Suvidyakumar Homkar

Daniele Preziosi

Xavier Devaux

Corinne Bouillet

Johanna Nordlander

Morgan Trassin

Francois Roulland

Christophe Lefèvre

Gilles Versini

Sophie Barré

François Roulland

Sophie Barre

Cédric Leuvrey

Marc Lenertz

Manfred Fiebig

Geneviève Pourroy

Nathalie Viart

Ultrathin regime growth of atomically flat multiferroic gallium ferrite films with perpendicular magnetic anisotropy

à la diffusion de documents scientifiques de niveau recherche, publiés ou non, émanant des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

I. INTRODUCTION

The electronics industry is yearning to benefit from the wide range of sizeable improvements potentially offered by multifunctional oxides. Magnetoelectric multiferroics are among the most promising ones, allowing the manipulation of magnetic properties by an electric field. Despite their scarcity, they inspire a strong research activity with high expectations mainly in the field of low-power spintronics [1][2][3][4][5][6]. The difficulty to retain and/or shift their magnetoelectric multiferroic properties towards room temperature, especially in the form of thin films [7], makes the growth optimization process of these oxides a necessary step towards their effective implementation. The literature presents various examples where heterostructuring and epitaxial strain have been used to modify the multiferroic properties of high quality bismuth ferrite [8][9][10] thin films. On the contrary, thin films of room-temperature magnetoelectric multiferroic Ga 2-x Fe x O 3 (0.7 x 1.4) (GFOx) have not been much studied in that perspective, despite the alternative interesting features this material can bring [11,12]. Bulk GFOx is ferrimagnetic with a net magnetization and a Curie temperature above room temperature for x > 1.3 [13]. These properties are retained in thin films and a room temperature magnetization of 100 emu/cm 3 is observed for GFO1.4 thin films [11]. GFO shows a linear magnetoelectric effect of approximately 10 -11 s/m [13,14] and a spontaneous polarization of about 25 μC/cm 2 [15,16]. It crystallizes in the polar orthorhombic Pna2 1 (equivalently Pc2 1 n) space group, different from the usual perovskite structure adopted by other magnetoelectric compounds, with a = 0.5086 (2) nm, b = 0.8765(2) nm, and c = 0.9422(2) nm for x = 1.4 [17]. Its structure is based on a double hexagonal compact stacking of oxygens with cations occupying four different sites: a tetrahedral one, Ga1, and three octahedral ones, Ga2, Fe1, and Fe2. The complexity of this structure makes its deposition as thin films with controlled interfaces difficult. Among the few studies in the literature reporting on the growth of thin films of GFOx [11,[18][19][20], none addresses the growth mode nor the possibility to obtain a two-dimensional (2D) growth mode. The need for atomically smooth thin films is, however, a fundamental prerequisite in spintronics to avoid a charge/spin transport degradation across the interface due to electrons scattering [21]. Sharp interfaces are mandatory for observing effects like a spin-orbit torque which is proposed for low-power spintronics devices [22]. A root-mean-square (rms) roughness higher than 1-2 nm ruins the reproducibility of spintronics effects, hindering the desired device integration [23].

In this paper we report on the elucidation of the growth mode of GFO thin films by reflection high energy electron diffraction (RHEED)-assisted pulsed laser deposition (PLD) from the very early stages of the growth. We show that room temperature ferrimagnetic and polar ultrathin films of GFO1.4 with rms roughness of less than a nanometer can be obtained with a control of their thickness down to a fraction of the unit cell. The magnetic anisotropy of the films is tunable via the sample thickness. The films already present an electric polarization for a thickness as low as 7 nm. This study lays the ground for realistic use of GFO thin films for spintronics applications.

II. EXPERIMENTAL DETAILS

Ga 0.6 Fe 1.4 O 3 thin films were deposited on strontium titanate, SrTiO 3 (STO) (111) substrates (Furuuchi Chemical Corporation, Japan, with rms roughness lower than 0.15 nm) using a PLD system having a base pressure of 2 × 10 -8 mbar. A polycrystalline target of stoichiometric Ga 0.6 Fe 1.4 O 3 composition was prepared by the ceramic method [24] and ablated with a 248-nm wavelength KrF excimer laser at a repetition rate of 2 Hz in an oxidizing atmosphere with a O 2 partial pressure of 0.1 mbar. The target-substrate distance was 5.5 cm and the energy density of the laser on the target was 4 J/cm 2 . The substrate temperature was kept at 900 °C during the growth. In situ RHEED was used to monitor the growth of the thin films. The growth was stopped after a number of pulses comprised between 500 and 24 000 to produce GFO samples of different thicknesses. Thickness and structure were studied by x-ray reflectometry and diffraction, respectively, using a Rigaku Smart Lab diffractometer equipped with a rotating anode and a monochromated copper radiation (1.54056 Å). The surface of the films was observed and their rms roughness quantified by atomic force microscopy (AFM) using a Bruker ICON microscope in tapping mode. Energy dispersive x-ray spectroscopy (EDX) coupled to a scanning electron microscopy technique (JEOL 6700 F) was used to check the overall composition of the films. The analysis was performed at 5 keV, ensuring a large surface sensitivity of the EDX signal. Observations of the films in cross sections were performed by high resolution transmission electron microscopy (HR TEM) and high-resolution scanning transmission electron microscopy (HR STEM) using a probe corrected microscope JEOL ARM200F equipped with the Gatan GIF Quantum SE system for electron energy loss spectroscopy (EELS). The lamellae were prepared by focused ion beam (FIB) and rethinned by low voltage (200 V) precision Argon ion milling. The microscope was operated at 80 kV. Multivariate statistical analysis has been used to improve the quality of the STEM-EELS mapping data. We studied the electric polarization state of the films by optical second harmonic generation (SHG). SHG is a nonlinear optical process denoting the emission of light at frequency 2ω from a crystal irradiated with light at frequency ω [25]. This is expressed by the equation P i (2ω) = ε 0 j,k χ (2) i jk E j (ω)E k (ω), where E j,k (ω) and P i (2ω) are the electric-field components of the incident light and of the nonlinear polarization, respectively, with the latter acting as the source of the SHG wave. The nonlinear susceptibility χ (2) i jk characterizes the ferroelectric state. The experiments were performed in transmission mode, after polishing the back of the substrate, with a fundamental wavelength of 1200 nm (see Fig. S1 in the Supplemental Material [26]). In our experiments the polarizer [setting E (ω)] and analyzer [selecting component of E (2ω)] are kept parallel to each other and both are rotated simultaneously by 360°to obtain the SHG anisotropy, where 0°and 90°correspond to the vertical and horizontal laboratory axes, respectively. The out-of-plane (OOP) SHG component was extracted from the sin 2 (x)-dependence of the SHG intensity on the sample tilt with respect to the incident laser beam, where this angle was varied between -30°and 30°. For probing the GFO films, we used light pulses emitted at 1 kHz from an amplified Ti:sapphire system with an optical parametric amplifier. The light pulses had a photon energy of 1.03 eV, a pulse length of 40 fs, and a pulse energy of 20 μJ. The setup for SHG is described in Ref. [27]. The magnetic properties of the films STO (111) substrate were studied with a superconducting quantum interference device (SQUID) vibrating sample magnetometer (SQUID VSM MPMS 3, Quantum Design) in both parallel and perpendicular configurations. Hysteresis loops were measured for magnetic fields up to 7 T, at 10, and 300 K.

III. RESULTS AND DISCUSSION

The in situ RHEED monitoring during the GFO1.4 thinfilm growth onto STO(111) single crystals rendered a 3D to 2D transition. In order to study this phenomenon, we investigated films of various thicknesses. Figure 1(a) shows some of the RHEED patterns imaged at the end of the deposition. A 3D growth mode is observed at 4 nm as indicated by the dotted RHEED pattern. The RHEED pattern exhibits a transition to a state with a modulated streaky feature at 7 nm, which then becomes fully streaky at 32 nm, indicating a 2D sample surface. The 3D to 2D growth-mode transition as observed via RHEED thus appears for a thickness of approximately 7 nm. The AFM images [Fig. 1(b)] also show an evolution of the surface morphology from randomized islands for 4 nm, through islands aligned in rows, and towards atomically flat films, when islands finally coalesce, in a process very similar to the one reported for the growth of SrRuO 3 onto STO (111) [33]. The rms roughness value is around 1 nm at the start of the growth and progressively decreases, with increasing film's thicknesses, to values comparable to the rms roughness of the substrate prior to deposition. The growth was also monitored by RHEED from the intensity variation of the spots in the zero-order Laue zone. Figure 2 covers the complete deposition of the 7-nm GFO1.4 film. The overall deposition lasted for 537 s, yielding a deposition rate of 0.0130 nm/s. The 2D growth sets in approximately 108 s after the growth starts, and lasts for 464 s, until the growth is stopped. It shows 24 oscillations. One RHEED oscillation is 19.3 s long. It corresponds to the deposition of 0.25 nm, i.e., about 1/4 th of a unit cell. Such a subunit cell growth is rare. It had until now been observed only for the growth of other complex cells, such as those of the spinel [28] or garnet [29] phases of iron oxides.

X-ray diffraction patterns of the deposited films, performed in the θ -2θ mode, are shown in the inset of Fig. 3. They indicate that, for all thicknesses, the GFO thin films are well crystallized, oriented along the Pna2 . This is in perfect agreement with the symmetry of the system as shown in Figs. S3 and S4 [26], and as already reported in the literature [19,[START_REF] Demchenko | Investigation of the Potential Offered by Gallium Iron Oxide Thin Films in Terms of Multiferroicity[END_REF][START_REF] Lefevre | [END_REF]. GFO a, b, and c cell parameters could be determined from the combination of θ -2θ scans and reciprocal space mapping of the 206 and 057 reflections (see Fig. S5 [26]). While c decreases with increasing thicknesses (as already observed from the GFO (004) peak in the θ -2θ scans), a and b remain essentially constant (see Fig. S6 [26]). The out-of-plane c parameter reaches the bulk value for thicknesses higher than 64 nm, after a decrease of about 0.2%. The out-of-plane expansion of the c parameter for lower thicknesses is thus not related to any substrate-induced in-plane strain, since it relaxes totally independently from the a and b parameters. The mismatch in this GFO (001) growth on STO FIG. 3. X-ray diffractograms of the deposited GFO thin films in the θ-2θ mode, with a focus on the GFO 004 peak, the vertical line at 2θ = 38.17 (111) system is large. The distances which have to be considered are (i) along the a GFO direction: a GFO = 0.5088 nm ≈ 3d STO 121 = 0.4780 nm (6.4% compressive strain), and (ii) along the b GFO direction: d GFO 060 = 0.1465 nm ≈ d STO 220 = 0.1380 nm (6.1% compressive strain) (see Figs. S3 and S4 [26]). The relatively high elastic energy introduced by this large strain value accounts for the absence of in-plane strain state in the GFO crystal structure even for the lowest thicknesses. This might also explain the initial 3D mode growth, where totally relaxed GFO islands coalesce while the deposition proceeds to obtain a flat surface. Such a 3D-to-2D metamorphic epitaxial growth [32] mechanism has already been reported for SrRuO 3 [33] onto STO and for ε -Fe 2 O 3 [34,35], which is isomorphic to GFO. In both cases, the initial 3D growth mode was related to the large lattice mismatch between substrate and thin film.

The electric polarization state of the films was studied by SHG. This contact-free technique is frequently used to probe the breaking of inversion symmetry related to the onset of ferroic order such as spontaneous polarization in bulk crystals and thin films [25]. While the multiferroic properties of bulk GFO have been extensively studied by this technique [36][37][38][39][40][41], reports on SHG probing of the polar properties of GFO thin films are limited [20,42]. The allowed SHG components related to the polarization in GFO are set by its orthorhombic mm2 point group with five independent elements. However, in the thin films, due to averaging over the three in-plane 120°-rotated crystallographic variants, the effective point-group symmetry increases to 6mm, which holds only three independent components, simplifying the SHG analysis in the thin-film case. By tilting the sample with respect to the incident laser beam, the OOP-SHG component related to the spontaneous polarization along c GFO can be accessed. The SHG dependence on the direction of the incident and detected light polarizations for the 32 nm GFO film at 30°tilt is shown in Fig. 4(a). The double-lobe symmetry mainly along the horizontal direction (along the projection of P s ) is in agreement with the mm2 point group averaged over the three crystallographic variants, and with previous bulk [37] and thin film [42] measurements. The offset of the lobes from the horizontal direction is due to a superposition of the STO substrate surface SHG together with the OOP-SHG contributions from the film. The thickness dependence of the GFO electric polarization was studied by performing tiltdependent SHG measurements [see Fig. each GFO thickness is plotted in Fig. 4(c). The SHG intensity increases with films thickness. We detect a polarization in our films down to a GFO thickness of 7 nm.

Figure 5 summarizes the magnetic characterization of our GFO thin films series, showing a thickness dependence of the magnetization easy axis. It is worth noting that magnetization lies in-plane for films thicker than 11 nm and is, surprisingly, oriented out-of-plane for the 7-nm-thick sample. This behavior, already observed at room temperature, is clearly confirmed by low temperature measurements [Fig. 5(b)]. Films presenting out-of-plane magnetization are of high interest for spintronic based applications. They indeed address the current challenge to develop high quality sub-10-nm-thick ferrimagnetic insulating films with perpendicular magnetic anisotropy, which are desired for efficient spin current transmission in ferromagnetic/Pt spin Hall effect driven systems [43]. We observed, for samples thicker than 11 nm, the saturation magnetization value of approximately 100 emu/cm 3 already observed for GFO1.4 thick films (100 nm) [11], while the 7-nm-thick GFO sample showed a value of approximately 80 emu/cm 3 .

Deeper insight into the structure of the GFO films was obtained with HR STEM. High angle annular dark field (HAADF) images of a cross section of the 32 nm GFO film is presented in Fig. 6. The interface between the STO substrate and the GFO film is well defined and the film shows the cationic pattern expected for GFO in its Pna2 1 space-group structure from the beginning of the growth. A zone of darker contrast, indicating a strong evolution of the local chemical composition (lower average atomic number), is clearly visible and delimitates the first five nanometers from the rest of the film. One should note that observations of other areas of the sample reveal a slight dispersion in the position of this delimitation at distances between 2 and 5 nm from the substrate. Mapping the convex or concave shape formed by the four Fe2 and Ga2 sites in a row (cf. GFO unit cell in Fig. 2 for the positions of Fe2 and Ga2) allows the determination of the polarization orientation within each unit cell [START_REF] Lefevre | [END_REF]. Interestingly, we observe a polarization reversal as we move away from the film/substrate interface. This corresponds to the first experimental observation of ferroelectric domains in ultrathin GFO films. The domain wall imaged in Fig. 6 corresponds to a nominally charged tail-to-tail configuration. While neutral (normal to the surface) ferroelectric 180°domain formation in the ultrathin regime is expected in absence of charge screening [44][45][46][47] for classical systems, the stability of the tail-to-tail domain wall type observed here might be attributed to the highly energetic switching path of GFO. The polarization switching is expected to occur along the Pna2 1to-Pnna phase-transition path with an activation energy in the 0.5-1 eV range per formula unit [16,19,48]. This high switching activation energy results in a high coercive field which would therefore stabilize electrostatically unfavorable domain architectures. Similar charge domain walls have only been observed recently in improper ferroelectrics such as YMnO 3 , in which the domain pattern is set by nonferroelectric primary order parameter [49]. The tail-to-tail polarization configuration, with a domain boundary at about 5 nm from the substrate, may be at the origin of the increased outof-plane cell parameter, observed for low thicknesses. The tail-to-tail polarization domains are indeed expected to repel each other and therefore lead to an increase of the overall cell parameters. The phenomenon will be less and less appreciable with increasing films thicknesses.

In order to further investigate the origin of the changes in the polarization and magnetization orientations during the early stage of GFO growth, we performed a chemical analysis. Atomically resolved quantitative elemental mappings of the film at the interface with the substrate were obtained by processing the STEM-EELS spectrum [Fig. 7]. Both the Ga and Fe maps in Fig. 7(b) are in perfect agreement with the expected positions of Ga majorly in the Ga1 (tetrahedral) sites, and Fe majorly in the three other (octahedral) sites [see Fig. 7(a)] [START_REF] Lefevre | [END_REF]. The zone of darker contrast, observed at the polarization reversal in Fig. 6(a), is due to the presence of Fe atoms only, on three consecutive cationic layers in this zone in Fig. 7(a). This results in a Ga depleted zone and, Ga being a heavier atom than Fe, to a relative darker contrast. Ti migrates from the STO substrate into the GFO film by an amount of about 6%, over the first 5 nm. The profile of O is very similar to that of Ti. It shows a depletion at the interface and in the very first nanometers of the deposited film (59 at. %, to be compared to the 60% expected for GFO), and then an increasing concentration up to the unit cell orientation reversal (65 at. %, to be compared to the 60% expected for GFO). After the polarization reversal region, the O content restores to the expected value and that of Ti becomes negligible.

When analyzing the profile of the Fe L 2,3 -edges spectra [Fig. 8], one can observe a significant variation between the very first deposited GFO and the rest of the film, in both 8]. Both the Fe excitation edge energy and the I (L 3 )/I (L 2 ) ratio are correlated to the Fe oxidation state [50,51]. Here, they indicate the presence of Fe 2+ in the early deposited layers of GFO. This is a second possible explanation to account for the increase of the c parameter observed for the very thin films. Indeed, Fe 2+ has a bigger radius than Fe 3+ , both in tetrahedral and octahedral sites (r Fe 2+ Oh = 0.78 nm > 0.645 nm = r Fe 3+ Oh = 0.78 nm, and r Fe 2+ T d = 0.63 nm > 0.49 nm = r Fe 3+ T d ) [52]. Both HR-TEM and EELS information can be combined to trace back the whole deposition process, from the very first steps of the growth. The following hypotheses may be formulated to describe the mechanism according to which GFO films grow onto STO substrates. At the early stages of the deposition, the Ti atoms of the STO substrate are attracted towards the oxidizing atmosphere at the surface of the sample and diffuse into the forming GFO film, always remaining as close as possible to the free oxygen-rich atmosphere. Such a diffusion phenomenon is promoted by the high temperature at which the sample is maintained during the growth (900 °C). Oxygen-driven cationic mobility has already been observed in other systems [53]. This results in an oxygen-deficient environment for the early grown GFO cells, and therefore to an orientation of their electric polarization towards the substrate. There is indeed a strong relationship between the oxygen concentration and the polarization orientation of polar films [54,55]. The control of the orientation of the polarization of a film through the oxygen partial pressure to which it is submitted has been shown by Wang et al. in ultrathin PbTiO 3 films [56]. Fe in the first GFO deposited layers is reduced to its +II valence state, as observed by EELS. Ga, which is less prone to reduction, tends to move away from this zone, which is therefore relatively enriched in Fe [see the Fe concentration in the first deposited nanometers in Fig. 7(b)]. The Ti migration from the substrate towards the oxidizing atmosphere of the surface continues as long as the electric conductance of the material allows it. It stops when the GFO film is 5-nm thick and becomes insulating enough to prevent any ionic mobility. The oxygen is then no more cornered by Ti, and the oxygen-rich atmosphere and its correlated negative charges allow a reversal of the polarization [57], which will now point towards the surface until the end of the growth.

Thickness dependent electric polarization orientation had already been observed [58]. The phenomenon was, however, ascribed to already well documented strain-driven effects [57,59,60]. Here the strain-relaxed nature of the GFO films excludes strain-driven effect on the observed polarization reversal within the first 5 nm of the films. Instead, it could have a chemistry-based mechanism and be related to the important ionic migration processes evidenced by HR-TEM and EELS at the substrate-film interface.

The peculiar composition of the first 5 nm of the films, evidenced here by the atomically resolved EELS study, could also account for the modification of the magnetic anisotropy of the GFO films from in-to out-of-plane. Some first principles calculations are currently under way to study the influence of the insertion of Ti in the GFO cell onto its magnetic properties.

IV. CONCLUSION

The pulsed laser deposition of thin films of the magnetoelectric multiferroic compound Ga (2-x) Fe x O 3 (x = 1.4) onto STO (111) substrates has been studied from the very early steps of the growth. The growth is initially 3D, because of the mismatch between the substrate and the film but rapidly becomes 2D. A rms roughness as low as 0.16 nm is observed for samples of thicker than 10 nm. The RHEED monitoring resolved one metal-oxygen-polyhedra layer deposition, corresponding to 1/4 th of a unit cell. All films showed an out-ofplane electric polarization directed towards the substrate for the first few nm and in the opposite direction for the rest of the thickness. Atomically resolved EELS measurements pointed at an ionic migration process driven by electrostatic effects, possibly explaining the magnetic properties for the ultrathin GFO films characterized by an out-of-plane anisotropy. This study therefore lays the ground for the use of multiferroic GFO1.4 thin films in spintronics heterostructures, such as spin-Hall-effect-based heavy-metal / ferrimagnetic oxide heterostructures. The 7-nm-thick sample, with its low rms roughness value and out-of-plane magnetization, is promising to allow efficient spin-current transmission from the heavy-metal (Pt for example) into the ferrimagnetic oxide. A configuration featuring both room temperature out-of-plane net magnetic moment and electric polarization is extremely rare; GFO is only the second oxide system unveiled up to now to display low rms roughness sub-10-nm-thick ferrimagnetic insulating films with out-of-plane magnetization. This opens new perspectives in the field of spin-orbit torque-based spintronics.

FIG. 1 .

 1 FIG. 1. (a) In situ RHEED patterns observed during the growth of GFO1.4 on STO (111). The patterns are taken along the [ 12 1] direction of the substrate. The visible lines for the bare substrate correspond to STO 10 1 reflections and harmonics. When the growth of GFO starts, some features appear corresponding to GFO 020 reflections and harmonics, the azimuth being along the GFO [100] direction (in the Pna2 1 space group). (b) AFM images of the surface of the samples.

FIG. 2 .

 2 FIG.2. Fluctuation of the intensity of a spot in the zero-order Laue zone during RHEED monitoring of the 7-nm GFO film growth. Each oscillation corresponds to the deposition of one fourth of a GFO cell, that is, to a layer of transition-metal/oxygen polyhedra.

  1 [001] axis, without any trace of spurious phase. The observation of Laue oscillations on the zoomed GFO (004) θ -2θ scans confirms high crystallinity and low roughness of the films [Fig. 3]. For the lowest thicknesses, the 00l peaks are shifted towards lower 2θ values, indicating a larger out-of-plane c parameter. The in-plane relationships between the STO (111) surface and the GFO (ab) face are determined from φ scans performed on both STO and GFO reflections (see Fig. S2 [26]). GFO may adopt three directions complying with the following epitaxial relationships with STO: [060] GFO (001) // [hkl] STO (111) with [hkl] equal to [2 20], [ 220] or [ 202]

  FIG.3. X-ray diffractograms of the deposited GFO thin films in the θ-2θ mode, with a focus on the GFO 004 peak, the vertical line at 2θ = 38.17 • indicates the angle observed for bulk GFO x = 1.4 (c = 0.9422 nm)[17].

FIG. 4 .

 4 FIG. 4. (a) OOP-SHG anisotropy of the 32 nm GFO film at a 30°tilt. The OOP-SHG is maximized along the projection of the polar axis, which is, along 90°, and exhibits a slight offset from this direction because of interference with the substrate-related SHG contribution. (b) Sample tilt dependence of the OOP SHG-component shown in (a). The sin 2 (x) dependence of the SHG intensity on the tilt angle confirms the OOP orientation of the spontaneous polarization in the material. (c) OOP-SHG at room temperature as function of GFO film thickness. The data point at 0 nm corresponds to measurements on the bare STO(111) substrate. The OOP-SHG component was extracted for each film by a sin 2 (x) fit of SHG tilt-dependence, as shown in (b).

FIG. 5 .

 5 FIG. 5. (a) Room temperature and (b) 10 K magnetic hysteresis loops measured with both in-plane and out-of-plane applied magnetic field for the 32-,11-,and 7-nm-thick samples.

FIG. 6 .

 6 FIG. 6. (a) HAADF HR STEM image of a cross section of the 32 nm GFO film deposited on STO (111) showing that GFO grows in its expected Pna2 1 structure from the growth start. A dark line delimitates a first 5-nm-thick GFO layer from the rest of the layer; (b) magnifications of zone 1, showing a polarization pointing towards the substrate, and zone 2, showing a polarization pointing outwards from the substrate. The GFO unit cell and expected atomic pattern for an observation along the [100] zone axis are superimposed on the magnifications; the orientation of the cell is recognizable from the convex or concave feature formed by the four Fe2 and Ga2 octahedral sites in a row (zoomed on the right of the figure).

FIG. 7 .

 7 FIG. 7. STEM-EELS analysis of the interface between the STO(111) substrate and the GFO (32 nm) thin film. (a) HAADF survey image and zoom recalling the atomic positions and orientation of the cells on each side of the polarization boundary. (b) Atomically resolved quantitative elemental maps for Fe, Ga, O, and Ti with elemental profiles integrated over the zone of interest (indicated by white dotted rectangles on the maps).

FIG. 8 .

 8 FIG. 8. (a) HAADF HR STEM image of the 32-nm GFO film highlighting four zones at various depths in the films, and (b) the Fe L 2,3 EELS profiles integrated on these zones.

ACKNOWLEDGMENTS

This work was funded by the French National Research Agency (ANR) through the ANR MISSION ANR-18-CE-CE24-0008-01 and through the Programme d'Investissement d'Avenir under Contract No. ANR-11-LABX-0058_NIE within the Investissement d'Avenir program ANR-10-IDEX-0002-02. S.H. also acknowledges financial support from ANR through Grant No. ANR-17-EURE-0024 EUR QMat. This research was supported by the EU European Research Council (Advanced Grant No. 694955-INSEETO) and by the Swiss National Science Foundation under Project No. 200021_-178825. The authors wish to thank D. Troadec (IEMN, Lille, France) and A.-M. Blanchenet (UMET, Lille, France) for the preparation of the FIB lamellae, as well as the the XRD, MEB-CRO, and TEM platforms of the IPCMS.


Cherif Baldé
email: cbalde@univ-zig.sn

Mame Seyni Sylla

Cédric Desplanches

Guillaume Chastanet

Chérif Baldé

High-spin → low-spin relaxation kinetics and metal dilution effects in 1D spin-crossover chain compounds

Keywords:

à la diffusion de documents scientifiques de niveau recherche, publiés ou non, émanant des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

Introduction

The recent developments in technology in which molecules or assemblies of molecules are used for information processing has enabled scientists to investigate their potential in several areas, namely for the development of display and memory devices [1]. Among them, the materials that undergo a spin crossover (SCO) phenomena upon application of external stimuli, such as tempera ture, pressure, light and electric field [2 6], have received a high interest within the construction of molecular devices with optical switch or/and magneto optical storage properties [7 17]. In this context, SCO coordination compounds of a 3d 4 to 3d 7 electronic configuration, displaying a central metallic ion in an octahedral environment, have been the focus of research for nearly 80 years [1 6]. Several aspects of the SCO phenomenon have been reviewed in the past [2 4,16 20]. Of these, by far the majority are pseudo octahedral 3d 6 iron(II) complexes, thus, as a function of the ligand field strength, the six electrons of the iron(II) ion can occupy the 3d orbitals in two different ways, giving rise to two different stable states: a diamagnetic low spin (LS) electronic configuration state t 2g 6 e g 0 ( 1 A 1g , S = 0) in the case of a strong ligand field and a paramag netic high spin (HS) electronic state t 2g 4 e g 2 ( 5 T 2g , S = 2) when the ligand field is weak. In the case of an intermediate ligand field, the two spin states may become so close in energy that a small external perturbation can switch one state into the other. LS M HS conversions are accompanied by profound changes in all the properties that depend on the distribution of the 3d valence electrons. Besides the most obvious change of magnetic properties [21], including switchable diamagnetism M paramagnetism [10], the reversible change of color [10], refractive index [22], electrical conductivity, [23] luminescence, [START_REF] Shepherd | Spin-Crossover Materials, Properties and Applications[END_REF] non linear optical [START_REF] Lacroix | [END_REF] and mechanical properties [26] may accompany the SCO. The bistabil ity of SCO materials can be addressed with the application of var ious external stimuli, like temperature [START_REF] Gütlich | [END_REF]27], pressure [28 30], magnetic field [31,32] or light irradiation [START_REF] Gütlich | [END_REF]27,33,34]. The optical switching is particularly regarded since ultrafast switches (fem toseconds) can be achieved [15]. Light can reversibly write and erase the LS and HS states according to the direct or reverse light induced excited spin state trapping (LIESST) effects [34 (d),35]. In 1998, Létard et al. introduced the T(LIESST) measurement, which aims to estimate the limiting temperature above which the photoinduced HS metastable state is erased [19,36]. From the determination of the T(LIESST) value of several compounds, a comparison has been made in order to draw trends for chemists to increase the lifetime of the photo induced state [37]. In particular, the correlation between T(LIESST) and T 1/2 was evidenced based on a [T 1/2 ; T(LIESST)] database. From this [T 1/2 ; T (LIESST)] database [36,37], Létard et al. observed a linear depen dence between T(LIESST) and T 1/2 and extracted the empirical rela tionship T(LIESST) = T 0 0.3 T 1/2 where T 0 is the initial value of the linear function. This phenomenological relationship appears to This result contrasts with that obtained for the zero dimen sional mononuclear compounds previously studied. Indeed, it was shown that Zn(II) dilution does not lead to any significant change of T(LIESST) [46]. In the polymeric systems under study, the metal dilution does not follow this trend and moreover it does not exactly follow the T(LIESST) versus T 1/2 empirical relationship. Even if the variations look minor, they are still not understood.

One clue to explain such behavior should lie in the relaxation rate, k 1 . Indeed, despite the great structural similarity between the two series of compounds endi and btzp, k 1 has values that vary in magnitude. Hauser, on the basis of the SCC approach (Single Configurational Coordinate), predicted values of k 1 of the order of 10 6 10 9 s 1 [START_REF] Gütlich | [END_REF]27,33,54]. In SCC theory, the Fe ligand distance is taken as a reaction coordinate with a single vibration mode. Regarding the values of k 1 , especially those relating to the endi and btzp series, it is reasonable to assume that a single reaction coordinate may not be sufficient to correctly describe the HS ? LS transition. Therefore, the SCC approach may be insufficient, especially to obtain values of k 1 of the order of 10 2 s 1 . It becomes essential to consider a system that considers several modes of vibration to describe correctly the HS ? LS relaxation. In 2006, Hauser has elsewhere proposed a model with two modes of vibration [53]. Compared to a single mode system, a system involving a low distortion mode vibration energy allows to the anticipation of low values of k 1 . Interestingly, from a crystallo graphic point of view, the multimode theory, which takes into account the concept of distortion, is the most suitable for describ ing the spin crossover phenomenon. Indeed, a direct relationship has been established between the variation of the octahedron [FeN 6 ] distortion, T 1/2 and T(LIESST) [20,55].

Conclusions

In the present work, we have investigated the effects of metal dilution on thermal properties and LIESST relaxation kinetics time/temperature dependence for [Fe x Zn 1 x (btzp) 3 ](BF 4 ) 2 and [Fe x Zn 1 x (endi) 3 ](BF 4 ) 2 SCO materials. As a first approximation, the results related to the compounds of the two series obey roughly a linear relationship between T(LIESST) and T 1/2 . Contrary to the dilution series of zero dimensional mononuclear compounds, the iron atoms in the compounds under study are not located in iso lated molecules, linked only by van der Waals forces or p p stack ing, but are linked by covalent bonds. This can justify the somewhat different behaviors between the systems studied in this article and the previously studied dilution studies on zero dimen sional mononuclear compounds.

As far as kinetic parameters governing the metastable HS ? LS relaxation process (activation energy, frequency factor) parameters are concerned, it appears that dilution increases the energy barrier, whereas the derived frequency factor is increased in the case of the [Fe x Zn 1 x (endi) 3 ](BF 4 ) 2 series and remains essentially constant in the case of the [Fe x Zn 1 x (btzp) 3 ](BF 4 ) 2 series. It seems that, in both cases, the increased activation energy is not completely compen sated by the higher frequency factor, so that globally T(LIESST) is slightly increased with dilution.

Acknowledgments

The authors acknowledge the Department of Cooperation and Cultural Action Embassy of France to Dakar Sénégal, the University of Ziguinchor and its Department of Cooperation and Research, the Ministry of Higher Education and Research of Sénégal (Program FIRST) and Campus France. This work was also supported by the University of Bordeaux, France, the CNRS, the Region Nouvelle Aquitaine and by the LabEx AMADEus (ANR 10 LABX 42) within IdEx Bordeaux (ANR 10 IDEX 03 02), i.e. the Investissements d'Avenir programme of the French government managed by the Agence Nationale de la Recherche. The ANR is also warmly acknowledged (ANR femtomat n°13 BS04 002). The authors also deeply thank the French network GdR MCM2 (Magnétisme et Com mutation Moléculaires) for the fruitful discussions the regular meetings provide. Finally, we warmly thank all the collaborators, visitors and experts for their confidence, interest and criticisms that helped a lot in the building of the database and the under standing of the phenomenon.

Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org /10.1016/j.poly.2018.11.046.


Shian Guan

Aline Rougier

Matthew R Suchomel

Nicolas Penin

Kadiali Bodiang

Manuel Gaudon

Geometric considerations of the monoclinic -rutile structural transition in VO 2

The mechanism of the displacive phase transition in VO2 near the transition temperature is discussed in terms of a geometrical approach, combining simple calculations based on the Brown's band valence model and in situ X-ray diffraction experimental results. Considering that the structural origin is well linked to the electrostatic potential optimization as in a Peierls model, our geometrical calculations and experimental studies are in agreement and suggest that VO2 phase transition is the consequence of very short atomic shifts mainly associated to a decrease of the 2 nd sphere coulombic interactions. Hence, at a given temperature, the allotropic form (monoclinic versus rutile form) offering the largest unit-cell volume is stabilized over the lower unit-cell volume allotropic, while the transition occurs at the intercept of the unit cell variation versus temperature of the two forms, which exhibit significantly different thermal expansion coefficients.

Introduction

The origin of the driving force(s) behind the metal-insulator transition (MIT) of VO2 has been the subject of controversy for decades. Structurally, the MIT transition at approximately 70 °C is accompanied by a change from a high-temperature tetragonal rutile structure (S.G. P42/mnm) to a lowtemperature monoclinic (M) form (S.G. P21/c) characterized by a structural twist and V 4+ cations shift along the rutile c axis to form V-V pairs (homopolar bonds) [1]. From a simplistic point of view considering the orthonormal referential constituted by the octahedral cage edges, the key aspect of the transition is an offcentring of vanadium ions along (111), with an antiferroelastic correlation. A crystal description of the transition was first reported by Goodenough et al. [2], associating the lattice distortion (pairing and tilting) in the VO2 (M) phase with an increase of the antibonding 3d band above the Fermi level. Herein lies the controversial point; whether the phase transition originates from the crystal distortion and unit-cell doubling (Peierls mechanism), or from the opening of a correlation gap due to the presence of electron-electron correlations (Mott mechanism) [3].

Recent structural calculations (Peierls approach) based on density functional theory within the local density approximation tends to predict the monoclinic phase as the most stable [4][5][6][7][8][9]. In a complex structural approach, i.e. using the Buckingham potential to describe the short-range interaction (V-O and O-O bonds), a Morse potential for V-V interaction (at the origin of the octahedral chain distortion) and additional optimization techniques, Woodley [10] concludes that monoclinic and rutile free energies cross at nearly 300K, in favour of a Peierls transition. Cavalleri et al. [11] showed evidence of a structurally driven transition with the use of ultrafast pump-probe spectroscopy. Conversely, several authors argue for a Mott-driven transition [12][13][14] based on experimental parameters such as switching time using femto-laser sources [12,13] or by measuring the electromagnetic response of a VO2 thin film by scattering scanning near-field infrared microscopy. [14] Current interpretations seem to propose a combination of both mechanisms simultaneously; meanwhile the appearance of new phases during the structural phase transition (well supported by nanoscale imaging techniques and/or optical/electronic spectroscopy) makes it increasingly critical to understand the true nature of this transition [15][16][17][18][19]. Despite the complex nature of previously reported theoretical structure calculations, a basic interpretation of the transition in terms of fundamental geometrical factors such as unit-cell volume and bond distances is often neglected in these studies. However, these parameters are linked to the thermal expansion and coulombic interactions of both VO2 allotropic forms and are thus potential significant driving forces to consider near the boundary of the structural transition.

Within the limit of a simple geometrical approach, we describe here a new approach for evaluating the main driving forces involved in the MIT transition of VO2. This approach compares measured unit-cell parameters to basic calculations for the cell volume and electrostatic potential of the 1 st and 2 nd sphere anion/cation pairs in monoclinic and rutile structure types near the phase transition. Calculations are based on the well-established bond valence model, [20] and can be rapidly performed with a standard spreadsheet software package. The experimental temperature dependence of monoclinic and rutile phase unit-cell parameters near the transition temperature for this study were obtained from in-situ synchrotron powder diffraction measurements of VO2 samples, the synthesis of which was previously reported. [21] Calculations Using Brown's bond valence model [20] to evaluate all atom positions in a structural unit cell with a basic set of geometric Please do not adjust margins Please do not adjust margins equations is a powerful tool which has been previously used with success to predict crystal structure anisotropy (as in Würtzite compounds [22]) or even multiple phase transitions (as in the BaTiO3 perovskite [23]). In this type of an approach, atomic positions are typically relaxed to minimize the electrostatic potential between first neighbouring cations and anions pairs. The constraints of the Brown equations result in convergence towards a single solution, i.e. the valence of the vanadium ion deduced from the bond valence model must be kept equal to 4. For the binary ionic compound as VO2, the Brown equation for each V-O pair relating bond lengths and bond valences is thus BV = exp[(r0 -dij)/0.37] with BV: bond valence, r0 = 1.78 Å (V-O bond empirical parameter given by Brown tables) and dij = V-O bond length in Å. In the VO2 crystal structure, as the vanadium cations occupies an octahedral coordination site, six bonds are summed up for the valence of the vanadium cations.

The initial step of this geometrical approach was to consider positional shifts of vanadium ions in the octahedral cages of the rutile (tetragonal) vs. monoclinic structures and to examine if V 4+ decentring is associated with an increase of the surrounding 1 st sphere coulombic interaction E(1st sphere) (i. e., a stabilization effect) and thus connected to the transition between these polymorphic forms. The parameters used to describe the rutile structure are shown in Fig. 1a andc, the unitcell parameters; x, describing the oxygen position at (x, x, 0); l', the edge length of the octahedral cages along the octahedral chains; l", the shared edge length of the octahedral chains; and h, the half-height of the octahedral cage (note that h is identical to the distance of apical V-O bonds, while the other 4 equivalent equatorial V-O bonds are noted as r in Fig. 1). The equations linking these parameters are: h = 2.a.x ; r = [( ½.l') 2 + ( ½.l'') 2 ] ½ ; l' = 𝑐 ; l'' = 2.2.(0.5 -x) and Vcell = a 2 .c. Considering an ionic approach, the first sphere potential is defined as E(1 st sphere) = K.(2/h + 4/r), where K is a constant.

A simplified geometric description of the pseudo-monoclinic VO2 structure can be obtained by considering only V 4+ decentring along [111] direction as oxygen site displacements in the distorted octahedral cages are negligible. This leads to the new set of equations: δ is the amplitude of the decentring in (111) direction; α is the angle of the [111] axis with respect to the equatorial plane of the octahedral cage;

h' = (h 2 + δ 2 -2.h.δ".sinα) ½ ; h'' = (h 2 + δ 2 + 2.h.δ".sinα) ½ ; r' = [(½l') 2 + (½l'') 2 + δ 2 -δ.l'.cosα] ½ ; r'' = [(½l') 2 + (½l'') 2 + δ 2 + δ.l'.cosα] ½ ; l' = 𝑐; l'' = 22(0.5 -x) ; Vcell = a 2 .
c. The first sphere potential becomes E (1 st sphere) = K.(1/h' + 1/h'' + 2/r' + 2/r''). c, a and x as have been considered as variables in order to maximize (considering positive terms for crystal energies) the E (1 st sphere) versus δ. From these calculation, E (1 st sphere) and the optimised unit cell volume can be plotted versus cation decentering (Fig. 2). Please do not adjust margins

Please do not adjust margins

As shown in Fig. 2a, the decrease of E(1 st sphere) with δ shows that our ionic approach fails to explain the rutile to monoclinic phase transformation. With a higher absolute energy, the tetragonal environment appears most stable in these geometrical calculations, with any monoclinic-type decentring (δ) appears to reduce the overall energy. However, based on the same calculations, it is noted that the unit-cell volume is predicted to slightly increase with vanadium cation decentring (Fig. 2b), suggesting that the transition may be eased by an expansion of the cell volume. This later result is somewhat unexpected, as crystal energies are normally to increase (stability increase) with crystal density (i.e. with shortening bond lengths).

In a second step, the origin of the tetragonal distortion is evaluated, i.e. the E(1 st sphere) and unit-cell volume of tetragonal phase are calculated versus the c/a ratio. The goal of this analysis is to probe why the VO2 rutile phase does not exhibit cubic symmetry; a discussion that could in principle be extended to all rutile compounds typically exhibiting a tetragonal distortion. The isotropic point, defined by a c/a ratio equal to 0.586 (i.e. 2.(2-1) ), corresponds to a "cubic-rutile form". For the calculation, the c/a value is fixed while a and x are variables adjusted in order to minimize the E(1 st sphere). Regardless of the tetragonal distortion a stable minimum for the E(1 st sphere) can be determined (Fig. 3a). Despite the tetragonal distortion of the unit-cell and consequently of the octahedral cages also, the six V-O bond lengths remain constant. Indeed, the distortion of the equatorial plane of the octahedral site can be followed via the evolution of the l' and l'' parameters (Fig. 3c), but this distortion occurs such as the r bond lengths remain equal to the h length, whatever the c/a ratio (fixed values for r = h = 1.930 Å). Thus, the tetragonal distortion from the cubic form of the rutile structure is apparently not due to a minimization of the E(1 st sphere). In addition, it is noted that the plot of unit-cell volume versus c/a adopts a bell-shape curve with a maximum at approximately c/a ≈ 0.626 (Fig. 3b), in good agreement with experimental measured values for the rutile form of VO2. Moreover, by fixing ab initio the value of h to 1.930 Å, since no deviation of this parameter was found with varying c/a ratio, it is possible to treat all a, c (and so c/a) and x parameters as variables, and to perform a calculation in order to maximize the unit-cell volume. This optimization leads to a crystal defined by a c/a ratio of 0.626, with a = 5.754 Å, c = 2.870 Å, with characteristic octahedral lengths of l' = 2.87 Å, l'' = 2.57 Å and r = h = 1.93 Å. Interestingly, these calculated lengths are very similar to the experimental parameters given in the literature for VO2 (R) phase [24][25][26][27]. The stability of the rutile form is therefore strongly linked to a maximization of the cell volume while simultaneously maintaining the E(1 st sphere). Indeed, the cell volume maximization is directly correlated to the minimization of the E(2 nd sphere), i.e. the decrease of the anion-anion or cation-cation coulombic repulsions. Moreover, these conclusions are supported by an additional calculation performed to minimize the E(2 nd sphere) using the set of shortest V-V and O-O bond lengths. In a manner analogous to that detailed above, the bond distances can be described as follows: 2 V-V bonds and the 2 O-O bonds = 2.l', the 8 V-V bonds = [(h + ½.l'') 2 + (½.l') 2 ] 1/2 , the 8 O-O bonds = (h 2 + r 2 ) 1/2 + 2.l', the 2 last O-O bonds = [(2.x.a) 2 + (a -2.x.a) 2 ] 1/2 . Then the E(2 nd sphere) minimization calculations are performed with a, c and x as variables, and h fixed to 1.93 Å (as above), the overall crystal parameters results are nearly identical to that found with the described cell volume maximization calculations.

Hence our basic geometric calculation clearly shows that by considering only a Peierls origin for the VO2 phase transition, with a pure ionic bond approach, the tetragonal distortion of the rutile phase is well explained by an energy minimization. The tetragonal distortion allows the decrease of the crystal energy which can be approximated as E(crystal) = E(1 st sphere) -E(2 nd sphere). The originality of the system is that the tetragonal distortion does not impact of E(1 st sphere); thus, the optimization of the electrostatic potential is governed by E(2 nd sphere) associated with cell volume maximization. This result on the driving forces of the tetragonal distortion origin has to be correlated to our first calculation showing that the monoclinic to rutile phase transition is not associated with an increase of the cation-anion interaction potential (1 st sphere maximization) but may be promoted by an increase in the unit cell volume (2 nd sphere minimization). This journal is © The Royal Society of Chemistry 20xx

Please do not adjust margins Please do not adjust margins

Experimental Results

Synthesis details of the VO2 (M) powder, which showed minor trace of a V3O7 impurity, has been described previously [1]. The phase transition as a function of temperature was followed by in-situ powder diffraction. High-resolution synchrotron X-ray powder diffraction (HR-SPD) data of VO2 have been recorded on both heating and cooling with 20 °C temperature steps between 280 and 380 °C using beamline 11-BM at the Advanced Photon Source (APS), Argonne National Laboratory. For each temperature scan, data were collected over a 0.5-35° 2θ range with a 0.001° step size using a wavelength of λ = 0.41417 Å; the powder sample being sealed in a 0.8 mm diameter Kapton capillary and was spun at 60 Hz during data collection (Fig. 4a). parameters were performed by profile matching and cell fractions from Rietveld refinements using the TOPAS software package, as illustrated on the low (monoclinic) and the high (tetragonal) temperature forms recorded at 5°C (Fig. 4b) and 105°C (Fig. 4c), respectively.

Determination of VO2 unit cell

The evolution of the rutile and monoclinic phase fractions over the studied temperature range are summarized in Fig. 5a. The unit-cell parameters of both rutile and monoclinic forms were refined, and normalized unit cell volumes vs. temperature are plotted to support the discussion (Fig. 5b). The rutile to monoclinic phase transition is described by a hysteresis loop, centred on 65-70 °C with a width of about 20°C : it can be seen that at 65°C the percentage of rutile phase / monoclinic phase ratio in the heating mode is about the same than the one calculated at 45°C in the cooling mode. in good agreement with past studies [2,[12][13][14][15][16][17][18][19]. The evolution of both cell parameters versus temperature shows that the volumetric coefficients of thermal expansion for the two phases are significantly different. Using a simple linear slope fit of the low temperature (monoclinic) and high temperature (tetragonal) unit cell volumes, we can estimate the volumetric thermal expansion coefficients (Vol.) using Vol. = (V/T)/V, where V/T is simply the linear fit slope. For the low temperature (monoclinic) phase, Vol. = 20.3x10 -6 K -1 and for high temperature (tetragonal) phase, Vol. = 37.2x10 -6 K -1 . Note that for ideal anisotropic materials, the vol. value is 3 times larger than the linear thermal expansion coefficient; these estimated volumetric thermal expansion coefficients are in good agreement with the range of VO2 literature values [28].

The rutile form exhibits a higher αVol than the monoclinic one. On the transition temperature, the volume of the two Please do not adjust margins Please do not adjust margins phases is similar, i.e. the intersection point of the unit-cell evolution is located at the transition temperature. Consequently, due to the hysteresis phenomenon linked to the 1 st order character of the phase transition, and due to the mismatch between the monoclinic thermal expansion coefficient (TEC) and the rutile TEC, both in cooling and heating modes, the transition is associated with a slight volume expansion. The volume expansion is experimentally with an amplitude of about 0.01 Å 3 in cooling mode: measured between the monoclinic (blue triangle) and rutile phase (blue square) at 45°C, and about 0.02 Å 3 in the heating mode: measured between the monoclinic (blue triangle) and rutile phase (blue square) at 45°C. Thus, the volume expansion in both heating and cooling mode produced by the phase transition is roughly 0.025% volume (or 0.3% of bond length variation). It is apparently a minor change, but it well governs the phase transition. The stable form is the one with the highest volume (per formula unit Z). Such a transition scheme is very different to most other 1 st order transitions in crystalline solids for which the volume difference between the two crystalline forms around the transition is the main aspect (as in the textbook austenite-martensitic example [29][30] or for the AMO4 phase transitions [31][32][33]). From these experiments, it appears that the driving force of the transition comes from this volume difference: a greater unit-cell volume leads to higher crystal structure stability.

Conclusion

This work considers the advantage of simplistic geometric examination of the much studied VO2 transition, in a marked contrast to the very complex literature debates on competing Peierls and Mott-type origins. From a solid-state chemistry point of view, there is often considerable benefit to interpret such transitions in terms of a structural origin and an optimization of electrostatic potentials. Through the positively correlated calculations and experimental data examined in this work, it is shown that the VO2 rutile to monoclinic phase transition is likely the consequence of very short atomic shifts, leading to non-significant (at least non-dominant) changes on the charge interactions of the cations and anions in first neighbouring, however allowing a decrease of the 2 nd sphere coulombic interactions.

Figure 1 :

 1 Figure 1: Rutile VO2 structure with l' and l'' are the two edges lengths of the octahedral equatorial rectangle, h the half height of the octahedron, and oxygen in (x, x, 0) Wyckoff position. Vanadium octahedral sites and their associated parameters in rutile, monoclinic or "simulated" monoclinic structure: i.e. considering a rutile octahedral cage with a cation  decentring in (111) direction. (r' and r'' represent two V-O bonds distance in the plane after V 4+ decentring)

Figure 2 :

 2 Figure 2: Effect of the consideration of  decentring amplitude in the simulated monoclinic structure on the electrostatic potential (1 st sphere energy) (a) and the unitcell volume (b).

Figure 3 :

 3 Figure 3: Effect of the c/a tetragonal distortion on the electrostatic potential (1 st sphere energy) (a); the unit-cell volume (b) and the l' and l'' octahedral edges lengths (c).

Figure 4 :

 4 Figure 4: Synchrotron powder diffraction pattern over the 5°C105°C temperature range showing the reversible transition between monoclinic phase (P21/c) and rutile (P42/mnm) structures on heating and cooling (a). Profile fits of VO2 HR-SPD data collected at 5 °C (b) and 105 °C (c), showing the observed (red points) and calculated (black line) intensities as a function of 2θ angle. The difference profile and unit cell hkl reflections are shown below as a blue line and green ticks, respectively. The exclude 2θ region (approximately 8 -9°) is indicated by light grey data points.

Figure 2 :

 2 Figure 2: (a) Phase fraction (weight %) of the monoclinic phase as a function of temperature on heating and cooling, (b) VO2 unit cell volume (per formula unit Z) as a function of temperature on heating and cooling.

  . CNRS, Univ. Bordeaux, Bordeaux INP, ICMCB, UMR 5026, F-33600 Pessac, France This journal is © The Royal Society of Chemistry 20xx

a

This journal is © The Royal Society of Chemistry 20xx J. Name., 2013, 00, 1-3 | 3

Acknowledgements

Use of the Advanced Photon Source at Argonne National Laboratory was supported by the U. S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357.

Conflicts of interest

There are no conflicts to declare.


Axel Martin

Laure Beven
email: laure.beven@inra.fr

Marie-Hélène Delville

Christine Grauby-Heywang
email: christine.grauby-heywang@u-bordeaux.fr

Marion Mathelié-Guinlet

Touria Cohen-Bouhacina
email: touria.cohen-bouhacina@u-bordeaux.fr

Ibtissem Gammoudi

Laure Bé- Ven

Silica nanoparticles

Keywords: immunosensor, cyclic voltammetry, E. coli, nanoparticles

come

The detection of pathogenic bacteria is of critical importance for health and safety issues in clinical diagnosis, food industry, water and environment quality, as well as struggle against potential bioterrorism [START_REF] Lazcka | Pathogen detection: A perspective of traditional methods and biosensors[END_REF]. Traditional and conventional methods for bacteria identification and detection are mainly based on culture and colony counting [START_REF] Ayçiçek | Assessment of the bacterial contamination on hands of hospital food handlers[END_REF], polymerase chain reaction [START_REF] Jeong | A novel multiplex PCR assay for rapid and simultaneous detection of five pathogenic bacteria: Escherichia coli O157:H7, Salmonella, Staphylococcus aureus, Listeria monocytogenes, and Vibrio parahaemolyticus[END_REF] and immunology tests [START_REF] Galikowska | Specific detection of Salmonella enterica and Escherichia coli strains by using ELISA with bacteriophages as recognition agents[END_REF]. Despite their high sensitivity, most of them require pre-enrichment of bacteria in samples and/or DNA-amplification and, thus, tend to be laborious, time consuming and expensive. Consequently, many efforts have been made to reduce the cost and detection time while improving the reliability, sensitivity and robustness of new detection techniques. In this context, biosensors, inspired from immunology tests like ELISA one, have raised an incredible interest as they allow this detection in a single autonomous and integrated device [START_REF] Minh | Les biocapteurs : principes, construction et applications[END_REF][START_REF] Taylor | Biomolecular sensors[END_REF][START_REF] Grieshaber | Electrochemical biosensors-sensor principles and architectures[END_REF]. Simultaneously, nanotechnologies have experienced, over the last decades, a fast and growing development, especially in biomedical fields [START_REF] Duncan | Nanomedicine(s) under the Microscope[END_REF][START_REF] Gao | Targeted drug delivery system for platinumbased anticancer drugs[END_REF][START_REF] Goesmann | Nanoparticulate Functional Materials[END_REF]. Combining the physical properties and high surface-to-volume ratios of nanoparticles (NPs) to surface modification and composition of transducers, the search for highly sensitive and selective biosensors can be strongly renewed. Indeed, NPs have already been successfully applied to detect molecules of biological interest such as metallic ions, nucleic acids and proteins [START_REF] Mani | Ultrasensitive Immunosensor for Cancer Biomarker Proteins Using Gold Nanoparticle Film Electrodes and Multienzyme-Particle Amplification[END_REF]. Nowadays, a worldwide threat concerns pathogens that can cause serious, even fatal, illnesses through the contamination of food, water or medical tools, such those induced by the well-known E. coli O157:H7, that is susceptible to even develop resistance to current antibiotic therapies [START_REF] Chemburu | Detection of pathogenic bacteria in food samples using highly-dispersed carbon particles[END_REF][START_REF] Van Den Bogaard | Epidemiology of resistance to antibiotics: links between animals and humans[END_REF]. Regarding this issue, research on biosensors for the detection and elimination of E. coli bacteria is fairly extensive not only because of the wide range of existing transducers but also because of the broad spectrum of NPs properties especially developed for this purpose [START_REF] Sanvicens | Nanoparticle-based biosensors for detection of pathogenic bacteria[END_REF]. Optical biosensors, essentially based on Surface Plasmon Resonance (SPR), and fluorescence measurements have been successfully applied to the detection of E. coli, with limits of detection (LOD) as low as 10 2 -10 3 CFU/mL [START_REF] Sun | Optical transduction of E. Coli O157:H7 concentration by using the enhanced Goos-Hänchen shift[END_REF], similar to electrochemical and gravimetric biosensors [START_REF] Chowdhury | Label free polyaniline based impedimetric biosensor for detection of E. coli O157:H7 Bacteria[END_REF]. Indeed, these LODs are strongly improved when integrating NPs as shown, for instance, by Joung et al. who designed an SPR biosensor for E. coli, based on gold NPs, and reduced its LOD by three orders of magnitude [START_REF] Joung | High sensitivity detection of 16s rRNA using peptide nucleic acid probes and a surface plasmon resonance biosensor[END_REF]. A milestone was reached with fluorescein-enriched silica NPs associated with magnetic Fe3O4-NPs that allow detecting 3 CFU/mL within 75 min. [START_REF] Hu | A novel biosensor for Escherichia coli O157:H7 based on fluorescein-releasable biolabels[END_REF]. Despite such a high sensitivity, optical tools require expansive setups and analyses of experts which are not really convenient for wide use. Gravimetric biosensors, mainly based on Quartz Crystal Microbalance, coupled with gold or iron oxide NPs are easier to use and can detect as low as 10² CFU/mL [START_REF] Mao | A nanoparticle amplification based quartz crystal microbalance DNA sensor for detection of Escherichia coli O157:H7[END_REF]. Compared to these tools, electrochemical biosensors appear to be the least restrictive, since they offer portability, low cost and easy manipulation. For instance, Setterington and Alocilja took advantage of Cyclic Voltammetry (CV) to detect down to 6 CFU/mL within 1 h thanks to immuno-functionalized magnetic/polyaniline core/shell NPs [START_REF] Setterington | Electrochemical Biosensor for Rapid and Sensitive Detection of Magnetically Extracted Bacterial Pathogens[END_REF]. In this work, we present the first results on the design of an electrochemical biosensor, in which the presence of silica-NPs, functionalized with specific polyclonal antibodies (Abs), is used to target and bind E. coli. Subsequently, the interaction inhibits the current response of the transducer, observed by CV. We highlight the existence of straightforward monitoring indicators for E. coli detection and quantification that can easily be screen-displayed on a future point of care device. Therewith, this new biosensor can be applicable to a wide variety of targeted organisms only by changing the bioreceptor grafted on the NPs. Although optimizations are still in progress in terms of NPs and bioreceptor concentrations, according to the expected LOD (in the range of 10 3 to 10 6 CFU/mL), the time of detection can be lowered to only 5 min., the best LOD of 10 3 CFU/mL being achieved in only 30 min. Such characteristics and tunable biosensor architecture pave the way for industrial purposes.

Materials -

Gram negative E. coli bacteria (MRE 162 strain) were a kind gift from the Centre d'Etudes du Bouchet, DGA (Direction Générale de l'Armement, France). Polyclonal antibodies (Abs) against E. coli were purchased from Thermo Scientific. Bovine Serum Albumin (BSA), ethanol, ammonium hydroxide, L-arginine, tetraethylorthosilicate (TEOS) and aminopropyl triethoxysilane (APS), used for silica nanoparticles (NPs) synthesis, were purchased from Sigma Aldrich. Polyelectrolytes, used as a sublayer for NPs fixation in the biosensor, respectively cationic poly (allylamine hydrochloride) (PAH, Mw = 56000) and anionic polystyrene sulfonate (PSS, Mw = 70000), were purchased from Sigma Aldrich. K3Fe(CN)6, K4Fe(CN)6.3H2O and phosphate buffered saline (PBS, 10 mM, 0.137 M NaCl, pH 7.4), used for cyclic voltammetry measurements, were also purchased from Sigma-Aldrich. Millipore ultrapure water (pH 5.5, resistivity > 18.2 MΩ.cm) was used for NPs cleaning and for preparing all the buffers. 2.2 Electrochemical measurements -Cyclic voltammetry (CV) measurements were performed with a 910 PSTAT mini from Metrohm and data were collected with the associated software PSTAT. Data analysis was controlled with a Matlab routine, estimating the oxidation and reduction currents and the associated potentials. Gold screen printed electrodes (4 mm in diameter), with carbon auxiliary and silver reference electrodes, also purchased from Metrohm, were used as sensors. Before any solution was applied, sensors were rinsed with ultrapure water and allowed to dry. Sensors were disposed off, after a set of bacteria detections. CV measurements were performed, at room temperature, in PBS in presence of 10 mM of K4Fe(CN)6 redox probe to estimate the blocking properties of the biolayers versus electron-transfer process. CVs were carried out from -0.4 V to 0.6 V, at a rate of 100 mV/s, those presented in this work being acquired after the stabilization of the system. 2.3. Preparation of SiO2-NPs + solutions -Synthesis of negatively charged home-made SiO2-NPs -was carried out following synthesis developed by Hartlen et al. [START_REF] Hartlen | Facile Preparation of Highly Monodisperse Small Silica Spheres (15 to >200 nm) Suitable for Colloidal Templating and Formation of Ordered Arrays[END_REF]. The first step consisted in synthesizing SiO2 seeds. For this purpose, 100 mL of a 6 mM L-arginine water solution were introduced in a vial at 60°C, under magnetic stirring (150 rpm). After temperature stabilization, 10 mL of TEOS were slowly added so as to respect the interface. Seeds of about 30 nm in diameter were obtained after complete consumption of TEOS. The second step consisted in the growth of these seeds. Typically, 10 mL of SiO2 seeds were introduced in a mixture of 455 mL of ethanol and 35 mL of ammonium hydroxide in water solution ([NH3] = 1 M), under stirring (700 rpm). TEOS was then introduced in the amount calculated to get the targeted diameter (100 nm in this study), at a rate of 0.5 mL/h [START_REF] Gaudon | Thermochromic Phase Transition on CuMo 0.9 W 0.1 O 4 @SiO 2 Core-Shell Particles[END_REF]. Surface modification was performed by adding APS (six equivalents assuming a rate of 2-3 µmol/m² of active OH groups at the surface of the silica particles) to the SiO2-NPs -solution; the mixture was then stirred for 48 h in a water bath at room temperature, and then for 2 h at 100°C. Final SiO2-NPs + were washed using ultrapure water and ethanol and three cycles of centrifugation (13,000 rpm for 20 min.). SiO2-NPs + were finally suspended and kept in ultrapure water. They were sonicated before any experiment to prevent their aggregation.

Preparation of bacteria suspensions -E. coli cells

were grown on LB agar plates and incubated for 16 h at 37 °C. After scraping, bacteria were suspended in PBS to obtain an optical density of 1. Serial dilutions were then made to adjust their concentration. 10 µL of each dilution were then spread on the biosensor, previously functionalized by the appropriate systems, described in the "Results and Discussion" section. Before any CV, bacteria were incubated 5, 15 or 30 min. on the biosensor, and the surface was then rinsed with PBS, to establish the needed detection time.

Architecture of the biosensor -

The architecture of the biosensor is briefly described here. Some details, mainly concerning optimization protocol, are described more precisely in the "Results and Discussion" section. After cleaning the sensor with ultrapure water, a first CV measurement was made on bare electrode that served as a control for the subsequent functionalization steps. A polyelectrolyte multilayer (PEM), constituted by the alternate deposition of PSS and PAH, was then formed by spin coating (45 s, 100 rpm, 200 s -2 ) on the gold working electrode, by spreading alternatively 10 µL of each polyelectrolyte (5 g/L in ultrapure water). Between each layer, the sensor was rinsed with ultrapure water, and a CV measurement was made to follow the PEM construction. Then, when NPs were implemented into the biosensor, 10 µL of a SiO2-NPs + , at a concentration of 10 g/L, were attached through electrostatic interactions to the PSS ending layer of the PEM, by spin coating (20 s, 200 rpm, 200 s -2 ) before being rinsed with ultrapure water. Finally, a drop of 7 µL of Ab, concentrated at 0.05 mg/mL in PBS, was deposited either on the PEM or on the NPs and left to incubate for 2 h at room temperature. Non-occupied sites were finally blocked with a 1 mg/mL solution of BSA for 30 min.

Biosensor architecture -

We here compare types of biosensor for E. coli bacteria detection: one is implemented with NPs grafted with Abs whereas the other one lacks NPs and only exhibits Abs as a bioreceptor. This comparison allows assessing the role of NPs in biosensor performances. Both systems exhibit a similar initial architecture (Figure 1a, Figure S1): the transducer, a gold electrode, is first functionalized with a polyelectrolytes multilayer (PEM) in order to electrostatically immobilize the second layer of either NPs grafted with Abs or Abs alone. The PEM has already been studied in details in [START_REF] Mathelié-Guinlet | Etude de l'intéraction nanoparticulesbactéries : application à l'élaboration d'un biocapteur[END_REF]. It is made of a spin-coated (45 s, 100 rpm) assisted Layer by Layer assembly of polystyrene sulfonate polyanions (PSS-) and poly(allylamine hydrochloride) polycations (PAH + ): 6 ((PSS-PAH)3 or PEM6) or 7 ((PSS-PAH)3-PSS or PEM7) layers are alternatively deposited from a 5 g/L polyelectrolyte solution by electrostatic interactions with the ending layer adjusted according to the charge of the targeted system: PEM7, ending with a PSS layer, allows immobilizing positively charged NPs whereas the PEM6, ending with a PAH layer, allows the immobilization of negatively charged bacteria or of the Abs through bonds with its carboxylate functions. j'ai mis des points à la place des virgules Following this first functionalization, a self-assembled layer of Abs is formed either directly on the PEM or on a monolayer of NPs previously deposited on the PEM. In the former case, Abs are incubated on the transducer for 2 h at a concentration of 0.05 mg/mL, unbounded sites being finally blocked with Bovin Serum Albumin (BSA). In the latter case, we used harmless positively charged silica NPs (coated with aminopropyltriethoxysilane (APS)) of 100 nm in diameter [START_REF] Mathelié-Guinlet | Probing the threshold of membrane damage and cytotoxicity effects induced by silica nanoparticles in Escherichia coli bacteria[END_REF]. They were integrated through a spin-coated (20 s, 200 rpm) deposit from a 10 g/L solution. Abs were then deposited on NPs with the same protocol exposed above. Then, the detection of bacteria is ensured through CV experiments (Figure 1b, Figure S1). From these voltammograms, various parameters are extracted and analyzed for quantification: the peak intensities for oxidation (Ia) and reduction (Ic) and the difference of potential between the two peaks (ΔE) which are indicative of respectively the chemical and electrochemical stabilities of the redox couple in the medium. Before any measurement in the presence of bacteria, the effect of these diverse layers added to the transducer were studied to unravel which modifications would be attributed to actual detection thereafter. Briefly, compared with voltammograms obtained on the bare electrode, those following the deposition of a PEM showed a slight decrease of both oxidation and reduction peak intensities (Figure 1b), as reported before [START_REF] Mathelié-Guinlet | Etude de l'intéraction nanoparticulesbactéries : application à l'élaboration d'un biocapteur[END_REF]. This suggests a decrease in the kinetics of charge transfer at the electrode surface, due to the presence of a "screening" thin film of polyelectrolytes. Besides, adding NPs on the PEM did not induce significant differences in the subsequent voltammograms (Figure 1b), certainly due to the surface positive charges that (i) do not add supplementary resistance and that (ii) might favor charge transfer, the redox couple being negatively charged. Despite a decrease in the CV signal subsequent to these functionalizations, its shape remains unchanged and peak intensities are still largely enough for the ongoing bacteria detection.

E. coli bacteria detection -

First, to test the robustness of our biosensors, we first had to eliminate any undesired non-specific detection. We therefore performed some controls by directly depositing E. coli bacteria on a PEM6, i.e. in the absence of either NPs or Abs and recording the CV. The subsequent CV showed no significant variation as compared to CV obtained on the PEM6 (see Figure S2). This shows that no detection took place through unspecific electrostatic interactions between the positively charged PAH last layer and the negatively charged bacteria. Consequently, thereafter, only the biorecognition events will lead to a change in the CV characteristics. Then, we compared the biosensor performance depending on the presence or the absence of NPs, E. coli being incubated at different concentrations on the functionalized transducer for 15 min. In the absence of NPs, when the Abs were implemented into the biosensor architecture, on a PEM6, the CV showed simultaneously a decrease in oxidation-reduction peak intensities and of their ratio (Figure 2b-c blue stars) and an increase in the difference of potential (Figure 2d blue stars). The same trend was observed after incubation with bacteria at least at low concentration, both parameters stabilizing at concentrations higher than 4.10 7 CFU/mL. This suggests first that, logically, the Ab-bacteria antigen (Ag) specificity guarantees the recognition events and a minimal detection of E. coli bacteria. Second, the saturation of the signal at concentrations above 4.10 7 CFU/mL limits the quantification range of the biosensor, making it hardly possible. On the opposite, the combination of Abs deposited on a layer of silica-NPs, previously attached on a PEM7, led to a progressive modification of the CVs following the same trend on the whole range of bacteria concentrations: (i) the oxidation-reduction peak intensities and their ratio decreased continuously with the addition of increasingly concentrated bacteria suspensions (Figure 2a-c) while (ii) the difference of potential continuously increased (Figure 2d) (red squares). Such phenomena convey a slow-down in the kinetics of the charge transfer, due to an increased adsorbed mass at the electrode surface. These modifications suggest that E. coli bacteria are actually detected all over the tested range of concentrations [10 6 ; 10 9 CFU/mL]. Consequently, coupling Ab with NPs did induce a significant increase of efficiency in detecting E. coli bacteria, in a non-saturated and continuous way, as compared to the self-assembled system functionalized by Abs alone. This study was also performed during shorter (5 min.) and longer (30 min.) incubation times in the presence of E. coli and identical linear trends were obtained (Figure 3). From such linear variations, it was thus possible to determine the sensitivity (slope of the linear regression) and the LOD (abscissa at the origin, in a first approximation) of the designed biosensors, from the intensity variations obtained between detected bacteria and Ab (Table 1 and Figure 3). Considering the initial and non-optimized Ab concentration of 0.05 mg/mL, the silica-NPs assisted biosensor achieves a LOD of 8.10 5 CFU/mL within 15 min. With exactly the same experimental conditions, a decrease of the time detection to only 5 min. allowed anyway to accurately detect E. coli with a LOD of 1.10 6 CFU/mL (Table 1 and Figure 3). Though these LODs seem relatively high, it has to be underlined that (i) in various diseases, such as urinary infections, the significant threshold to qualify an infection is above 10 5 CFU/mL [25] pourquoi la ref est-elle passé en gras????and (ii) no usual detection technique is able to detect such a quantity in less than 15 min. Nevertheless, an optimization of this parameter was undertaken to meet the expected higher sensitivity. By increasing the time of detection to only 30 min., the Abs silica-NPs based biosensor was able to detect E. coli concentrations as low as 2.10 3 CFU/mL but, surprisingly with a lower sensitivity than the previous ones (Table 1 and Figure 3).

In the literature, lower LODs, of the order of tens to hundreds of CFU/mL, were obtained with electrochemical sensors incorporating magnetic NPs. They required however at least 1 h to reach this detection level [START_REF] Setterington | Electrochemical Biosensor for Rapid and Sensitive Detection of Magnetically Extracted Bacterial Pathogens[END_REF][START_REF] Geng | A DNA sequence-specific electrochemical biosensor based on alginic acid-coated cobalt magnetic beads for the detection of E. coli[END_REF]. At the same time, also using magnetic NPs, Varshney and Li reached a higher LOD of 10 5 CFU/mL within 35 min. [START_REF] Varshney | Interdigitated array microelectrode based impedance biosensor coupled with magnetic nanoparticle-antibody conjugates for detection of Escherichia coli O157:H7 in food samples[END_REF]. Based on these reports, it seems that there is a real need to systematically specify the biosensor efficiency not only in terms of sensitivity or LOD, but also in terms of required incubation time, all these parameters depending on the objective of the end-user who will pay more attention to the detection time while another one will expect a high sensitivity whatever the duration. In that sense, the NPs-based biosensors presented here offer nice perspectives. Finally, a last parameter is important, namely the sensor specificity. To test this parameter, same experiments were performed with another strain of Gram negative bacteria, Pseudomonas fluorescens. This strain did not trigger any significant variations of the CV shapes, whatever their concentration (see Figure S4). Indeed, the oxidation-reduction peak intensities and the difference of potentials did not decrease more than 5%, even for the highest bacteria concentration, which is well below the modifications reported for E. coli. These observations showed that Pseudomonas fluorescens were not detected by this biosensor, in agreement with the specificity of the Abs towards some Ag O units of E. coli membrane lipopolysaccharides, which differ in Pseudomonas species. This confirms a first step in this biosensor specificity.

In summary, this study illustrates a proof of concept of a new type of NPs assisted biosensor. It first shows that a rather simple and time-saving functionalization might be applied to an electrochemical transducer for the detection of E. coli. Most of all, it emphasizes the undeniable contribution of biocompatible silica NPs to the biosensor sensitivity. Indeed, in the absence of NPs, bacteria are detected to a small extent, however no quantification is possible due to non-continuous and random variations in the CV characteristic parameters. On the contrary, when NPs are implemented into the biosensor, bacteria are not only detected in the tested ranges (from low to high amounts) but always in a continuous way. This increase in sensitivity is likely linked to the much higher specific area of the NPs that can be grafted with Abs, as compared to a plan substrate. It might also be associated with a better orientation of the bioreceptors on such APS modified NPs. The variations of indicators, such as Ic, Ia, Ic/Ia and ΔE, qualitatively and straightforwardly point to the actual detection of E. coli. In addition, based on the linear variations of intensities, the quantification of detected E. coli is fulfilled. Extensive optimizations are still in progress to achieve the Holy Grail of an ultra-sensitive and fast biosensor. For instance, the reported current performance can be highly enhanced by the control of the orientation (through EDC-NHS protocol) and concentration of Abs grafted on the NPs. Nevertheless, the initial optimizations presented in this work highlight the power of NPs integrated into biosensors as E. coli are specifically detected in a very short time (5 min.) and the detection is enhanced to 10 3 CFU/mL by simply adjusting the time of detection (30 min.) according to the expectations of potential industrial applications.

Supporting Information. Details about all results obtained for control, specificity and other experiments (Figures S1- S3).

Figure 1 .

 1 Figure 1. a) Schematic illustration of the architecture of the NPs-based biosensor designed for E. coli detection. On the gold electrode are successively deposited a polyelectrolyte multilayer (PEM), a layer of NPs (100 nm in diameter) welcoming the Abs specific to E. coli. In the absence of NPs, the Abs are directly deposited on the PEM. For bacteria detection, the biosensor is immersed in a PBS solution containing the Fe(CN)6 3-/4-redox couple for voltammetry measurements. b) Cyclic voltammetry experiments of Fe(CN)6 3-/4- performed after each functionalization step of the transducer.

Figure 2 .

 2 Figure 2. Influence of the presence of NPs in the electrochemical detection of E. coli. a) Representative cyclic voltammetry experiments based on the redox couple Fe(CN)6 3- /4-performed after incubation of increasing concentrations of bacteria during 15 min. on the transducer functionalized by NPs/Ab. Variation with the concentration of E. coli of (b) the reduction peak intensity Ic, (c) the ratio between oxidation and reduction peak intensities Ic/Ia and (c) the difference of potential ΔE. Results are normalized to the values obtained on the bare gold electrode and thus expressed in %. Errors bars account for three independent measurements.

Figure 3 .

 3 Figure 3. Variation of the reduction peak intensity obtained on the NPs-Abs assisted biosensors during the detection of various concentrations of E. coli bacteria and for three different times of detection (5, 15 and 30 min.). Error bars account for three independent measurements.

Table 1 .

 1 Performance of the biosensors integrating Abs grafted on NPs, depending on the detection time. Sensitivities and LOD are determined from the variations of Ic.

	Detection Time	5 min.	15 min.	30 min.
	Sensitivity (µA/CFU/mL)	18	20	12
	LOD (CFU/mL)	1.10 6	8.10 5	2.10 3
	Tested range (CFU/mL)	8.10 6 -5.10 8	8.10 4 -8.10 6

Authors to whom correspondence should be addressed. Christine Grauby-Heywang: christine.grauby-heywang@ubordeaux.fr (Univ. Bordeaux, LOMA, CNRS UMR 5798, 351 crs Libération, 33405 Talence, France) and Marie-Hélène Delville: marie-helene.delville@icmcb.cnrs.fr (Univ. Bordeaux, CNRS, ICMCB, UMR 5026, 87 avenue du Dr Albert Schweitzer, 33608 Pessac, France)

No competing financial interests have been declared. The research presented in this article was funded by the Direction Générale de l'Armement (DGA, Ministère de la Défense, France grant number: 2014017), the Région Aquitaine (France) and LOMA through the Ph.D. grant to M. Mathelié-Guinlet.

The authors thank Alexandre Vilquin for his technical help with Matlab, allowing an automatization of all the analyzed data, Zeineb Ben Abdallah and Asma Soussou for their discussions and helpful advices during electrochemical studies.