Siham Difi

Moulay Ismael Saadoune

Tahar Sougrati

Rachid Hakkou

Kristina Edstrom

Pierre-Emmanuel Lippens

Moulay Tahar Sougrati

Mechanisms and Performances

INTRODUCTION

Electrochemical rechargeable batteries are considered one of the leading technologies for electrical energy storage, especially for the development of sustainable and renewable energy sources, electric vehicles, and portable electronic devices. [START_REF] Ibrahim | Energy Storage Systems-Characteristics and Comparisons[END_REF] Liion batteries provide high energy density, 2 but lithium resources are unevenly distributed worldwide and might not meet the increasing demand for energy. 3 In addition, cost and safety are major drawbacks, particularly for large-scale applications such as stationary energy storage systems. [START_REF] Pan | Room-Temperature Stationary Sodium-Ion Batteries for Large-Scale Electric Energy Storage[END_REF] In this light, rechargeable Na-ion batteries can be considered a serious alternative because of the natural abundance and low cost of sodium, but also for the safety of the electrochemical systems that use, for example, aqueous electrolytes. [START_REF] Kim | Aqueous Rechargeable Li and Na Ion Batteries[END_REF] The main disadvantage of these batteries is the lower energy density compared to Li-ion batteries due to lower operation voltage and higher atomic mass of Na compared to Li. However, Naion batteries could be the best option for applications when cycle life and low cost are more essential factors than energy density.

The electrochemical properties of Na 2 FeTi(PO 4 ) 3 obtained by solid state reaction were first reported by Tillement et al. [START_REF] Guler | A Rapid Synthesis of Sodium Titanium 1082 Phosphate, NaTi 2 (PO 4 ) 3 by Using Microwave Energy[END_REF] The potential curve of the first discharge obtained by these authors in the galvanostatic mode at C/50 (1 Na/f.u. in 50 h) shows a two-phase plateau at 2.3 V versus Na + /Na 0 until 0.5 inserted Na/f.u. followed by a continuous decrease until 1.5 V versus Na + /Na 0 for about 2 inserted Na/f.u. The incremental capacity curve shows peaks at 1 and 1.5 inserted Na/f.u. that were attributed to local orderings. A material with the same composition was also synthesized by solid state reaction, but with a somewhat different procedure, by Patoux et al. [START_REF] Tillement | Electrochemical Studies of Mixed Valence NASICON[END_REF] The structure was accurately determined by neutron diffraction. The same rhombohedral space group (R3̅ c) as that of Na 3 Ti 2 (PO 4 ) 3 was obtained with slightly different lattice parameters. The potential curve obtained in the galvanostatic regime at C/10 for the first cycle shows neither a well-defined plateau as in the case of NaTi 2 (PO 4 ) 3 nor a two-step mechanism as observed by Tillement et al. [START_REF] Guler | A Rapid Synthesis of Sodium Titanium 1082 Phosphate, NaTi 2 (PO 4 ) 3 by Using Microwave Energy[END_REF] but a single-phase reaction. This result was also confirmed by potential intermittent titration technique (PITT) and by in situ X-ray diffraction (XRD). Recently, Aragon et al. 28 used sol-gel method and citric acid as carbon source to obtain Na 1+x Fe x Ti 2-x (PO 4 ) 3 /C composites with x = 0, 0.2, 0.4, and 0.6. For all the compositions, the structure is rhombohedral and the composites were formed by aggregation of carbon-coated nanoparticles. The potential curves measured in the galvanostatic mode at C/10 show a two-step reaction mechanism which differs from those described by Tillement et al. [START_REF] Guler | A Rapid Synthesis of Sodium Titanium 1082 Phosphate, NaTi 2 (PO 4 ) 3 by Using Microwave Energy[END_REF] and by Patoux et al. [START_REF] Tillement | Electrochemical Studies of Mixed Valence NASICON[END_REF] Along the discharge, there is a continuous decrease of the potential followed by a well-defined plateau at about 2.1 V versus Na + /Na 0 . The first domain is limited to about 0.1, 0.25, and 0.3 inserted Na/f.u. for x = 0.2, 0.4, 0.6, respectively, although the PITT curve for x = 0.6 shows a more extended domain until 0.6 inserted Na/f.u. In the latter case, the PITT curve exhibits a series of successive small plateaus characteristic of local Na orderings over the M2 interstitial sites. This first domain was attributed to change in Fe oxidation state from Fe 3+ for the pristine material to Fe 2+ based on ex situ [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] Fe Mossbauer experiments carried out for x = 0.4. For 1 inserted Na/f.u., the Fe 2+ Mossbauer component represents only 80% of the total spectrum area, which indicates that although most of the Fe 3+ were reduced in the first half of the discharge, there is no clear evidence that this change took place only in the first domain (<0.4 inserted Na/f.u.). The authors pointed out that substitution of Fe for Ti improves electrochemical performances compared to Na 3 Ti 2 (PO 4 ) 3 . The specific capacities of the first discharge were found to be 120, 130, 119, and 117 mAh g -1 , and the Coulombic efficiency were 79%, 91%, 87%, and 85% for x = 0, 0.2, 0.4, 0.6, respectively.

The observed decrease of the capacity with increasing x ≥ 0.2 was attributed to the increase of Na amount in the pristine material. Improved capacity retention and power rate were also observed with a Coulombic efficiency in the range of 85-92% and a decrease of the capacity lower than 10% from C/10 to C/ 2 that have been related to low internal impedance. Thus, the comparison of the results obtained by these authors suggests that different insertion mechanisms operate in Fe-Ti-based NASICON electrode material for Na-ion batteries that not only depend on the composition but also on other aspects of the pristine material like crystallinity, particle size, carbon coating, etc. This also depends on the experimental conditions used for electrochemical tests as observed, for example, from the differences between galvanostatic and PITT results that could be due to kinetic effects. 28 The Journal of Physical Chemistry C were recorded in the 100-2000 cm -1 range, at room temperature using a LabRAM ARAMIS IR 2 spectrometer equipped with a 632 nm laser diode, a 1200 gr mm -1 grating with X50 long work distance objective. Each spectrum was recorded with an acquisition time of 30 min.

Magnetic Measurements.

The magnetic susceptibility was measured with a Superconducting Quantum Interference Design (SQUID) magnetometer MPMS XL7.

The samples were first cooled to 2 K under a zero magnetic field (ZFC). Then, a magnetic field of 100 mT was applied, and the susceptibility was recorded from 2 to 300 K. Finally, the susceptibility was also measured by cooling the samples to 2 K with the same field (FC).

2.5. Mossbauer Spectroscopy. The 57 Fe Mossbauer spectra were recorded in transmission geometry and a constant acceleration mode at room temperature with [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] Co (Rh) as the γ source. The velocity scale was calibrated using the magnetic sixline spectrum of a high-purity iron foil absorber as a standard.

The values of isomer shift (δ), quadrupole splitting (Δ), full line width at half-maximum (Γ), and relative area (RA) were determined by fitting Loretzian lines to experimental data with a nonlinear least-square method. The values of the isomer shift (δ) are given with respect to α-Fe, and the goodness of the fits was controlled by a classical χ 2 test. A specific electrochemical cell was used for in situ [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] Fe Mossbauer spectroscopy in the transmission mode. [START_REF] Leriche | An Electrochemical Cell for Operando Study of Lithium Batteries Using Synchrotron Radiation[END_REF] Two methods were considered to study changes in the Fe oxidation state during charge-discharge cycles. In the first one, the operando Mossbauer data were collected along the first cycle performed in the galvanostatic mode at C/25 during time steps of 5 h. Thus, each spectrum corresponded to the insertion (discharge) or extraction (charge) of 0.2 Na/f.u. However, the evaluation of the hyperfine parameters was made difficult by the rather poor signal-to-noise ratio due to the small amount of [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] Fe (natural abundance: 2.12%) in the electrode material, and only the relative contributions of the two oxidation states to the whole spectra were evaluated. The other method was based on GITT mode in order to record in situ Mossbauer spectra at equilibrium potentials. This allowed for the reduction of the influence of the kinetic effects and an increase in the recording time without increasing the number of inserted Na as in galvanostatic mode. In that case, one full discharge-charge cycle and the second discharge were conducted at C/10, interrupted every 4 h for the first cycle and every 1 h for the second discharge by relaxation periods of 12 h to collect the in situ Mossbauer data.

2.7. Theoretical Method. The theoretical structural optimizations of NaTi 2 (PO 4 ) 3 and Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 were performed with the pseudopotential method as implemented in CASTEP code [START_REF] Segall | First-Principles Simulation: Ideas, Illustrations and the CASTEP Code[END_REF] based on the density functional theory (DFT) [START_REF] Kohn | Self-Consistent Equations Including Exchange and Correlation Effects[END_REF][START_REF] Hohenberg | Inhomogeneous Electron Gas[END_REF] and the generalized-gradient approximation (GGA) using an exchange-correlation potential by Perdew, Burke, and Ernzerhof for solids (PBESOL). [START_REF] Perdew | Restoring the Density-Gradient Expansion for Exchange in Solids and Surfaces[END_REF] Ultrasoft pseudopotentials [START_REF] Vanderbilt | Soft Self-Consistent Pseudopotentials in a Generalized Eigenvalue Formalism[END_REF] and the reciprocal space representation were used. For Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 , a supercell of 37 atoms was used with the P1 space group in order to fully relax all the atomic positions. All the different configurations of Na atoms within the interstitial sites M1 and M2 were considered. The plane-wave kinetic-energy cutoff was set to 400 eV, and the kspace sampling was performed on a Monkhorst-Pack grid with 14 k points in the irreducible wedge of the Brillouin zone. Both the lattice parameters and the atom positions were relaxed until the mechanical equilibrium was achieved through conjugategradient minimization of the total energy to a tolerance of 10 -5 eV/atom, the forces to a tolerance of 0.03 eV Å -1 , the stress to a tolerance of 5 × 10 -2 GPa, and the atomic positions to a tolerance of 10 -3 Å. For the two most stable configurations of Na atoms, the convergence of the structural parameters was checked by increasing the energy cutoff (800 eV) and the number of k points (32 k), while the convergence parameters were changed to 10 -6 eV/atom, 5 × 10 -3 eV Å -1 , 10 -2 GPa, and 2 × 10 -4 Å. A density-mixing scheme with a conjugategradient Pulay solver was used for the energy minimization and a BFGS algorithm for the internal-coordinate optimization. The [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] Fe Mossbauer parameters of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 were evaluated with the linearized augmented plane-wave method (LAPW) as implemented in the WIEN2k code [START_REF] Blaha | WIEN2K An Augmented Plane Wave Plus Local Orbitals Program for Calculating Crystal Properties[END_REF] by considering the structural parameters of the most stable phase obtained from pseudopotential calculations. With this method, the unit cell is partitioned into atomic spheres centered at the atomic positions and an interstitial region. The muffin-tin radii of the atomic spheres were R mt (Na) = 2.2 au, R mt (Fe) = 2.0 au, R mt (Ti) = 2.0 a.u, R mt (P) = 1.4 au, and R mt (O) = 1.4 au. To improve the energy linearization, the basis set was extended with local orbitals. In the interstitial region, the wave functions were expanded in plane waves with wavenumbers k such as min(R mt )•max(k) = 8, and the charge density was expanded in a Fourier series with G max = 16 Ry -1/2 .

RESULTS AND DISCUSSION

3.1. Theoretical Structure of the Pristine Material. It has been reported previously that Na 1+x Fe x Ti 2-x (PO 4 ) 3 compounds with 0 ≤ x ≤ 1 are isostructural to NaTi 2 (PO 4 ) 3 . [START_REF] Guler | A Rapid Synthesis of Sodium Titanium 1082 Phosphate, NaTi 2 (PO 4 ) 3 by Using Microwave Energy[END_REF][START_REF] Tillement | Electrochemical Studies of Mixed Valence NASICON[END_REF][28] The structure of NaTi 2 (PO 4 ) 3 can be described in the rhombohedral system (space group R3̅ c, n°167). The rhombohedral primitive cell contains 36 atoms (Z = 2). In the most commonly used hexagonal description, Ti and P are in 12c and 18e Wickoff positions, respectively, and there are two different crystallographic sites for the oxygen atoms, O(1) and O(2), both in 36f Wyckoff positions. The structure consists of a three-dimensional framework of corner sharing PO 4 tetrahedra and TiO 6 octahedra, which are both slightly distorted. There are two possible types of interstitial sites that can be occupied by Na: M1 in 6b and M2 in 18e Wickoff positions, respectively. The M1 site is surrounded by 6 O firstnearest neighbors forming a distorted octahedron. These O atoms belong to the triangular faces of two adjacent TiO 6 octahedra along the hexagonal c axis. The M2 sites are located between the PO 4 tetrahedra and have 8 O first-nearest 377 neighbors that form a strongly distorted polyhedron. The 378 published experimental lattice parameters are found in the 379 ranges a hex = 8.48 ± 0.01 Å and c hex = 21.80 ± 0.02 Å, 380 depending on the synthesis method and the experimental 381 conditions of measurements. [START_REF] Delmas | Nasicon-Type Phase as Intercalation Electrode: NaTi 2 (PO 4 ) 3[END_REF][START_REF] Hagman | The crystal Structure of NaMe 2 IV (PO 4 ) 3 ; Me IV = Ge, Ti[END_REF][START_REF] Rodrigo | Thermal Expansion of NaTi 2 (PO 4 ) 3 Studied by Rietveld Method from X-ray Diffraction Data[END_REF]

d(Na-O) (Å) 2.451 × 6 2.457 × 6 d(Ti-O) (Å) 1.906 × 3 + 1.967 × 3 1.902 × 3 + 1.958 × 3 d(P-O) (Å) 1.523 × 2 + 1.524 × 2 1.520 × 2 + 1.528 × 2
obtained with the PBESOL functional. Moreover, GGA and LDA calculations were also performed with the functionals of Perdew, Burke, and Ernzerhof (PBE) [START_REF] Perdew | Generalized Gradient Approximation Made Simple[END_REF] and Perdew and Zunger (PZ), [START_REF] Perdew | Self-Interaction Correction to Density-Functional Approximations for Many-Electron Systems[END_REF] respectively. In the two cases, the most stable Na configuration is the same as that obtained with the PBESOL functional. However, the rhombohedral cell volumes were found to be 465.6 Å 3 (PBE) and 438.8 Å 3 (PZ), which overestimates and underestimates the experimental volume by about 3%, respectively. Finally, the effect of Ti 3d electron correlation was checked by a GGA (PBESOL) + U spin polarized calculation, with U = 4 eV. The calculated rhombohedral cell volume (459.1 Å 3 ) is slightly higher than that obtained with U = 0 eV, which does not improve the agreement with the experimental value. These results show the high accuracy of the DFT-PBESOL method for the geometry optimization of NaTi 2 (PO 4 ) 3 , which justifies its use for the structural analysis of the substitution of Fe for Ti in this compound.

A supercell of 37 atoms with the composition Na 3 FeTi 3 (PO 4 ) 6 was considered for the structural optimization of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 . The lattice parameters and the atomic positions of the rhombohedral cell of NaTi 2 (PO 4 ) 3 were used to initialize the calculations, but one Ti atom was replaced by Fe in the supercell and relaxation of all the atoms was allowed by considering the P1 space group. There are 56 possible distributions of the 3 Na atoms over the 2 M1 and 6 M2 sites (3 sites close to Fe: M2/Fe and 3 sites close to Ti: M2/Ti), among them some are equivalent due to the symmetry of the starting supercells. All the calculations provided high-spin Fe 3+ , whatever the Na distribution. The most stable configuration was obtained with 3 Na atoms in the 2 M1 sites and in 1 M2 site close to Fe (M2/Fe). However, the total energy is only 0.065 eV/f.u. lower than that obtained with Na in the 2 M1 sites and in the 1 M2/Ti site. The values of the total energy for the different configurations of Na in M1 + 2 M2 sites are about 0.2-0.4 eV/f.u. higher than those with Na in 2 M1+M2/Fe, while the values for the configurations of Na in 3 M2 sites are about 0.3-0.6 eV/f.u. higher than those with Na in 2 M1+M2/ Fe. The supercell volumes are about 462 Å 3 (461 Å 3 with improved computational parameters) for Na in 2 M1+M2, 466-470 Å 3 for Na in M1 + 2 M2, and 473-480 Å 3 for Na in 3 M2. Thus, the most stable Na configuration was obtained with fully occupied M1 sites and Na in a M2 site close to the FeO 6 octahedra. For a comparison with NaTi 2 (PO 4 ) 3 and the experimental data, the theoretical structural parameters of an equivalent R3̅ c crystal were evaluated from the P1 optimized supercell with improved computational parameters by averag-t2 ing the lattice parameters and atomic coordinates (Table 2).

The theoretical rhombohedral cell volume of the most stable configuration obtained for Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 is 5 Å 3 higher than the theoretical volume of NaTi 2 (PO 4 ) 3 . This increase of 1% is due to the increase of a hex (and b hex ) from 8.521 Å for NaTi 2 (PO 4 ) 3 to 8.582 Å since c hex decreases from 21.748 Å for NaTi 2 (PO 4 ) 3 to 21.685 Å. These changes are rather small but reflect a significant trend in the lattice parameters that was experimentally observed for other Fe contents. [START_REF] Tillement | Electrochemical Studies of Mixed Valence NASICON[END_REF]28 The decrease of c hex is mainly due to the decrease of the ionic charge from Ti 4+ to Fe 3+ , which lowers the repulsive interactions between the transition metal cations that are located along the c hex axis. As a result, the Fe-Ti and Fe-N a ( M 1 ) i n t e r a t o m i c d i s t a n c e s a l o n g c h e x i n Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 are smaller than those of Ti-Ti and f1 Ti-Na(M1) in NaTi 2 (PO 4 ) 3 , respectively (Figure 1a). Due to (Table 3). As an illustrative example, the 3). These local changes induced In addition, the relative occupations of the 12c sites with 25 at%

Fe and 75 at% Ti, the M1 sites with 100 at% Na and the M2 sites with 17 at% Na were considered for the initialization. An overall correct agreement is obtained between the experimental t4 (Table 4) and theoretical (Table 2) values of the lattice (Figure 2b). There are no additional peaks due to the sucrose pyrolysis that could be assigned to crystalline impurities. The structural parameters obtained from Rietveld refinement are close to those of the pure crystalline phase and are not reported here.

The [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] Fe Mossbauer spectrum of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 measured at room temperature was fitted with a symmetrical f3 doublet (Figure 3a). The small line width indicates that Fe local environments are similar within the overall crystalline phase t5 (Table 5). The values of the isomer shift, 0.437(1) mm s -1 , and the quadrupole splitting, 0.326(2) mm s -1 , are close to previously published values for iron-based NASICON compounds and can be assigned to high spin Fe 3+ ions in slightly distorted FeO 6 octahedra. [START_REF] Lajmi | Synthesis and Structural Study of a New Iron Phosphate K 3 Fe(PO 4 ) 2[END_REF][START_REF] Jirak | Magnetic and Mossbauer Resonance Investigations of the Weak Ferrimagnet Iron Molybdate (Fe 2 (MoO 4 ) 3 )[END_REF][START_REF] Beltran-Porter | Etude Magnetique et par Reśonance Mossbauer de l'Orthophosphate Na 3 Fe 2 (PO 4 ) 3 α et d'une Phase Vitreuse Deŕiveé[END_REF] Higher values of the quadrupole splitting (∼0.8 mm s -1 ) were recently reported for Na 1+x Fe x Ti 2-x (PO 4 ) 3 with x ≤ 0.6. 28 To confirm the present value, LAPW calculations of the electric field gradients at the Fe nucleus were performed for the theoretically optimized P1

supercell obtained for the most stable Na configuration with the pseudopotential method. The theoretical approach is described in section 2 and was previously used for different Mossbauer isotopes. [START_REF] Blaha | First-Principles Calculation of the Electric Field Gradient of Li 3 N[END_REF][START_REF] Lippens | Interpretation of the 119 Sn Mossbauer parameters[END_REF][START_REF] Lippens | Calculation of 121 Sb and 125 Te Mossbauer Spectra[END_REF][START_REF] Weibel | Local Atomic and Electronic Structure in Nanocrystalline Sn-Doped Anatase TiO 2[END_REF][START_REF] Yahla | First-Principles Calculations of the Electronic Structure and Mossbauer Parameters of Sb-Doped SnO 2[END_REF] The quadrupole splitting is given by

η Δ = + ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ eQV 1 2 1 3 zz 2 1/2 (1)
where e is the electron charge, Q is the [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] Fe nuclear quadrupole moment of the first excited state (I = 3/2), V ZZ is the main a The polyhedron distance distortion Δ poly is defined by 552 component of the diagonalized EFG tensor, and η is the 553 asymmetry parameter defined by

Δ = (1/n)∑ i = 1 n [(d i -d av )/d av ] 1/2 .
η = - V V V XX YY
ZZ 554 (2) 555 where V XX and V YY are the other two components of the 556 diagonalized EFG tensor with V XX < V YY < V ZZ . We obtained 557 V ZZ = 0.9 V m -2 and η = 0.63, which gives Δ = 0.16 mm s -1 by considering Q = 0.16 b. [START_REF] Dufek | Determination of the Nuclear Quadrupole Moment of 57 Fe[END_REF] To take into account the Fe 3d electron correlation, we also performed a GGA+U calculation which gave V ZZ = 1.9 V m -2 and η = 0.62, leading to Δ = 0.32 mm s -1 in better agreement with the present experimental value. In the two cases, the rather small value of Δ originates from the weak asymmetry of the electronic density around Fe. This is due to the weak distortion of the FeO 6 octahedral environment (Figure 1b) and the high spin Fe 3+ electronic configuration. This result confirms the present interpretation of the Mossbauer spectrum. The recently published higher values of Δ could be due to the small size of the particles obtained by the sol-gel method, which provides higher surface-to-bulk ratio and more asymmetrical Fe local environments, or to structural disorder. [START_REF] Beltran-Porter | Etude Magnetique et par Reśonance Mossbauer de l'Orthophosphate Na 3 Fe 2 (PO 4 ) 3 α et d'une Phase Vitreuse Deŕiveé[END_REF] The 57 Fe Mossbauer spectrum of the composite Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C shows an additional small peak at about 2.4 mm s -1 and was fitted with two doublets (Figure 3b).

The Mossbauer parameters of the main doublet (91% relative area), δ = 0.420(7) mm s -1 and Δ = 0.275(5) mm s -1 , are close to those of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 , while the isomer shift of the small doublet (9% relative area): δ = 1.178(7) mm s -1 is typical of high spin Fe 2+ (Table 5). The high value of the quadrupole splitting obtained for the latter doublet, Δ = 2.368(2) mm s -1 , indicates a strong anisotropy of the electron density around Fe. This is mainly due to the partial occupation of the Fe 3d spin down states but could also be related to the more distorted Fe local environment. This additional component arose from sucrose pyrolysis and could be related to intrinsic defects or external impurities such as olivine Na x FePO 4 , the Mossbauer parameters of which are close to the present values for Fe 2+ . [START_REF] Boucher | Elucidation of the Na 2/3 FePO 4 and Li 2/3 FePO 4 Intermediate Superstructure Revealing a Pseudouniform Ordering in 2D[END_REF] The absence of additional XRD peaks is in favor of intrinsic defects or impurities with small size and/or amorphous.

Whatever its exact nature, this defect will be called "Fe 2+ impurity" later in this paper, and its influence upon the reaction mechanism and electrochemical performances will be analyzed in the section 4.

The inverse magnetic susceptibility of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 progressively increases with temperature from the origin and f4 follows a Curie-Weiss law above about 30 K (Figure 4). This curve differs from that of Na 3 Fe 2 (PO 4 ) 3 , which shows a typical antiferromagnetic order at low temperature below the Neél temperature of 47 K. [START_REF] Beltran-Porter | Etude Magnetique et par Reśonance Mossbauer de l'Orthophosphate Na 3 Fe 2 (PO 4 ) 3 α et d'une Phase Vitreuse Deŕiveé[END_REF][START_REF] Greaves | The Nature of Magnetic Ordering at 47 K in Na 3 Fe 2 (PO 4 ) 3[END_REF] The present Curie temperature: θ NFTP = -8 K is close to 0 K and strongly differs from that of Na 3 Fe 2 (PO 4 ) 3 at -85 K. [START_REF] Beltran-Porter | Etude Magnetique et par Reśonance Mossbauer de l'Orthophosphate Na 3 Fe 2 (PO 4 ) 3 α et d'une Phase Vitreuse Deŕiveé[END_REF] This clearly indicates that the antiferromagnetic order is negligible in our samples and [START_REF] Beltran-Porter | Etude Magnetique et par Reśonance Mossbauer de l'Orthophosphate Na 3 Fe 2 (PO 4 ) 3 α et d'une Phase Vitreuse Deŕiveé[END_REF][START_REF] Greaves | The Nature of Magnetic Ordering at 47 K in Na 3 Fe 2 (PO 4 ) 3[END_REF] This value is typical of high spin Fe 3+ in the octahedral FeO 6 environment and is lower than the spin-only value of the Fe 3+ free ion (5.9 μ B ). The variations of the inverse magnetic susceptibility of the Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C composite are similar to those of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 except for a small decrease of the Curie constant: C = 3.57 emu K mol -1 , leading to a magnetic moment μ eff = 5.34μ B per Fe atom. If we consider, for simplicity, that the magnetization of the composite originates only from those of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 and Fe 2+ impurity, the Curie-Weiss law in the paramagnetic regime can be written:

χ θ θ θ = - = - + - + + C T C T C T 0.91 0.09 NFTP NFTP Fe Fe 2 2 (3)
where C Fe 2+ and θ Fe 2+ denote the Curie constant and the Curie temperature of the Fe 2+ impurity, respectively. Since the Curie t e m p e r a t u r e s o f N a 1 . 5 F e 0 . 5 T i 1 . 5 ( P O 4 ) 3 a n d Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C are very close, we have considered, for simplicity, that all the Curie temperatures are identical in eq 3 and obtained C Fe 2+ = 2.56 emu K mol -1 . The resulting magnetic moment, μ eff = 4.5μ B per Fe atom, is lower than the spin-only magnetic moment of Fe 2+ (4.9 μ B ). This result confirms the existence of high spin Fe 2+ in the composite as observed from Mossbauer spectroscopy.

The SEM images show that Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 particles have an average size of several micrometers and rather smooth f5 surfaces (Figure 5a), while Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C is composed by micrometer-sized secondary particles formed by the

f7

The Raman spectrum of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 (Figure 7) is 650 similar to those of LiTi 2 (PO 4 ) 3 and NaTi 2 (PO 4 ) 3 with rather 651 close frequencies. [START_REF] Yue | Hydrothermal Synthesis of MTi 2 (PO 4 ) 3 (M = 1080 Li, Na, K)[END_REF][START_REF] Cretin | Study of Li 1+x A1 x Ti 2-x (PO 4 ) 3 for Li + Potentiometric Sensors[END_REF][START_REF] Burba | Vibrational Spectroscopic Study of Lithium Intercalation into LiTi 2 (PO 4 ) 3[END_REF][START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF][START_REF] Yue | Hydrothermal Synthesis of MTi 2 (PO 4 ) 3 (M = Li, Na, K)[END_REF][START_REF] Bamberger | Synthesis and Characterization of Sodium-Titanium Phosphates, Na 4 (TiO)(P0 4 ) 2 , Na(TiO)P0 4 , and NaTi 2 (P0) 3[END_REF] On the basis of previously published 652 analyses, all the structures of the present Raman spectrum can 653 be attributed to the NASICON material. The broad band 654 around 1000 cm -1 and the small peak at 1093 cm -1 can be 655 assigned to intramolecular stretching modes of PO 4 : the 656 symmetric modes (ν 1 ) can be assigned to the 962 and 985 cm -1 657 frequencies and the antisymmetric modes (ν 3 ) to the 1005 and 658 1093 cm -1 frequencies. The small structures in the 500-700 659 cm -1 range and the sharp peak at 436 cm -1 arise from 660 antisymmetric (ν 4 ) and symmetric (ν 2 ) bending modes of PO 4 , Table 5. [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] respectively. The structures below 400 cm -1 are due to external modes and reflect mixtures of different atomic motions. However, it is possible to attribute the main contributions that are due to translational and rotational modes of PO 4 at 336, 310, 225, 195, and 157 cm -1 and to translational modes of Ti 4+ /Fe 3+ at 272 and 225 cm -1 . It is worth noting that frequencies in the range of 300-400 cm -1 are sensitive to metal substitution such as the Fe substitution for Ti in the octahedral environment. [START_REF] Barj | Raman and Infrared Spectra of Some Chromium Nasicon-Type Materials: Short-Range Disorder Characterization[END_REF] Compared to NaTi 2 (PO 4 ) 3 , some small shifts (<10 cm -1 ) were observed that can be explained by the difference between the atomic masses of Fe and Ti. In addition, the peaks are overall broadened as expected from the chemical disorder on the 12c crystallographic site that is responsible of bond length and bond angle distributions in both MO 6 (M = Fe, Ti) and PO 4 units as discussed above from DFT calculations (Table 3). Similar effects were observed previously for other cation-substituted NASICON compounds. [START_REF] Cretin | Study of Li 1+x A1 x Ti 2-x (PO 4 ) 3 for Li + Potentiometric Sensors[END_REF][START_REF] Junaid Bushiri | Raman and FTIR Studies of the Structural Aspects of NASICON-Type Crystals[END_REF] The spectrum of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C is similar except for a broad band at ∼1350 cm -1 and a narrower band at ∼1580 cm -1 , with notably higher intensities than the vibrational modes of the NASICON phase (Figure 7). There are also some small changes for frequencies higher than about 700 cm -1 , mainly due to overlapping with the main broad band at 1350 cm -1 .

The two bands at 1350 and 1580 cm -1 can be assigned to the D and G bands of residual carbon arising from pyrolysis, respectively. Following the approach of Sadezki et al., [START_REF] Sadezky | Raman Microspectroscopy of Soot and Related Carbonaceous Materials: Spectral Analysis and Structural Information[END_REF] these two bands were fitted to 5 peaks attributed to the vibrational modes of crystalline graphite (G), disordered graphite (D1, D2, D4), and amorphous carbon (D3) (Figure 7). The frequencies of the peaks D4, D1, D3, G, and D2 are 1231, 1346, 1464, 1575, and 1604 cm -1 , respectively, in line with the values reported for different types of disordered carbons with the same fitting procedure. The relative integrated intensities with respect to G are 0.63 (D4), 2.47 (D1), 0.84 (D3) and 0.43 (D2), respectively. There is no clear correlation between these values and structural information, [START_REF] Sadezky | Raman Microspectroscopy of Soot and Related Carbonaceous Materials: Spectral Analysis and Structural Information[END_REF] but the intensity ratio between D1 + D4 and G + D2 (∼2.2) is similar to the values obtained for different carbon blacks [START_REF] Jawhari | Raman Spectroscopic Characterization of Some Commercially Available Carbon Black Materials[END_REF] and carbon coating onto LiFePO 4 , [START_REF] Doeff | Effect of Surface Carbon Structure on the Electrochemical Performance of LiFePO 4[END_REF][START_REF] Julien | Characterization of the Carbon Coating onto LiFePO 4 Particles Used in Lithium Batteries[END_REF] revealing the existence of small crystalline domains. The rather large D1 line width (150 cm -1 ) suggests a lower degree of graphitization than highly graphitic carbon blacks [START_REF] Sadezky | Raman Microspectroscopy of Soot and Related Carbonaceous Materials: Spectral Analysis and Structural Information[END_REF][START_REF] Jawhari | Raman Spectroscopic Characterization of Some Commercially Available Carbon Black Materials[END_REF] but is comparable to that found for carbon coating on electrode materials. [START_REF] Doeff | Effect of Surface Carbon Structure on the Electrochemical Performance of LiFePO 4[END_REF][START_REF] Julien | Characterization of the Carbon Coating onto LiFePO 4 Particles Used in Lithium Batteries[END_REF] Thus, TEM and Raman spectros- more intense peak at 2.18 V (labeled B) that can be related to 724 the potential variations in R1 and R2, respectively (Figure 8b).

725

The first charge shows two structures A′ and B′ similar to A 726 and B, respectively, except for a broadening of the peak A′ 727 compared to A that is also observed in the following cycles. There is a voltage shift between B′ and B of about 20 mV, which indicates a small potential polarization of the electrode material due to low internal impedance and reflects the efficiency of the carbon coating.

The GITT curve of the first cycle was recorded in the 1.6-2.3 V range. At each insertion stage, a current of C/10 was applied for 1 h and followed by an 8 h relaxation period in f9 order to approach the equilibrium potential (Figure 9). This curve confirms the two-step mechanism with a change of the slope around 0.5 inserted Na/f.u. and also indicates that potential in R2 does not strictly form a plateau even at equilibrium. Such a two-step potential shape differs from the voltage profile of NaTi 2 (PO 4 ) 3 formed by a well-defined plateau due to a two-phase reaction between the pristine material NaTi 2 (PO 4 ) 3 (R3̅ c) and Na 3 Ti 2 (PO 4 ) 3 (P1̅ ) that reflects the Ti 4+ /Ti 3+ redox couple. 15,[START_REF] Delmas | Nasicon-Type Phase as Intercalation Electrode: NaTi 2 (PO 4 ) 3[END_REF] Previous studies on LiTi 2 (PO 4 ) 3 and Li 3 Fe 2 (PO 4 ) 3 NASICON-based electrode materials for Li-ion batteries show that Ti 4+ /Ti 3+15,18 and Fe 3+ /Fe 2+66 redox couples operate at about 2.5 and 2.8 V versus Li + /Li 0 , respectively. From the difference between the redox potentials of Li + /Li 0 and Na + /Na 0 (0.33 V), we can expect that Ti 4+ /Ti 3+ and Fe 3+ /Fe 2+ redox couples operate at 2.2 and 2.5 V versus Na + /Na 0 , respectively. For Ti 4+ /Ti 3+ , the former value is in line with that of the potential plateau observed for NaTi 2 (PO 4 ) 3 . The present study suggests that potential variations for Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C in R1 and R2 could be related to Fe 3+ /Fe 2+ and Ti 4+ /Ti 3+ redox couples, respectively.

To understand more deeply the effect of the substitution of Fe for Ti on the electrochemical process, we made use of [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] Fe Mossbauer spectroscopy during charge-discharge cycles in two different modes. Na insertion was performed in galvanostatic (C/25) and GITT (C/10 and 12 h open-circuit relaxation every 4 h) regimes, respectively, in order to investigate possible kinetic effects and improve the signal-to-noise ratio of the Mossbauer spectra. In the galvanostatic regime, the Mossbauer spectra were recorded every 0.08 Na/f.u. during the first cycle in the range of 1.6-3 V. All the spectra were fitted with two doublets assigned to Fe 3+ and Fe 2+ , respectively. Due to the poor signal-to-noise ratio of the spectra, unexpected randomly large variations of the isomer shift and quadrupole splitting were observed. This is mainly due to the small amount of [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] Fe in the sample and the short recording time for each spectrum (4 h). However, it was possible to extract values of the relative Na/f.u. (Figure 11a). More spectra were recorded at the 796 beginning of the second discharge, and this peak grows more 797 progressively (Figure 11c). Through the charge, the spectra 798 change significantly after the extraction of about 1 Na/f.u., 799 showing a decrease of the peak at 2.2 mm s -1 until the end of 800 the charge (Figure 11b). All the spectra were fitted with three arising from the quality of the spectra (Table 6). The average value is δ = 1.23 ± 0.04 mm s -1 . The quadrupole splitting exhibits stronger variations that could be related to changes in the local environments of the Fe atoms due to the continuous filling (emptying) of the M2 sites by Na along the discharge (charge). However, these variations do not exhibit a significant trend, which suggests a random occupation of the M2 sites. The average value is Δ = 2.05 ± 0.12 mm s -1 . The value of the isomer shift is close to that found for the Fe 2+ impurity, while the quadrupole splitting is significantly lower. However, these two parameters are both typical of high spin Fe 2+ . Almost all the Fe 3+ are transformed into Fe 2+ at the beginning of the discharge for less than 0.8 Na/f.u and only several percent of Fe 3+ remained in the second part of the discharge (Figure 10). This mechanism is reversible, since the amount of Fe 3+ increases only after the extraction of 1 Na/f.u. At the end of the charge, about 10 at % Fe 2+ remained in addition to the Fe 2+ impurity, which proves that sodium extraction was not complete. However, such a value could be due to the specific preparation of the electrode material for Mossbauer measurements as described in the Experimental and Theoretical Methods section (powdered sample with only carbon black additive) and the geometry of the in situ cell. This explains, for example, the observed maximum amount of 1.6 Na/f.u. inserted Na at the end of the first discharge. During the charge, 1.55 Na/f.u. were extracted and this difference of 0.05 Na/f.u. could affect 10% of the Fe 2+ atoms within the composite in line with the remaining amount observed by Mossbauer spectroscopy. Finally, the second discharge shows the decrease of the amount of Fe 3+ for Fe 2+ , as along the first discharge, but more accurately since more spectra were recorded (Figure 10). Compared to the galvanostatic mode, it is worth noting that the GITT mode provided more reliable experimental values of the Mossbauer subspectrum areas, as a result of a better signal-to-noise ratio, and larger amounts of Fe 2+ formed during the discharge probably due to kinetic effects.

For both galvanostatic and GITT regimes, the operando Mossbauer spectra clearly show the progressive reduction from Fe 3+ to Fe 2+ at the beginning of discharge. In the GITT mode, about 95% of Fe 3+ was transformed along the discharge until a Fixed. b Constrained to be equal.

2. 6 .

 6 Electrochemical Measurements. The electrochemical tests were performed in coin cells. The working electrodes were prepared by mixing 75 wt % active material, 15 wt % Super P carbon powder, and 10 wt % PVDF (polyvinylidene fluoride) binder in NMP (1-methyl-2-pyrrolidinone). The mixture was spread on Al foil substrate and dried at 120 °C for 12 h under vacuum to remove solvent and water. Coin cells were assembled inside filled drybox under controlled argon atmosphere by using a sodium foil as the counter electrode, a borosilicate glass microfiber sheet (Whattman) as the separator, and an organic electrolyte with 1 M NaClO 4 dissolved in a solution of propylene carbonate (PC) and 5 vol % of fluoroethylene carbonate (FEC) as the additive. The cells were cycled in a voltage range between 1.6 and 3.0 V, under galvanostatic conditions with C/n current rate (1 Na/f.u. in n hours), where n = 10, 5, 2, 1, and 0.2.

and 1

 1 389 determine the most stable configuration of the Na atoms in the 390 M1 and M2 interstitial sites and to evaluate the structural 391 parameters. There are 28 possibilities to place the 2 Na atoms 392 over the 8 interstitial sites of the rhombohedral cell of 393 NaTi 2 (PO 4 ) 3 . However, only several starting configurations 394 have been considered because of the symmetries of the initial 395 rhombohedral primitive cell. Previously published experimental 396 lattice parameters and atomic positions obtained for 397 NaTi 2 (PO 4 ) 3 41 were used as starting values of the structural 398 optimizations except for Na in M2, where the coordinates 399 (0.65, 0, 0.25) close to those obtained for Na 2 FeTi(PO 4 ) 3 27 400 were considered. On the basis of the total energy values 401 obtained for the optimized structures, three types of Na 402 configurations were distinguished: 2 Na in M1, and 1 Na in M1 403 Na in M2, and 2 Na in M2. The total energy obtained for 404 Na in M1 was about 0.3 eV/f.u. lower than that obtained for Na 405 in both M1 and M2 and about 0.5 eV/f.u. lower than that 406 obtained for Na in only M2. This clearly indicates that the most 407 stable NaTi 2 (PO 4 ) 3 phase predicted by DFT is obtained with 408 the full occupation of the M1 sites by Na in agreement with 409 previously published experimental data. In addition, the 410 theoretical volumes of the rhombohedral primitive cells are 411 456.2, 461.7, and 466.7 Å 3 for M1, M1+M2, and M2 412 occupations, respectively. The comparison with the averaged 413 experimental value (452.5 Å 3 ) indicates that the best agreement 414 is also obtained for Na in M1. This agreement is slightly better 415 by considering improved computational parameters (see 416 Section 2) since, in this case, the rhombohedral cell volume 417 is V rho = 455.9 Å 3 , which overestimates the experimental 418 volume by less than 1%. The theoretical atomic positions and 419 the main interatomic distances of the most stable phase are 420 similar to the experimental data obtained in this work (

  485 local relaxations, all the bond lengths and bond angles in 486 Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 are distributed leading to local distortions 487 t3

  FeO 6 local 488 environment shows significant variations from perfect octahe-489 dron (Figure 1b). The average values of Ti-O and Fe-O 490 interatomic distances in TiO 6 and FeO 6 octahedra are 1.94 and 491 2.00 Å, respectively. This increase is in line with that of the 492 ionic radius from 0.605 Å for Ti 4+ to 0.645 Å for Fe 3+ . The 493 average P-O and Na(M1)-O interatomic distances are close 494 to those in NaTi 2 (PO 4 ) 3 , but there are local distortions around 495 Na(M1) and Na(M2) (Table

  DFT expected crystallographic site a (Å), c (Å) 8.582, 21.685 V rho (Å 3 ) 461.0 Na(M1) -0.0103, -0.0079, -0.0013 0, 0, 0 (6b) Na(M2) 0.6347, 0.0006, 0.2470 x, 0, 1/4 (18e) M -0.0012, -0.0002, 0.1458 0, 0, x (12c) P 0.2884, -0.0004, 0.2508 x, 0, 1/4 (18e) O(1) 0.1775, 0.9810, 0.1926 x, y, z (36f) O(2) 0.1921, 0.1663, 0.0874 x, y, z (36f) a The expected corresponding general coordinates of the crystallographic sites in R3̅ c are given for comparison.

Figure 1 .

 1 Figure 1. (a) Theoretical optimized rhombohedral unit cell of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 with Na(M1) (light green), Na(M2) (green), Fe (red), P(pink), Ti (blue), and O (yellow). The dots indicate the interatomic distances (in Å) between the cations located along the hexagonal axis c hex . Local structure of Fe bonded to 6 O first-nearest neighbors forming a distorted octahedron and the two closest Na in M1 and M2. (b) The interatomic distances (in Å) and the bond angles (in degree) are also shown.

  constants and atomic coordinates. For example, the experimental volume of the rhombohedral cell (456.8 Å 3 ) is only 1% lower than the theoretical one (461.0 Å 3 ). In addition, the theoretical increase of 1.05% of the unit cell volume from NaTi 2 (PO 4 ) 3 to Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 matches the present experimental value of 0.95%. Finally, the experimental occupation rates of Ti, Fe, and Na are close to the predicted values. This is a clear indication of the accuracy of the theoretical model and confirms the validity of our structural analysis given above, including the random substitution of Fe for Ti and the Na occupation of the M1 and M2 sites. The XRD pattern of the composite Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C is similar to that of the crystalline phase Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3

Figure 2 .

 2 Figure 2. XRD patterns of (a) Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 and (b) Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C: experimental data (red •), Rietveld refinement (black line), Bragg positions (green bars), and difference between experimental data and fitted curve (blue line).

aFigure 3 .

 3 Figure 3. 57 Fe Mossbauer spectra of (a) Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 and (b) Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C at room temperature. In the later case, the Fe 3+ and Fe 2+ subspectra are plotted in red and green, respectively.

  633 aggregation of smaller particles of several hundred nanometers 634 (Figure 5b). This indicates that sucrose pyrolysis produced 635 porosity and reduced the particle size, improving the Na + 636 diffusion length. The TEM image of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C 637 f6 shows light-gray zones at the surface of the particles (Figure 638 f6 6b) that were not observed for Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 (Figure 639 6a) and can be assigned to carbon. The amount of carbon was 640 estimated from TGA to 1.5 wt %. Thus, the sucrose pyrolysis 641 produced carbon coating, which improved the electrical 642 conductivity of the composite and decreased the polarization. 643 The BET specific area of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 and 644 Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C are 0.17 and 20 m 2 g -1 , respectively. 645 This strong increase of the specific area by 2 orders of 646 magnitude confirms the high porosity of the secondary particles 647 of the composite that is expected to improve electrolyte 648 impregnation and diffusion of Na + ions.

a

  Fe Mossbauer Parameters of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 and Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C at Room Temperature: Isomer Shift Relative to α-Fe (δ), Quadrupole Splitting (Δ), Line Width at Half Maximum (Γ), Relative Area of the Subspectra (RA), and Assigned Fe Sites δ (mm/s) Δ (mm/s) Constrained to be equal.

Figure 4 .

 4 Figure 4. Temperature dependence of the inverse magnetic susceptibility of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 (•) and Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 / C (red •).

Figure 5 .

 5 Figure 5. SEM images of (a) Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 and (b) Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C.

Figure 6 .

 6 Figure 6. TEM images of (a) Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 and (b) Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C.

  704copy confirmed the existence of graphitic carbon at the surface 705 of the Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 particles, which is crucial to 706 improve the electronic conductivity of the composite material. 707 3.3. Insertion Mechanisms and Electrochemical 708 Performances. The potential curve of the first cycle of 709 Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C-based electrode obtained in galvano-710 static mode (C/10) in the range of 1.6-3 V versus Na + /Na 0 is 711 f8 displayed in Figure 8a. The first discharge shows a continuous 712 decrease of the potential from 2.6 to 2.2 V in the range of 0-713 0.5 inserted Na/f.u. (region R1), followed by a "pseudo-714 plateau" slightly decreasing from 2.2 to 2.1 V in the range of 715 0.5-2 inserted Na/f.u. (region R2). The specific capacity for 716 the first discharge is 125 mA h g -1 , which is close to the 717 theoretical value of 128 mA h g -1 . This small difference can be 718 related to the capacity loss due to the 9 at % electrochemically 719 inactive Fe 2+ . A similar potential curve was obtained for the first 720 charge, except for an irreversible capacity loss of about 5 mA h g -1 . The incremental capacity curve of the first discharge 722 exhibits a rather broad band at about 2.3 V (labeled A) and a 723

Figure 7 .

 7 Figure 7. Raman spectra of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 (black line) and Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C (red line). Inset: curve fitting of D and G bands.

Figure 8 .

 8 Figure 8. Galvanostatic voltage curve (C/10, 1.6-3.0 V vs Na + /Na 0 ) for the first cycle of (a) Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C and (b) derivative curves for different numbers of cycles. R1 and R2 denote the regions of the two-step insertion mechanism.

  771 areas (RA) of the two components that can be assimilated to 772 the relative amounts of Fe 3+ and Fe 2+ by assuming they have 773 the same Lamb-Mossbauer factor at room temperature. The 774 variations of these relative amounts versus number of inserted 775 Na/f.u. show a strong increase of the Fe 2+ contribution until 776 80% for 0.5 inserted Na/f.u. followed by some erratic changes 777 f10 in the range of 80-95% until the end of the discharge (Figure 778 f10 10). The inverse trend is obviously obtained for the Fe 3+ 779 contribution. This mechanism is reversible since the Fe 2+ 780 amount shows a rather constant trend from the beginning of 781 the charge up to about 0.6 inserted Na/f.u. and then strongly 782 decreases until the end of the charge. In the GITT experiment, 783 the spectra were recorded every 0.4 Na/f.u. for the first cycle 784 and 0.14 Na/f.u. for half the second discharge. The recording 785 time of 12 h (open-circuit relaxation) was sufficient to extract 786 reliable values of the Mossbauer parameters and subspectrum 787 areas. As indicated in the previous section, the Mossbauer 788 spectrum of the pristine Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C is formed by 789 two components assigned to Fe 3+ (δ = 0.42 mm s -1 , Δ = 0.28 790 mm s -1 , RA = 91%) and Fe 2+ (δ = 1.2 mm s -1 , Δ = 2.4 mm s -1 , 791 RA = 9%), originating from high spin Fe 3+ in crystalline 792 Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 and Fe 2+ impurity, respectively. At the 793 beginning of the first discharge, there is a strong increase of the 794 intensity of a peak at 2.2 mm s -1 until the insertion of about 0.8 795 f11

  801 doublets. One doublet was assigned to Fe 3 + in 802 Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 and the Mossbauer parameters, δ = 0.42 803 mm s -1 and Δ = 0.28 mm s -1 , were kept fixed during the fitting 804 procedure of all the spectra. The second doublet was attributed 805 to an Fe 2+ impurity with the following values, δ = 1.2 mm s -1 , 806 Δ = 2.4 mm s -1 , RA = 9%, that were also kept fixed. The last 807 doublet was associated with Fe 2+ arising from the reduction of 808 Fe 3+ due to Na insertion, and its Mossbauer parameters were 809 allowed to vary. In the latter case, only small and insignificant 810 variations of the Mossbauer isomer shift with inserted Na 811 amount were observed that could be assigned to uncertainties 812 t6

Figure 9 .

 9 Figure 9. GITT potential curve of the first discharge of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C. R1 and R2 denote the regions of the twostep insertion mechanism.

Figure 10 .

 10 Figure 10. Variations of the relative areas of the subspectrum of Fe 3+ (red) and of the sum of the two subspectra assigned to Fe 2+ in Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C and to Fe 2+ impurity (blue). The results are given for operando galvanostatic (circles) and GITT (squares) modes. The solid lines are guides for the eye.

Na 1 . 5

 15 852 0.8 inserted Na/f.u., which suggests that the remaining of the 853 discharge is mainly concerned with the Ti 4+ for Ti 3+ change. 854 This two-step mechanism, involving first the Fe 3+ /Fe 2+ change 855 followed by the Ti 4+ /Ti 3+ reduction, is in line with the 856 electrochemical results and can be related to the peaks A and B 857 of the derivative curve (Figure 8b). In the same way, the 858 beginning of the charge corresponds to the Ti 3+ /Ti 4+ oxidation, 859 indicating that the mechanism is reversible. The variations of 860 the relative amounts of Fe 3+ and Fe 2+ obtained in galvanostatic 861 and GITT regimes at the beginning of first or second discharge 862 are quasi-linear and match the expected linear variations for the operando experiments provide quantitative 866 evidence that reaction 4 proceeds at the beginning of discharge 867 in all the particles of the composite. This unambiguously proves 868 the one-phase reaction suggested by the shape of the 869 galvanostatic and GITT potential curves. However, it is 870 worth noticing that reaction 4 should be fully completed in 871 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C for 0.

  884retention can be related to the good reversibility of the charge-885 discharge cycles as shown by the stability of both the position 886 and the shape of the peaks A and B of the derivative curves 887 during cycling (Figure 8b). The rate capability test was 888 performed by 5 cycles at each rate C/10, C/5, C/2, 1C, 2C, and 5C, to return to the same rates 2C, 1C, C/2, C/5, and C/

Figure 11 . 3 and

 113 Figure 11. Operando 57 Fe Mossbauer spectra of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C recorded during the first cycle and at the beginning of the second discharge in GITT mode.

Figure 12 .

 12 Figure 12. Specific capacity and Coulombic efficiency of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C measured in the galvanostatic mode at C/10 in the range 1.6-3.0 V vs Na + /Na 0 .

Figure 13 .

 13 Figure 13. Rate and cycling performances of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C.

Table 421

 421

1). It is worth noting that such an excellent agreement was

Table 1 .

 1 Theoretical (DFT) and Experimental (XRD) Values of the Structural Parameters of NaTi 2 (PO 4 ) 3 : Lattice Parameters (a, c) and Atomic Coordinates (x, y, z) Given in the Hexagonal System, Rhombohebral Unit Cell Volume (V rho ), and Interatomic Distances (d)

		DFT	XRD
	a (Å), c (Å)	8.521, 21.748	8.482(9), 21.794(3)
	V rho (Å 3 )	455.9	452.5
	Ti (x,y,z)	0, 0, 0.1442	0, 0, 0.1447
	P (x,y,z)	0.2870, 0, 0.25	0.2857, 0.0000, 0.25
	O(1) (x,y,z)	0.1754, 0.9803, 0.1928	0.1740, 0.9755, 0.1926
	O(2) (x,y,z)	0.1925, 0.1652, 0.0878	0.1910, 0.1661, 0.0884

Table 2 .

 2 Averaged Theoretical Values of the Structural Parameters of the Most Stable Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 Phase Obtained by DFT Calculations: Lattice Parameters (a, c), Atomic Positions (x, y, z) Given in Hexagonal System, Rhombohedral Unit Cell Volume (V rho ) a

Table 3 .

 3 Theoretical Interatomic Distances d i (i = 1,n where n is the Number of Nearest Neighbors), Average Distances d av and Polyhedron Distance Distortions Δ poly Obtained from DFT Calculations for MO 6 Octahedra with M = Ti or Fe (n = 6), Na(M1)O 6 (n = 6), and Na(M2)O 8 (n = 8) Polyhedra of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3

							a
				interatomic distances (Å)				d av (Å)	Δ octa
	Ti1-O	1.871	1.872	1.948		1.957	2.025	2.027		1.950	1.05 × 10 -3
	Ti2-O	1.891	1.906	1.914		1.953	1.959	1.961		1.931	2.08 × 10 -4
	Ti3-O	1.905	1.936	1.938		1.951	1.966	1.974		1.945	1.36 × 10 -4
	Fe-O	1.914	1.915	1.993		2.012	2.088	2.089		2.002	1.27 × 10 -3
	Na1-O	2.374	2.388	2.414		2.458	2.555	2.600		2.465	1.20 × 10 -3
	Na2-O	2.343	2.410	2.414		2.427	2.475	2.583		2.442	9.10 × 10 -4
	Na3-O	2.294	2.343	2.374	2.400	2.409	2.532	2.741	2.913	2.501	6.63 × 10 -3

Table 4 .

 4 Structural Parameters of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 Obtained from Rietveld Refinement a

	space group	R3̅ c
	a = 8.520(8) Å
	c = 21.794(2) Å
	vol = 456.8(1) Å 3
	atom	site symmetry	x	y	z	occupation rate	B iso (A 2 )
	Na(1)	6b	0.0	0.0	0.0	0.9799(1)	4.627 (1)
	Na(2)	18e	0.6376(9)	0.0	1/4	0.5200(9)	2.374(1)
	Ti(1)	12c	0.0	0.0	0.1439(4)	1.500(1)	2.375(1)
	Fe(1)	12c	0.0	0.0	0.1439(4)	0.500(1)	2.375(1)
	P(1)	18e	0.2875(9)	0.0	1/4	3.000(1)	1.909(1)
	O(1)	36f	0.1747(8)	0.9759(1)	0.1926(2)	6.000(1)	0.207(1)
	O(2)	36f	0.1979(1)	0.1668(6)	0.0883(4)	6.000(1)	0.281(1)

Table 6 .

 6 [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] Fe Mossbauer Parameters of Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C-Based Electrode at Different Stages of the Electrochemical Discharge-Charge Process in Operando GITT Mode (See Text for Details): Isomer Shift Relative to α-Fe (δ), Quadrupole Splitting (Δ), Line Width at Half Maximum (Γ), Relative Area of the Subspectra (RA), and Assigned Fe Sites first discharge: number of inserted Na/f.u. δ (mm/s) Δ (mm/s) Γ (mm/s)

	RA
	(%)	site

a 0.460(1) b 0.5 Fe 3+ [Oh]

DOI: 10.1021/acs.jpcc.5b07931 J. Phys. Chem. C XXXX, XXX, XXX-XXX I

The performances can also be related to the low internal impedance leading to a low potential polarization for the charge-discharge cycles. The substitution of Fe for Ti produces more complex insertion mechanisms than the two-phase and one-phase mechanisms observed for NaTi 2 (PO 4 ) 3 and Na 3 Fe 2 (PO 4 ) 3 , respectively. [START_REF] Butt | Raman Spectroscopy of Superionic Ti-Doped Li 3 Fe 2 (PO 4 ) 3 and LiNiPO 4 Structures[END_REF] Fe Mossbauer spectroscopy used in operando mode quantitatively shows linear variations of the Fe 3+ and Fe 2+ amounts in Na 1.5 Fe 0.5 Ti 1.5 (PO 4 ) 3 /C at the beginning of discharge and at the end of charge for less than about 0.5 inserted Na/f.u. in line with one-phase mechanism suggested by the potential profiles.

■ AUTHOR INFORMATION

Corresponding Author *E-mail: lippens@univ-montp2.fr.

Notes

The authors declare no competing financial interest.


Xavier Sauvage
email: xavier.sauvage@univ-rouen.fr

W Lefebvre

C Genevois

S Ohsaki

K Hono

Complementary use of TEM and APT for the investigation of steels nanostructured by severe plastic deformation

Keywords: Atom Probe Tomography, Transmission Electron Microscopy, Severe Plastic Deformation, Nanostructures, Steels

The properties of bulk nanostructured materials are often controlled by atomic scale features like segregation along defects or composition gradients. Here we discuss about the complimentary use of TEM and APT to obtain a full description of nanostructures. The advantages and limitations of both techniques are highlighted on the basis of experimental data collected in severely deformed steels with a special emphasis on carbon spatial distribution.

Introduction

Strength of materials has been known to increase as inverse proportional to the crystal grain size following the Hall-Petch relationship. Thus, continuous efforts have been made to improve the strength of structural materials by refining crystal grain sizes by controlling recrystallization by thermomechanical treatments. In steels and aluminum alloys, fine grain size means 1 to 10 m, and there have not been effective processing technique to reduce the grain size of these materials to less than 100 nm until recently. Although mechanical milling and alloying can process the powders containing nanosized grains, grain growth cannot be suppressed during consolidation processes like sintering and hot pressing. Therefore, processing bulk nanostructured materials (BNM) for structural applications is still a big challenge, in particular using an industrially viable process. Recent developments of various severe plastic deformation (SPD) processes [1,2] like equal channel angular pressing (ECAP), accumulative roll bonding (ARB), and high pressure torsion (HPT) straining opened up a new way to process nanosized grain microstructures to steels and aluminum alloys, and it is in principle capable of scaling up ECAP and ARB to an industrial scale. The main advantage of these SPD techniques is that materials are free of porosity unlike the powder consolidation processes. The final grain size is typically less than 300 nm in most of metals and alloys which gives rise to a significant strengthening; however the main drawbacks of BNMs are the lack of ductility and of thermal stability [3]. Therefore, various attempts are being made to achieve both high strength and ductility by modifying the nanostructures.

Many studies on SPD have been applied to single phase materials to refine grain size.

Applications of SPD to multiphase microstructures have also been attempted to explore the possibility of obtaining ultrafine composite nanostructures, e.g. [4][5][6][7][8][9][10]. One classical example of severe plastic deformation of two phase alloys can be seen in well known pearlite wire that is widely used as suspension cables, tire cords and fishing wires. By cold drawing pearlite to a high strain of more than 4, the cementite and ferrite lamellar spacing is reduced to a few tens nanometers, and a strength exceeding 3 GPa can be commonly obtained [11][12][13]. Since cementite lacks plasticity, they are fragmented in the initial drawing stage and eventually decomposed to carbon and ferrite. A similar concept was used to develop ultrahigh strength electric wires for pulsed magnets by cold drawing Cu-Ag eutectic microstructures. Although Cu and Ag are insoluble each other, they are considered to be mixed each other after large strain deformation like mechanical alloying [14]. To observe such decomposition and mixing phenomena that occur during severe plastic deformation of multiphase alloys, analytical techniques that have a rather high spatial resolution and quantitative analytical capabilities are required. Transmission electron microscopy (TEM) with energy dispersive X-ray spectroscopy (EDS) as well as electron energy loss spectroscopy (EELS) is the most versatile technique for qualitative analysis of nanostructures processed by SPD, but atom probe tomography (APT) has superior spatial resolution as well as quantitative analytical capability of embedded particles and light elements. Effective complimentary applications of TEM and APT will provide critical information regarding the mechanism of nanostructure evolution of multi-phase alloys processed by SPD. In this article, we critically overview the current problems as to the nanostructures of multiphase alloys processed by SPD, using as example the typical case of steels.

APT and TEM

APT is a projection type field ion microscope (FIM) combined with a time of flight mass spectrometer [START_REF] Miller | Atom Probe Tomography[END_REF]. Atoms are ionized from a hemispherical surface of a sharp needle specimen (FIM tip) of approximately 50 nm in radius by the field evaporation phenomenon.

Positively charged ions accelerated along the electric field are detected on a position sensitive detector located in front of the specimen. Since atoms are field-evaporated layer-by-layer, APT has an atomic layer resolution in the depth direction of the elemental map with optimized analysis conditions. The mass resolution depends on the energy deficit of ions as well as the flight length, but is usually improved by energy compensator like reflectrons [START_REF] Bémont | [END_REF].

The area for analysis depends on the acceptance angle 2 of the detector which is determined by the tip radius, tip detector distance and the diameter of the detector (~tan -1 (D/d), AA~2r 2 , where D: diameter of detector, d: tip detector distance, AA: analysis area, r: tip radius). Because of recent wide angle configuration of APT, typical size of analysis volumes has expanded from 10x10x200 nm 3 in the 90's to 50x50x200 nm 3 nowadays, meaning that the number of collected atoms has expanded from typically half a million to several tens of millions. Although it is possible to achieve an atomic resolution in the depth direction, the lateral resolution is subject to the evaporation aberration, which is about 0.3 nm. Nevertheless, the APT is the only analytical microscope that can provide three dimensional real elemental maps of constituent elements in alloys.

During the last twenty years APT has been used widely to measure the composition of nanoscale particles in various metallic materials and to display their distribution and morphology. It is also a powerful instrument to reveal chemical gradients or to highlight segregation along structural defects like grain boundaries or dislocations. Recently, it has been demonstrated that femtosecond laser could successfully assist the field evaporation of materials analyzed by APT [17]. It reduces the stress applied to the specimen, thereby reducing the probability of specimen failure. This feature is extremely important for the atom probe analysis of nanostructured materials that are prone to rupture during voltage pulse field evaporation due to the presence of a high number density of defects like dislocations and grain boundaries. APT specimens have been prepared for a long time by standard electropolishing techniques but nowadays micro-milling by the focused ion beam (FIB) technique is widely used [18], which made site specific specimen preparation possible, for example from surface layers, grain boundaries and powders [19].

Compared to APT, TEM provides images of micro-or nano-structures with a larger field of view, and chemical mapping can easily be achieved by energy filtering TEM (EFTEM) or scanning TEM (STEM) using EDS, high angle annular dark field (HAADF) or EELS.

Moreover, crystallographic information can be collected using electron diffraction techniques or high resolution TEM (HRTEM) images. Nevertheless, one should note that both EDS and EELS are rather qualitative. This makes difficult the measurement of concentrations typically lower than 1at.% in nanoscaled grains or the detection of segregation along crystallographic defects like grain boundaries. Moreover, the composition of nanosized particles embedded in the matrix cannot be quantitatively measured. Although TEM tomography was recently developed to build up the three-dimensional tomography of microstructures, most of the case they give only morphological information without any chemical information [20].

Complimentary TEM and APT analysis of cold drawn pearlitic steel

Cold drawn pearlitic steel wires are among the strongest commercial steels (yield stress up to 3GPa) and are widely used in the tire industry [11][12][13]. This extremely high strength results from their unique nanostructure that is progressively formed during the cold drawing process: cementite lamellae are aligned along the wire axis, and the interlamellar spacing is reduced in a range of 10 to 30 nm, while the thickness of the lamellae is typically less than 5 nm [11][12][13].

At the end of the drawing process, the typical cumulated true strain applied to the material is about 3 for commercial products and higher than 4 for laboratory material. It was often reported that such high level deformation also leads to the decomposition of cementite [21][22][23][24][25][26][27]. This feature was first discovered by Mössbauer spectroscopy and then confirmed by APT measurements in the late 90's. However, the physical mechanisms leading to this decomposition are still under debate [9,12,[27][28][29], especially because the description of the resulting nanostructure and of the carbon distribution is not fully understood. In some cases, carbon supersaturated solutions in the -Fe matrix exhibiting some features similar to quenched martensite have even been reported [22,23,30].

The driving force and the kinetics of this phase transformation are still controversial. Some authors point out the possible effect of moving dislocations that cross cementite lamellae.

They may trap carbon atoms of the lamellae and then redistribute them inside the ferrite phase [21]. Others argue that the driving force for the decomposition could be the dramatic increase of the Fe 3 C/-Fe interfacial area resulting from the lamellae thinning. This feature may lead to a strong increase of the carbon solubility in the ferrite through the well known Gibbs-Thomson effect [12,27]. Cementite decomposition is still a challenging issue since it strongly affects the deformation mechanisms (dislocation nucleation and motion) in the drawn wires, and thus its mechanical properties and especially the ductility and the ageing behaviour. In spite of the large amount of experimental data published in the literature, the influence of metallurgical parameters (original interlamellar spacing, pearlite colony size, alloying elements) and of processing parameters is not yet well established. In the following, we demonstrate that linking HR-STEM data to advanced APT data treatment lead to a better understanding of such unique nanostructure.

The nanostructure of drawn pearlitic steels is usually imaged using bright field TEM, but as shown in Fig. 1, it can also be observed by FIM (Fig. 1(a)) or HAADF STEM (Fig. 1(b)). The dim contrast of cementite in the FIM image results from the concavity of the specimen surface due to the lower evaporation field of cementite compared to that of the -Fe matrix. When the image is observed at a lower temperature like 20K, these contrasts disappear. This means that no information on the concentration of carbon in each phase can be obtained from the FIM images. On the other hand, HAADF images are sensitive to the atomic number, so the darkly imaged regions in Fig. 1 (b) are considered to contain higher amount of carbon compared to the matrix. Thus, in the HAADF image, the contrast of the recorded images can be linked to the carbon concentration gradient within the nanostructure. This indicates that even if cementite is partly decomposed as it was pointed out by Mössbauer spectroscopy long time ago [21], a lamellar structure is kept and carbon atoms are not homogeneously distributed.

Using APT, it is possible to obtain the three-dimensional carbon map of such a lamella (Fig. 2(a)). These data confirm that most of carbon atoms are located in thin lamellae of cementite or former cementite with a thickness in a range of 2 to 4 nm. However, as seen in the concentration profile (Fig. 2 (c)) that was calculated across the lamella of the volume displayed in the Fig. 2 (a), the carbon concentration in the -Fe phase is in a range of 0.5 to 1 at.%, i.e. much higher than the equilibrium solubility limit. Note that the maximum carbon concentration is well below the 25at.% for the stoechiometry of Fe 3 C, confirming the strain induced decomposition of this carbide. However, due to the lower evaporation field of cementite, local magnification effects and ion trajectory aberrations could significantly affect the distribution of carbon atoms within the reconstructed volume [31]. Thus, the large carbon gradients exhibited on each side of the interface by the composition profile (Fig. 2(c)) may not reflect the actual concentration of carbon.

To check this point, EELS mapping of a carbon rich region with a similar interlamellar spacing was performed (Fig. 3(a)). The profile of the intensity of the carbon K edge (Fig. 3(b)) was computed (with the three windows method) across the lamella with a sampling volume similar to the one used for the carbon concentration profile of the APT data (Fig.

2(c))

. Due to carbon contamination on the TEM sample surface, the carbon concentration cannot be accurately measured in the -Fe phase. Despite this, the profile unambiguously shows that there are large carbon gradients in the -Fe phase. The gradients spread over a distance of about 2 nm in each side of the interface, in accordance with the APT data. Thus, we can conclude that local magnification effects do not significantly affect the distribution of carbon atoms within the APT reconstructed volume. From the carbon EELS map (Figure 3(a)), it is clear that the carbon concentration is not homogeneous within the probed lamella (there are less bright dots on the left). A careful treatment of the APT data can confirm this point too. The original data set displayed in the Fig. 2(a) was filtered to highlight in grey the regions containing more than 10at.% carbon (see details of the procedure in [START_REF] Sauvage | Ultrafine Grained Materials III[END_REF]). Some clusters with a diameter in range of 2 to 5 nm do appear within the carbon rich lamellae. They could be some remaining cementite nano-particles. This inhomogeneous concentration fluctuation of carbon in the former cementite lamellae is considered to be due to the fragmentation of the cementite lamellae into nanocrystalline cementite as was observed by the dark field TEM image in [25].

Therefore, HR-STEM was carried out to image the nanostructure at the atomic scale in a region where the interlamellar spacing is only about 10 nm (Fig. 4). Carbon rich lamellae are darkly imaged on the HAADF image (arrowed in the Fig. 4(b)). Due to lattice distortions, the bright field STEM image (Fig. 4(a)) is difficult to interpret directly thus it was Fourier filtered (Fig. 4(c)). Nine dislocations are arrowed (in white), but due to distortions all of them are probably not visible in this image. Thus, in this region, the dislocation density is at least 2 10 16 m -2 . One should note that dislocations are not exclusively located in the regions with the highest carbon concentration (dark zones in the HAADF image). Thus, one can conclude that carbon atoms resulting from the cementite decomposition are not homogeneously distributed on interstitial sites in the bcc lattice of the -Fe phase but obviously trapped by stress fields of dislocations. This confirms the critical role of dislocations in the decomposition process of cementite [9,21,28].

This can be further supported by the APT results of carbon distribution in both mechanically milled pearlite and mechanically milled martensite. As reported by Ohsaki [START_REF] Ohsaki | [END_REF], cementite in pearlite powders are completely decomposed into carbon, resulted in the formation of nanocrystalline ferrite. Figure 5 shows (a) TEM, (b) FIM, and (c) APT results of Fe-3.8at.%C pearlite powder that was mechanically milled for 100 h. The TEM image shows complete decomposition of cementite. The ring pattern observed inside the {110} ring is due to the formation of oxide, and it does not correspond to cementite. The FIM image shows nanocrystalline features of ferrite, i.e., the brightly imaged regions correspond to ferrite and the darkly imaged regions correspond to the grain boundaries. The contrast mechanism is similar to that for Fig. 1 (a). The APT data show that the carbon distribution is not uniform regardless of the complete decomposition of cementite; carbon atoms are segregated along the grain boundaries. Within the ferrite, a small amount of C is detected, but majority of carbon are detected from the grain boundaries. This suggests that carbon tend to be segregated at defects. When nanocrystalline ferrite is formed, most of dislocations are swept away from the nanocrystals to form high angle grain boundaries; hence the carbon atoms segregated at dislocations are also dragged to the grain boundaries. This scenario can be further supported by the carbon distribution in nanocrystalline ferrite that was formed by mechanical milling of martensite powder. Figure 6 shows (a) TEM, (b) FIM and (c) APT results obtained from the nanocrystalline ferrite powder that was processed by mechanically milling Fe-3.6at.%C martensite steel powder [34]. Since it was martensite before mechanical milling, we assume that C distribution was uniform in the unmilled condition. Although the nanostructure formation does not involve Fe 3 C decomposition, after long time milling that is enough to produce nanocrystalline ferrite, the microstructure observed by TEM, FIM and APT are nearly the same as was produced from the mechanically milled pearlite. This suggests that the final microstructures were nearly the same; i.e., most carbon atoms are segregated along the grain boundaries with a small amount of carbon in the ferrite. Since the carbon distribution was uniform in the original martensite structure, the carbon depletion inside ferrite nanocrystals strongly suggests that it was swept away by dislocations to form high angle grain boundaries. Since there is little chance for a dislocation can be present within the nanocrystalline ferrite, carbon dissolution in the nanocrystalline ferrite is considered to be low.

Concluding remarks

TEM is without any doubt a more versatile characterization technique than APT, especially because it provides both structural and chemical information. Moreover, recent developments in analytical TEM (STEM, EELS, EDS) have led to chemical analysis down to the atomic column scale. However, there are still numerous applications where the complimentary use of APT is effective, especially for nanocrystalline alloys, in particular when low concentrations of light elements play a critical role in nanostructuring mechanisms or stability. Indeed, only this technique is able to quantitatively analyze nanoscaled features that are embedded in the matrix phase and smaller than the thin foil thickness. In the present paper, we highlighted a unique nanostructure feature of cold drawn pearlitic steels, and demonstrated that the complementary characterization by TEM and APT is useful. The solute segregations along structural defects is also a typical example where the APT has a superior analytical ability. It is especially relevant for nanostructured alloys with an inherent high volume fraction of grain boundaries and it is also a critical point for nanostructured materials processed by SPD with high dislocation densities. However, APT reconstructions are not always free from artifacts like local magnification effects, which may affect particle shapes and related composition gradients. As shown in this paper, analytical STEM may help to clarify doubts. It points out once more the evident complimentarily of APT and TEM that may continue to grow in the future, especially in the field of nanostructured alloys. Today, the main drawback is that TEM foils and APT needle shape samples are very different, thus one cannot analyze in optimal conditions exactly the same feature with both microscopes. However, recent instrumental developments like FIB nano-machining [18] or TEM specimen holder designed for tips [35] may push ahead the complimentary use of APT and TEM, especially in the field of tomography at the nanometer scale [36]. Reproduced from reference [34] Figure 1 13

Figure captions

Figure 1

 1 Figure 1 Cold drawn pearlitic steel (Fe-3.6at.%C, true strain 3.5) (a) FIM image showing the nanostructure of the cold drawn pearlitic steel. Cementite lamellae (arrowed) have a dark contrast due to their lower evaporation field comparing to the -Fe matrix. (b) STEM HAADF image of the same nanostructure. Cementite lamellae (arrowed) have as well a dark contrast because the atomic number of C (Z=6) is much smaller than that of Fe (Z=26).

Figure 2

 2 Figure 2 Cold drawn pearlitic steel (Fe-3.6at.%C, true strain 3.5) (a) High resolution STEM BF image of the nanostructure with FFT (inset). (b) STEM HAADF image of the same region where dark lamellae corresponding to cementite are arrowed. (c) Fourier filtered image of the BF. Some dislocations are arrowed in white.

Figure 3

 3 Figure 3 Cold drawn pearlitic steel (Fe-3.6at.%C, true strain 3.5) (a) Carbon K edge STEM EELS map, bright dots correspond to high carbon concentration. (b) EELS profile computed across the carbon rich lamella exhibited on the EELS map (obtained in a section of 5 nm as imaged on the EELS map).

Figure 4 Figure 5

 45 Figure 4 Cold drawn pearlitic steel (Fe-3.6at.%C, true strain 3.5) (a) 3D reconstructed volume of a small volume analyzed by APT. Only carbon atoms are plotted to exhibit a carbon rich lamella. (b) Same volume with regions containing more than 10at.% highlighted in grey. (c) Composition profile computed across the lamella (averaged across a section of 5x5nm 2 as imaged on the reconstructed volume and with a sampling volume thickness of 1 nm)

Figure 6

 6 Figure 6 (a) TEM, (b) FIM, and (c) APT map of the nanocrystalline ferrite produced from the Fe-3.6at.%C martensite steel wire by mechanically milling for 100 h.

Figure 2 14 Figure 3

 143 Figure 2

Figure 4 15 Figure 5

 155 Figure 4

Figure 6

 6 Figure 6

Acknowledgements

Dr Eiji Okunishi from JEOL Ltd is gratefully acknowledged by the authors for the TEM observations. The authors thank Mr. Guillaume Lathus and Yuichi Matsumoto from JEOL Ltd for providing access to the JEOL 2100F TEM at Akishima (Japan). This work was in part supported by the Grant-in-Aid for Scientific Research on Priority Areas ''Giant Straining Process for Advanced Materials Containing Ultra-High Density Lattice Defects", MEXT, Japan, the World Premier International Research Center Initiative (WPI Initiative) on Materials Nanoarchitronics, MEXT, and CREST, JST.


R odents are the most diversifi ed order of wild mammals and are also the prevailing mammal lineage associated with human-inhabited socioecosystems. In addition to their destructive behaviors, rodents are involved in maintaining, disseminating, and transmitting zoonotic pathogens, impacting both animal and human health [START_REF] Colombe | A new approach to rodent control to better protect human health: first international meeting of experts under the auspices of WHO and the Pan American Health Organization[END_REF].

Hantaviruses are transmitted to humans by inhalation of aerosols contaminated by rodent excreta, including urine, feces, and saliva. They are the causative agents of hemorrhagic fevers and are considered emerging pathogens that impact public health worldwide. Hantaviruses are enveloped with a tripartite single-stranded RNA genome of negative polarity comprising small (S), medium (M), and large (L) segments.

Because hantaviruses are host-specifi c, their geographic distribution is tightly linked to that of their host. However, the emergence of hantaviruses in new geographic regions is still possible by the spread of the rodent reservoir [START_REF] Lin | Migration of Norway rats resulted in the worldwide distribution of Seoul hantavirus today[END_REF]. Transport-mediated dissemination of rodent-borne hantaviruses is of critical importance in their distribution and constitutes a critical health concern [START_REF] Klempa | Hantaviruses and climate change[END_REF]. Seoul virus (SEOV), an orthohantavirus fi rst identifi ed in South Korea in 1982, has had a particular impact on global human health attributable to its worldwide dispersal [START_REF] Lin | Migration of Norway rats resulted in the worldwide distribution of Seoul hantavirus today[END_REF][START_REF] Clement | Wild rats, laboratory rats, pet rats: global Seoul hantavirus disease revisited[END_REF]. Some outbreaks are hypothesized to be driven by the sporadic introduction of its now cosmopolitan host, the Norway or brown rat (Rattus norvegicus), at seaports or from pet and laboratory rats [START_REF] Childs | A case study of two rodentborne viruses: not always the same old suspects[END_REF][START_REF] Dupinay | Detection and genetic characterization of Seoul virus from commensal brown rats in France[END_REF]. Little is known concerning the circulation of SEOV in Africa, although a recent study reported its presence in southeastern Senegal [START_REF] Diagne | Seoul orthohantavirus in wild black rats, Senegal, 2012-2013[END_REF]. We screened rats in the Autonomous Port of Cotonou, Benin, to determine the presence of SEOV in these rodents.

The Study

We trapped rodents in the seaport of Cotonou using Sherman line capture traps and locally made wide mesh traps that were set for 3 consecutive nights in April 2018. We transported the animals to the Laboratoire de Recherche en Biologie Appliquée laboratory in closed containers and processed them the same day. We anesthetized the rodents with diethyl ether and subsequently euthanized them by cervical dislocation.

We screened blood samples from 32 brown rats and 37 house mice (Mus musculus) for the presence of hantavirus-reactive antibodies by using an immunofl uorescence assay as previously described [START_REF] Van Cuong | Rodents and risk in the Mekong Delta of Vietnam: seroprevalence of selected zoonotic viruses in rodents and humans[END_REF]. A total of 12 (37.5%) rats were seropositive, and the lungs of 2 rats were confi rmed positive by reverse transcription PCR as previously described [START_REF] Klempa | Hantavirus in African wood mouse, Guinea[END_REF]. This discrepancy cannot be explained by the presence of maternal antibodies because all seropositive rats were adults except for 1 subadult. However, the difference could be because of the limited sensitivity of the panhantavirus PCR or to a viral load decrease over time that fell below detection limits in rats that were infected several months earlier.

Genetic Characterization of Seoul Virus in the Seaport of Cotonou, Benin

We treated reverse transcription PCR-positive samples with DNase I (Thermo Fisher Scientific, https:// www.thermofisher.com), and purified samples with Agencourt RNA Clean XP magnetic beads (Beckman Life Sciences, https://www.beckmancoulter.com). We removed ribosomal RNA using a NEBNext rRNA depletion kit (New England BioLabs, https://www. neb.com). We prepared the sequencing library with a NEBNext Ultra II RNA library prep kit and quantified it using a NEBNext Library Quant kit for Illumina (Illumina, https://www.illumina.com). We sequenced pooled libraries on a MiSeq platform using a MiSeq v3 reagent kit with 300 bp paired-end reads. Raw sequence reads were trimmed and low-quality (quality score <15) or short (<36 nt) sequences were removed using Trimmomatic [START_REF] Bolger | Trimmomatic: a flexible trimmer for Illumina sequence data[END_REF]. The trimmed sequence reads were assembled against reference sequences (GenBank accession nos. NC_005237.1, NC_005236.1, NC_005238.1) using Bowtie2 algorithm [START_REF] Langmead | Fast gapped-read alignment with Bowtie 2[END_REF] and some in-house scripts.

We deposited S, M, and L segment sequences of strain Benin1368 into GenBank (accession nos. MW561221-3) and analyzed them by using BLAST (https://blast.ncbi.nlm.nih.gov). BLAST revealed that the best matches were with SEOV strain CSG5 (accession nos. AB618112-30) from Vietnam for the S segment (97.51% nt identity) and the M segment (97.70% nt identity) and with SEOV strain Lyon/Rn/ FRA/2013/LYO852 (accession no. KF387723) from France for the L segment (96.82% nt identity).

We performed phylogenetic analyses on 3 datasets composed of complete or nearly complete coding regions of S, M, and L segment sequences of SEOV from different geographic areas available in GenBank. We used sequences of Hantaan and Anjozorobe viruses as outgroups in all analyses. We performed phylogenetic analyses as previously described ( 12) using the general time-reversible plus gamma distribution plus invariant sites model (S and M segment) or the general time-reversible plus gamma distribution model (L segment).

The 3 datasets produced broadly concordant phylogenetic topologies (Figure 1, https://wwwncd.cdc. gov/EID/article/27/10/21-0268-F1.htm). All datasets grouped the SEOV strains from Cotonou within a cluster that included strains from Europe (France and Belgium) and from Southeast Asia (Indonesia, Singapore, Vietnam, and Cambodia), referred to as SEOV lineage 7 [START_REF] Ling | Genetic analyses of Seoul hantavirus genome recovered from rats (Rattus norvegicus) in the Netherlands unveils diverse routes of spread into Europe[END_REF]. Variants of this lineage belonged to SEOV phylogroup A; this group originated in China and subsequently spread to other parts of the world (2). More specifically, SEOV lineage 7 may reflect the historical connections between regions of Southeast Asia and France through critical trade routes [START_REF] Heyman | Seoul hantavirus in Europe: first demonstration of the virus genome in wild Rattus norvegicus captured in France[END_REF].

From Africa, only short SEOV sequences from conserved parts of S (226 nt) and L (347 nt) segments were available from wild black rats (R. rattus) from Senegal [START_REF] Diagne | Seoul orthohantavirus in wild black rats, Senegal, 2012-2013[END_REF]. For this reason, we did not include them in our datasets using complete segments. However, phylogenetic analyses on the basis of datasets including these short sequences place them in SEOV lineage 4 (S segment) or 3 (L segment), with low branch support (Figure 2, https://wwwncd.cdc.gov/EID/ article/27/10/21-0268-F2.htm). This inconsistency is potentially attributable to the short length and high conservation of these sequences; although it could indicate a distinct introduction event from Benin, this interpretation must be considered with caution because of the low level of phylogenetic information provided by these sequences. We calculated estimates of evolutionary divergence between strains from Senegal and Benin using MEGAX (http://www.megasoftware. net). Analyses showed 95.7% nt homology for the S segment and 94.8% nt homology for the L segment. Analyses showed 97.85% amino acid-level homology for the nucleocapsid protein and 100% amino acid-level homology for the RNA polymerase.

Conclusion

Because of insufficient testing for hantavirus infections and unreported mild cases [START_REF] Heyman | Seoul hantavirus in Europe: first demonstration of the virus genome in wild Rattus norvegicus captured in France[END_REF], the exact circulation of hantaviruses on the continent of Africa is unknown. Whereas SEOV is not widely considered a public health issue in Africa by local health authorities, the presence of SEOV-like agents in humans and wild rats is strongly suspected in at least 17 different countries (4). Recent and unambiguous sequencingbased identification of SEOV in Senegal [START_REF] Diagne | Seoul orthohantavirus in wild black rats, Senegal, 2012-2013[END_REF] and in Benin with our study confirms that SEOV should be anticipated as a possible cause of illness, such as hemorrhagic fever with renal syndrome. Seaports and ships have already been identified as potential entry points for hantaviruses [START_REF] Lin | Migration of Norway rats resulted in the worldwide distribution of Seoul hantavirus today[END_REF][START_REF] Lokugamage | Epizootiological and epidemiological study of hantavirus infection in Japan[END_REF], which is a likely cause in Africa, as our study shows. SEOV strains recovered from brown rats from the Cotonou seaport are phylogenetically similar to strains from Southeast Asia and Europe, regions where many maritime trade exchanges occur that could explain the presence of these strains in Cotonou. The accidental transportation of SEOV-carrying rats at seaports could lead to local emergence of SEOV infections among port workers. Regular sanitary control of rats within seaports could prevent rodentborne and arthropodborne pathogen dissemination through sea trade.

  Seoul virus is a zoonotic pathogen carried by the brown rat Rattus norvegicus. Information on its circulation in Africa is limited. In this study, the virus was detected in 37.5% of brown rats captured in the Autonomous Port of Cotonou, Benin. Phylogenetic analyses place this virus in Seoul virus lineage 7.

	of Helsinki, Haartmaninkatu 3, Helsinki FI-00290, Finland
	(R. Kant, T. Smura, T. Sironen); University of Helsinki, Agnes
	Sjö bergin katu 2, Helsinki FI-00790 (R. Kant, T. Sironen);
	Laboratoire de Recherche en Biologie Appliquée, Unité de
	Recherche sur les Invasions Biologiques, Ecole Polytechnique
	d'Abomey-Calavi, Université d'Abomey-Calavi, Cotonou, Benin
	(S. Badou, HJ. Dossou, J. Etougbétché, G. Houéménou); La-
	boratoire de Génétique Moléculaire et d'Analyse des Génomes,

Guillaume Castel, Ravi Kant, Sylvestre Badou, Jonas Etougbétché, Henri-Joël Dossou, Philippe Gauthier, Gualbert Houéménou, Teemu Smura, Tarja Sironen, Gauthier Dobigny 2704 Emerging Infectious Diseases • www.cdc.gov/eid • Vol. 27, No. 10, October 2021 DISPATCHES Author affi liations: CBGP, INRAE, CIRAD, IRD, Institut Agro, Univ Montpellier, Montpellier, France (G. Castel, S. Badou, J. Etougbétché, H.-J. Dossou, P. Gauthier, G. Dobigny); University Faculté des Sciences et Techniques,Université d'Abomey-Calavi, Cotonou (S. Badou); Institut de Géographie, d'Aménagement du Territoire et d'Environnement, Université d'Abomey-Calavi, Cotonou (H.-J. Dossou). DOI: https://doi.org/10.3201/eid2710.210268

Emerging Infectious Diseases • www.cdc.gov/eid • Vol. 27, No. 10, October 2021

Acknowledgments

We thank the Autonomous Port of Cotonou authorities and staff who facilitated our access to their infrastructures. We also thank Serenal Dool for editing the manuscript. This study was supported by the VEO European Union's Horizon 2020 (grant no. 874735), Academy of Finland (grant no. 318726), and the Jane and Aatos Erkko Foundation. Field and laboratory work was conducted under the research agreement between the Republic of Benin and the French Institute of Research for Sustainable Development (September 30, 2010) and the Scientific and Technical Cooperation Agreement (July 3, 2019) between the University of Abomey-Calavi and the Institut de Recherche et de Développement. The collection and use of tissue samples (spleen, lung, kidney, brain, dried blood spots) and genetic data was authorized under the Access and Benefit Sharing, Nagoya protocol (permit no. 608/DGEFC/DCPRNF/ PF-APA/SA) (December 2019). This study is part of a long-term partnership between Cotonou Autonomous Seaport, Port of Antwerp International, the Polytechnic School of Abomey-Calavi, the French Institute of Research for Sustainable Development, and the International Institute of Tropical Agriculture on maritime trade-associated biological invasions.

About the Author

Dr. Castel is a researcher in virology at the French National Research Institute for Agriculture, Food and Environment. His research interests include the diversity of hantaviruses and the evolutionary processes that shape it using phylogenetic and phylogeographic approaches.