimentally through IFSS testing.

Interfacial shear strength

The interfacial shear strength values obtained from microdroplets pull-out tests are shown in Figure 7. The interfacial shear strength between PA6 and VF is about 40 % higher than with VFT or RCF. These results confirm that, in the case of short CF and thermoplastic composites, sizing has a pronounced beneficial effect on interfacial properties [START_REF] Zhang | Influence of sizing molecular weight on the properties of carbon fibers and its composites[END_REF].

Figure 7 -IFSS values obtained for each fibre type

It raises the question of RCF re-sizing or more generally surface treatments. The viability of surface treatments to restore the RCF capacity to adhere to a thermoplastic matrix is a promising path for future research.

CONCLUSION

The intrinsic mechanical properties of sized (VF), de-sized (VFT), and reclaimed carbon fibres (RCF) were measured and compared. The recovery process was found to have a limited impact on the mechanical properties of the fibres, as a minor 15 % decrease in their tensile strength and an unchanged elastic modulus were observed. RCF are not limited in their reuse by their mechanical properties, but by their degraded surface properties that prevent adhesion to the matrix as well as interfacial properties.

SEM and AFM scans highlight very different surface morphologies between RCF, VFT and VF. The presence of polymeric residues on RCF alongside smooth portions of bare fibre very similar to VFT surface result in a higher mean roughness that could facilitate the adhesion with the matrix. Moreover, RCF present higher values of surface energy compared to VFT that should potentially improve the adhesion with the matrix. In order to explain the different ratios of dispersive versus polar components a finer analysis of the fibre surfaces chemical state could be useful. Although RCF exhibited a higher surface energy than VFT, no significant advantage in cohesion could be observed via IFSS tests.

The results presented in this study contribute to identify the key parameters that influence fibre/matrix interfacial properties at the microscale. Further investigations at the larger scale of composite constitute the perspective of this work. Additionally, surface treatments such as plasma functionalization, re-sizing or chemical oxidation of the surface could be useful to further enhance adhesion between fibres and matrix and thus the mechanical performance of the fibre/matrix interface.

Figure 5 -

 5 Figure 5 -Topography of a 3 × 3 µm ROI of a) VF, b) VFT, c) smooth RCF, d) rough RCF