Fernando Igoa

Emile Defoy

Daniel Janisch

Gwenaëlle Rousse

Pierre-Olivier Autran

Anissa Ghoridi

Amandine Séné

Marzena Baron

Leopoldo Suescun

Yann Le Godec

David Portehault
email: david.portehault@sorbonne-universite.fr

Revealing the elusive structure and the reactivity of iron boride α-FeB

Crystal structures can strongly deviate from bulk states when confined into nano-domains. These deviations may deeply affect properties and reactivity, and then call for a close examination. In this work, we address the case where extended crystal defects spread through a whole solid and then yield an aperiodic structure and specific reactivity. We focus on iron boride α-FeB, whose structure has not been elucidated yet, thus hindering the understanding of its properties. We synthesize the two known phases α-FeB and β-FeB in molten salts at 600 and 1100 °C, respectively. The experimental X-ray diffraction (XRD) data cannot be satisfyingly accounted for by a periodic crystal structure. We then model the compound as a stochastic assembly of layers of two structure-types. Refinement of the powder XRD pattern by considering the explicit scattering interference of the different layers allows quantitative evaluation of the size of these domains and of the stacking faults between them. We therefore demonstrate that α-FeB is an intergrowth of nanometer-thick slabs of two structure types:

β-FeB and CrB-type structures, in similar proportions. We finally discuss the implications of this novel structure on the reactivity of the material and its ability to perform insertion reactions, by comparing the reactivities of α-FeB and β-FeB as reagents in the synthesis of a model layered material: Fe 2 AlB 2 . By using synchrotron-based in situ X-ray diffraction, we elucidate the mechanisms of formation of Fe 2 AlB 2 . We highlight the higher reactivity of the intergrowth α-FeB, in agreement with structural relationships.

Introduction

Nanoscale solids can yield large deviations of crystal structures compared to bulk states: as a result of the interface energy contribution to the overall energy landscape, nano-domains can accommodate larger strains, [START_REF] Gilbert | Nanoparticles: Strained and Stiff[END_REF] deviate from bulk compositions, [START_REF] Delacroix | Electron Precise Sodium Carbaboride Nanocrystals from Molten Salts: Single Sources to Boron Carbides[END_REF] and even show size-dependent polymorphism. [START_REF] Christiansen | Size Induced Structural Changes in Molybdenum Oxide Nanoparticles[END_REF] The latter has been known for decades in the case of oxides, where ϵ-Fe 2 O 3 , 4 λ-Ti 3 O 5 , [START_REF] Ohkoshi | Synthesis of a Metal Oxide with a Room-Temperature Photoreversible Phase Transition[END_REF] bixbyite-V 2 O 3 , [START_REF] Bergerud | Synthesis and Phase Stability of Metastable Bixbyite V2O3 Colloidal Nanocrystals[END_REF] and more recently ϵ'-Li x V 2 O 5 [START_REF] Luo | Effect of Crystallite Geometries on Electrochemical Performance of Porous Intercalation Electrodes by Multiscale Operando Investigation[END_REF] could be isolated as a consequence of their low surface energy. These materials show unique properties for magnetic storage and shielding, [START_REF] Tuček | ε-Fe2O3: An Advanced Nanomaterial Exhibiting Giant Coercive Field, Millimeter-Wave Ferromagnetic Resonance, and Magnetoelectric Coupling[END_REF] light-driven information storage, [START_REF] Ohkoshi | Synthesis of a Metal Oxide with a Room-Temperature Photoreversible Phase Transition[END_REF] oxygen storage, [START_REF] Bergerud | Synthesis and Phase Stability of Metastable Bixbyite V2O3 Colloidal Nanocrystals[END_REF] and Li-ion batteries, [START_REF] Luo | Effect of Crystallite Geometries on Electrochemical Performance of Porous Intercalation Electrodes by Multiscale Operando Investigation[END_REF] respectively. Fewer examples pertain to non-oxides, with the notable cases of cubic Pb 2-x Sn x S 2 [START_REF] Soriano | No Title[END_REF] and of Pb m Sb 2n Te m+3n nanocrystals that enhance thermoelectric energy conversion. [START_REF] Soriano | Nanoscale Stabilization of New Phases in the PbTe-Sb2Te3 System: PbmSb2nTem+3n Nanocrystals[END_REF] CdSe nanoparticles also exhibit structure, morphology and defectdependent optical properties. [START_REF] Masadeh | Quantitative Size-Dependent Structure and Strain Determination of CdSe Nanoparticles Using Atomic Pair Distribution Function Analysis[END_REF] Especially, a high density of stacking faults explains the unusual white light emission of ultra-small CdSe nanoparticles. [START_REF] Yang | Confirmation of Disordered Structure of Ultrasmall CdSe Nanoparticles from X-Ray Atomic Pair Distribution Function Analysis[END_REF] The confinement of solids into nano-domains can then yield specific structures that require close examination to interpret unique properties.

Iron borides have been extensively investigated for structural [START_REF] Geng | A Data-Driven Machine Learning Approach to Predict the Hardenability Curve of Boron Steels and Assist Alloy Design[END_REF][START_REF] Li | Microstructural Evolution and Strain Hardening Mechanism of a Boron-Containing Metastable Austenitic Steel[END_REF] and functional [START_REF] Gupta | Metal Boride-Based Catalysts for Electrochemical Water-Splitting: A Review[END_REF][START_REF] Ni | Boron-Iron Nanochains for Selective Electrocatalytic Reduction of Nitrate[END_REF][START_REF] Bocarsly | Magnetostructural Coupling Drives Magnetocaloric Behavior: The Case of MnB versus FeB[END_REF] applications.

Beyond Fe-B alloying, multiple stoichiometric compounds are known, with compositions ranging from Fe 9 B to FeB 49 . 17 When manufactured at the nanoscale, iron borides exhibit distinct properties: FeB 2 nanoparticles are bifunctional electrocatalysts for overall water splitting surpassing the Pt/RuO 2 benchmark, [START_REF] Li | Earth-Abundant Iron Diboride (FeB2) Nanoparticles as Highly Active Bifunctional Electrocatalysts for Overall Water Splitting[END_REF] FeB nanoparticles are harder magnets than their bulk analogues. [START_REF] Rades | Wet-Chemical Synthesis of Nanoscale Iron Boride, XAFS Analysis and Crystallisation to α-FeB[END_REF][START_REF] Rades | Size and Crystallinity Dependence of Magnetism in Nanoscale Iron Boride, α-FeB[END_REF] Besides, the high temperature form of FeB, β-FeB, is also an efficient precursor for the synthesis of the MAB phase Fe 2 AlB 2 , [START_REF] Tan | Magnetocaloric Effect in AlFe2B2: Toward Magnetic Refrigerants from Earth-Abundant Elements[END_REF][START_REF] Kádas | Ni): A Group of Nanolaminated Materials[END_REF][START_REF] Levin | Effects of Al Content and Annealing on the Phases Formation, Lattice Parameters, and Magnetization of AlxFe2B2 (X=1.0,1.1,1.2) Alloys[END_REF][START_REF] Barua | Anisotropic Magnetocaloric Response in AlFe2B2[END_REF][START_REF] Cedervall | Magnetic Structure of the Magnetocaloric Compound AlFe2B2[END_REF][START_REF] Cedervall | Mössbauer Study of the Magnetocaloric Compound AlFe2B2[END_REF][START_REF] Cedervall | Magnetic and Mechanical Effects of Mn Substitutions in AlFe2B2[END_REF] which is a potential magnetic cooler material [START_REF] Tan | Magnetocaloric Effect in AlFe2B2: Toward Magnetic Refrigerants from Earth-Abundant Elements[END_REF] and a precursor to two-dimensional MBenes. 28-35 36 While the reactivity of β-FeB is well documented and rationalized at the atom-scale, [START_REF] Liu | Rapid Synthesis and Characterization of a Nanolaminated Fe2AlB2 Compound[END_REF][START_REF] Bennett | Magnetic and Magnetocaloric Properties of Fe2AlB2 Synthesized by Single-Step Reactive Hot Pressing[END_REF] its low temperature counterpart, α-FeB, [START_REF] Bocarsly | Magnetostructural Coupling Drives Magnetocaloric Behavior: The Case of MnB versus FeB[END_REF][START_REF] Rades | Wet-Chemical Synthesis of Nanoscale Iron Boride, XAFS Analysis and Crystallisation to α-FeB[END_REF][START_REF] Rades | Size and Crystallinity Dependence of Magnetism in Nanoscale Iron Boride, α-FeB[END_REF][START_REF] Tsurin | α-FeB: Structure, Hyperfine Interactions, and Magnetic Properties[END_REF][START_REF] Deyoung | A Mössbauer Effect Study of 57Fe in Transition Metal Monoborides[END_REF][START_REF] Jothi | A Simple, General Synthetic Route toward Nanoscale Transition Metal Borides[END_REF] remains elusive: its reactivity and properties have been seldom documented, [START_REF] Rades | Wet-Chemical Synthesis of Nanoscale Iron Boride, XAFS Analysis and Crystallisation to α-FeB[END_REF][START_REF] Rades | Size and Crystallinity Dependence of Magnetism in Nanoscale Iron Boride, α-FeB[END_REF] which is linked to the lack understanding of its atom-scale structure. [START_REF] Rades | Wet-Chemical Synthesis of Nanoscale Iron Boride, XAFS Analysis and Crystallisation to α-FeB[END_REF] There are strong hints that α-FeB is a modification at the nanoscale of the β-FeB structure displayed in Figure 1a-b. First, the X-ray diffractogram of α-FeB resembles that of β-FeB, with the exception of some missing diffraction lines and the presence of a broad diffuse scattering signal. [START_REF] Rades | Wet-Chemical Synthesis of Nanoscale Iron Boride, XAFS Analysis and Crystallisation to α-FeB[END_REF] Second, transmission electron microscopy images of α-FeB particles exhibit contrasted parallel stripes. [START_REF] Rades | Size and Crystallinity Dependence of Magnetism in Nanoscale Iron Boride, α-FeB[END_REF] The same streaks are observed only in nano-domains (30 -100 nm) of β-FeB. [START_REF] Rades | Wet-Chemical Synthesis of Nanoscale Iron Boride, XAFS Analysis and Crystallisation to α-FeB[END_REF][START_REF] Rades | Size and Crystallinity Dependence of Magnetism in Nanoscale Iron Boride, α-FeB[END_REF][START_REF] Tsurin | α-FeB: Structure, Hyperfine Interactions, and Magnetic Properties[END_REF][START_REF] Deyoung | A Mössbauer Effect Study of 57Fe in Transition Metal Monoborides[END_REF][START_REF] Jothi | A Simple, General Synthetic Route toward Nanoscale Transition Metal Borides[END_REF] Three hypotheses have been proposed to account for the elusive structure of α-FeB: [START_REF] Rades | Wet-Chemical Synthesis of Nanoscale Iron Boride, XAFS Analysis and Crystallisation to α-FeB[END_REF][START_REF] Rades | Size and Crystallinity Dependence of Magnetism in Nanoscale Iron Boride, α-FeB[END_REF][START_REF] Tsurin | α-FeB: Structure, Hyperfine Interactions, and Magnetic Properties[END_REF][START_REF] Deyoung | A Mössbauer Effect Study of 57Fe in Transition Metal Monoborides[END_REF][START_REF] Jothi | A Simple, General Synthetic Route toward Nanoscale Transition Metal Borides[END_REF][START_REF] Dorofeev | Local Magnetic Structure of α-Iron Monoboride Produced by Mechanical Fragmentation[END_REF][START_REF] Yelsukov | Mössbauer Study of Solid State Reactions under Mechanical Grinding of the Fe2B and FeB Borides[END_REF][START_REF] Kanaizuka | Interpretation of Mössbauer Spectra of β-FeB and Its Low Temperature Modification, α-FeB[END_REF][START_REF] Mohn | The Calculated Electronic and Magnetic Properties of the Tetragonal Transition-Metal Semi-Borides[END_REF] (i) antisite mixing between Fe and B atoms; 46 (ii) crystallization in a higher symmetry space group (S.G.); 47 (iii) intergrowth through stacking faults between two structure-types. [START_REF] Kanaizuka | Phase Diagram of Pseudobinary CrB-MnB and MnB-FeB Systems: Crystal Structure of the Low-Temperature Modification of FeB[END_REF] Notwithstanding, none of these attempts has provided a satisfactory solution able to reproduce experimental diffractograms.

In the present work, we study an α-FeB material synthesized in molten salts. We analyze the different hypotheses of structural order by means of conventional powder diffraction modelling and refinement. The planar defects proposed in hypothesis (iii) above cannot be averaged in a long-range periodic structure, and thus cannot be accounted for by classical diffraction analysis techniques. We then study powder X-ray diffraction patterns through the elaboration and refinement of a statistical layered model to account for the occurrence of stacking faults. [START_REF] Casas-Cabanas | a New Program for Refinement of Powder Diffraction Patterns from Layered Structures[END_REF][START_REF] Casas-Cabanas | A Program for Refinement of Structures with Extended Defects[END_REF] We show that this approach can accurately and quantitatively describe disorder and bidimensional defects. We then reveal the atomscale and nano-scale structure of α-FeB and demonstrate its potential as a precursor for the synthesis of model layered materials like Fe 2 AlB 2 .

Experimental section

Materials

All reagents were stored and handled under an argon-filled glove box. NaBH 4 (98 % min., Alfa Aesar), NaAlCl 4 (99.9%, Alfa Aesar), and FeCl 2 (99.5%, Alfa Aesar) were used as B, Al and Fe sources, respectively. NaH (90 %, Aldrich) was used as reducing agent. All reagents were used as received. A eutectic mixture of LiCl (99.9% Alfa Aesar)/KCl (99.9% Sigma Aldrich) (45:55 wt. %) was used as solvent.

Deionized H 2 O or MeOH (VWR Normapur grade) were used for washing steps.

Methods

The heat treatments were performed in vertical tubular ovens connected to temperature controllers from Eraly®. These ovens were calibrated every 100 °C between 100 and 1100 °C. The temperatures we indicate are the effective temperatures in the synthesis crucible.

Synthesis of α-FeB and β-FeB 850 mg (6.71 mmol) of FeCl 2 and 254 mg (6.71 mmol) of NaBH 4 were mixed with 5.1 g of the eutectic LiCl/KCl mixture. All solids were finely ground together at 20 Hz for 2 min in a Retsch MM400 ballmiller (airtight vials of 50 mL, one steel ball of 62.3 g and a diameter of 23 mm). The solid mixture was later placed in a Mo crucible (Ø15 × H200 mm), which was placed into a sealed quartz tube. The quartz tube was connected to a Schlenk line under Ar flow and placed inside a vertical oven. A heating ramp of 10°C•min -1 was used up to 600 °C for α-FeB and 1100 °C for β-FeB, followed by a dwell time of 3 h. We also tested different dwell times from 0.5 h to 12 h, and did not observe a significant effect of this parameter. After dwelling, the sample was let to cool down naturally and subsequently washed with methanol by sequences of redispersion, sonication (1 min) and magnetic separation of the solid, until the conductivity of the supernatant dropped below 50 μS•cm -1 . Methanol-washed samples required preliminary grinding of the solidified mixture in order to extract the product due to the low solubility of the salt in the solvent. The powders were dried under vacuum at room temperature until the pressure dropped below 1.0•10 -3 mbar and were then heated to 150 °C with an oil bath still under vacuum overnight.

Synthesis of Fe 2 AlB 2

Impregnation of FeB. In order to optimize the reactivity of FeB particles, they were pre-dispersed in the eutectic salt mixture before melting. To do so, 62 mg (0.925 mmol) of FeB (either the α or β modifications) were placed in a Schlenk tube and connected to a Schlenk line. 1 mL of methanol was added and the suspension sonicated for 20 min to assure proper dispersion. Then, 2.55 g of previously ground LiCl/KCl were added abruptly. The mixture was then evacuated under vacuum, first at room temperature and then in a stepwise heating to 50, 75, 100, 120 and finally 150 °C. Each heating step was maintained until the vacuum pressure dropped below 10 -2 mbar, except the final step (150 °C) that was maintained until the vacuum pressure dropped below 1.0•10 -3 mbar.

Final mixing and synthesis. Impregnated FeB and LiCl/KCl were mixed with of 266 mg NaAlCl 4

(1.388 mmol), 111 mg of NaH (4.625 mmol) and 2.55 g more of LiCl/KCl eutectic mixture. All solids were finely ground together at 20 Hz for 2 min in a Retsch MM400 ballmiller (airtight vials of 50 mL, one steel ball of 62.3 g and a diameter of 23 mm). The solid mixture was later placed in a Mo crucible (Ø15 × H200 mm), which was placed into a quartz tube. The quartz tube was connected to a Schlenk line under Ar flow and placed inside a vertical oven. A heating ramp of 10°C•min -1 was used up to 1000 °C, followed by a dwell time of 1 h. After dwelling, the sample was let to cool down naturally and subsequently washed in sequences of solvent mixing, sonication (1 min), centrifugation at 21000 RPM for 2 min and separation of the solid, until the conductivity of the supernatant dropped below 50 μS•cm -1 . The powder was dried under vacuum at 60 °C overnight.

Characterization

Powder X-ray diffraction patterns were collected in a Bruker D8 Discover diffractometer, with a thetatheta Bragg-Brentano geometry and Ni-filtered CuK α sealed-tube radiation source (λ average = 1.5418 Å) operating at 40kV/30mA. The diffractometer radius is 280 mm.The X-ray intensity was measured by 9.2 s collection steps of Δ2θ = 0.0128°. Diffracted intensity was collected using a LynxEye linear detector. The reference crystallographic structure of β-FeB and CrB structures were taken from the ICSD database (391329 and 44249 cards respectively), while Al(OH) 3 , Al 9.75 Fe 3 , Fe 3 Al and Fe 2 AlB 2 references were taken from the PDF-4+ database (00-020-0011, 01-082-9854, 00-050-0955 and 04-007-5354 cards, respectively). Rietveld refinements were performed using the GSAS-II software. [START_REF] Toby | GSAS-II: The Genesis of a Modern Open-Source All Purpose Crystallography Software Package[END_REF] Stacking faults modelling and refining was performed using the FAULTS software. [START_REF] Casas-Cabanas | a New Program for Refinement of Powder Diffraction Patterns from Layered Structures[END_REF][START_REF] Casas-Cabanas | A Program for Refinement of Structures with Extended Defects[END_REF] Transmission electron microscopy analyzes were performed in a Tecnai Spirit 2 apparatus, equipped with a LaB 6 electron gun, operating at 120 kV. Samples were dispersed by sonication on ethanol and deposited on top of Cu grids covered with a carbon membrane.

In situ angular X-ray diffraction was carried out at the ESRF synchrotron, on the ID11 beamline, in transmission mode at a fixed energy (wavelength) of 93.3159 keV (0.1329 Å). A Double Crystal Monochromator equipped with two Si(111) crystals was used to tune the energy and the beam size was set to 300x300 microns. The detector was a Frelon4M 2048x2048 16Bit fast CCD Camera. We used a homemade capillary oven. In brief, it enables performing transmission X-ray diffraction in situ during controlled heating of the reaction mixture in a 1.5 mm-diameter quartz capillary opened under argon flow, to mimic the conditions of ex situ lab synthesis.

Results and discussion

Synthesis of α-FeB and β-FeB FeB was synthesized by adapting a previously reported synthesis in molten salts: [START_REF] Portehault | A General Solution Route toward Metal Boride Nanocrystals[END_REF][START_REF] Wei | Synthesis and Properties of Fe-B Powders by Molten Salt Method[END_REF] iron chloride acts as metal precursor and NaBH 4 as both reducing agent and boron source, according to Equation 1. The eutectic LiCl/KCl mixture was chosen as a cheap, polar and low-melting point solvent (347°C) to trigger reactions in the liquid state. The powders were mixed at room temperature and heated under argon flow. After cooling, the powders were washed with methanol. X-ray diffraction patterns (Figure 2) indicate that α-and β-FeB are obtained for syntheses at 600 and 1100°C, respectively. Especially, we recognize the XRD pattern of α-FeB from the typical absences of the (101) β , (201) β , (111) β , (211) β and (301) β reflections (shown in red indexes in Figure 2, the β subindex in the crystallographic planes and directions' notation refers to the β-FeB structure, in contrast to the CrB subindex used later, referring to the CrB-structure type) of the β-FeB pattern, as well as the broad diffuse scattering at 2θ (CuK α ) ≈ 41°. [START_REF] Rades | Wet-Chemical Synthesis of Nanoscale Iron Boride, XAFS Analysis and Crystallisation to α-FeB[END_REF] No crystalline oxidized species could be detected by XRD (Figure 1a) in the corresponding products, which contain mostly α/β-FeB as crystalline phases. No amorphous oxides were observed by TEM (Figure 3a-b), but the β-FeB sample contains traces of Fe 2 B impurities. The characteristic parallel stripes of α-FeB are recognized in TEM images (Figure 3). where both Fe and B occupy the 4c Wyckoff position (Figure 1). [START_REF] Kapfenberger | Structure Refinements of Iron Borides Fe2B and FeB[END_REF] Its most prominent structural features are the BFe 6 trigonal prismatic units (Figure 1a inset), which share faces. These units form boron chains in trans configuration, as commonly observed for many other B-rich compounds. [START_REF] Gouget | New Route toward Nanosized Crystalline Metal Borides with Tuneable Stoichiometry and Variable Morphologies[END_REF][START_REF] Hirt | Synthesis and Crystal Structures of the New Ternary Borides Fe3Al2B2 and Ru9Al3B8 and the Confirmation of Ru4Al3B2 and Ru9Al5B8-x (x ≈ 2)[END_REF] The ensemble of the B-chains and its coordinated Fe atoms (named trigonal chains hereafter) extend along the [010] β direction (Figure 1a). The trigonal chains are packed along the [100] β direction by sharing edges (Figure 1b). All units are tilted along the [010] β direction by alternating angles of ± 21.2° versus the a axis (Figure 1b). On the other side, the structure of α-FeB is believed to be related to β-FeB due to the similarities of the powder patterns. To decipher its actual crystal structure, we analyze below the hypothesis raised in previous works.

Structure analysis of α-FeB

Hypothesis (i): Fe/B antisite mixing

Fruchart proposed partial anti site mixing of Fe and B, resulting in mixed occupations of the Fe and B positions. [START_REF] Fruchart | Cementites of Boron, Iron Borides, and the Nickel Boride, Ni3B[END_REF] This hypothesis is likely unreliable as the Fe-Fe distance would drop below that of elemental iron (2.15 Å against 2.84 Å respectively) and boron covalent chains would be broken, [START_REF] Rades | Wet-Chemical Synthesis of Nanoscale Iron Boride, XAFS Analysis and Crystallisation to α-FeB[END_REF][START_REF] Deyoung | A Mössbauer Effect Study of 57Fe in Transition Metal Monoborides[END_REF][START_REF] Kanaizuka | Phase Diagram of Pseudobinary CrB-MnB and MnB-FeB Systems: Crystal Structure of the Low-Temperature Modification of FeB[END_REF] which would impact the Curie temperature, contrary to observations. [START_REF] Kanaizuka | Phase Diagram of Pseudobinary CrB-MnB and MnB-FeB Systems: Crystal Structure of the Low-Temperature Modification of FeB[END_REF] Besides, models of Fe/B antisite mixing (Fe x B 1-x )_4c (B x Fe 1-x )_4c every 10 % exchange (Figure S1a) do not enable reproducing the experimental diffractogram. Rietveld refinement by also allowing asymmetric substitutions and the presence of vacancies (Figure S1b and Table 1) could not reach satisfactory agreement with the experimental pattern. All in all, we discard the hypothesis of Fe/B substitution. 1c-d. [START_REF] Barinov | Structure and Magnetic Properties of the α-FeB Phase Obtained by Mechanical Working[END_REF] This model is indeed closely related to the β-FeB structure and made of the same building units: BFe 6 trigonal prismatic units. In the CrB-type structure, the trigonal chains extend along the [001] CrB direction and exhibit a face-sharing planar packing in the [100] CrB direction, the layers are then piled along the [010] CrB direction.

For their analysis, Barinov et al. relied on a highly strained sample characterized by the large width of the diffraction peaks. [START_REF] Barinov | Structure and Magnetic Properties of the α-FeB Phase Obtained by Mechanical Working[END_REF] Our synthesis procedure yields samples with higher crystallinity. The thinner diffraction peaks offer better resolved experimental data to test hypothesis (ii). Rietveld refinement using the cell previously proposed structure as starting model does not yield a satisfactory fit, neither did manual increasing nor decreasing of the lattice parameters. The best fit, yielding lattice parameters a = 2.9460(4) Å, b = 7.6100(3) Å, c = 2.9491(2) Å, is shown in Figure S2a and the retrieved crystallographic parameters summarized in Table 1.

By extending the proposition of Barinov et al., we raised the question whether another space group could explain the reflection absences. For this reason, we have attempted to determine the unit cell and space group of α-FeB ab initio with the EXPO2014 software. [START_REF] Altomare | EXPO Software for Solving Crystal Structures by Powder Diffraction Data: Methods and Application[END_REF] The most plausible model found was an orthorhombic cell of space group Pbcn. For direct comparison to the β-FeB structure, this space group is treated under the unconventional Pnca setting, with cell parameters: a = 5.498( 4 Hypothesis (iii): Faulted intergrowth of the β-FeB and the CrB-type structures.

Kanaizuka suggested a recursive intergrowth between the CrB-type and β-FeB structures to model α-FeB. [START_REF] Kanaizuka | Phase Diagram of Pseudobinary CrB-MnB and MnB-FeB Systems: Crystal Structure of the Low-Temperature Modification of FeB[END_REF] The author relied on the earlier discoveries by Parthé and collaborators [START_REF] Hohnke | FeB and CrB Type Structures of Monosflieides and Germanides[END_REF][START_REF] Klepp | Yttrium-Nickel YNi with the FeB Structure Type[END_REF][START_REF] Klepp | Orthorhombic Gadolinium Nickel Silicide Gd3NiSi 2 with a Filled-up Hf3P2 Structure Type[END_REF][START_REF] Roy | Structures of the Rare-Earth-Platinum Compounds R7Pt3, R2Pt, RsPt 3 and RPt[END_REF][START_REF] Parthé | The CrB and Related Structure Types Interpreted by Periodic Unit-Cell Twinning of Close-Packed Structures[END_REF] that (a) several rare earth-metalloid alloys (DySi, HoSi, ErSi and PrGe) crystallized in the CrB structure at low temperature, which transitioned to the β-FeB structure upon heating, and that (b) β-FeB and CrB-type structures were intermixed in periodically faulted stacks in YNi and GdNi-DyNi. On these bases, Kanaizuka built a 4-layers model, whose binary intermixing renders either pure CrB-type, pure β-FeB or a mixed β-FeB/CrB structure (Figure 4). [START_REF] Kanaizuka | Phase Diagram of Pseudobinary CrB-MnB and MnB-FeB Systems: Crystal Structure of the Low-Temperature Modification of FeB[END_REF] The layers A-D (Figure 4a) are made from the unit cell of the original β-FeB structure with origin shifts in the ab plane such that they fit one on top of the other while keeping bond distances reasonable. Conventional Rietveld refinement cannot account for stochastic stacking faults, as it only considers periodic structures. However, following the original idea of Kanaizuka, the powder diffractogram can be simulated and refined based on the explicit interference from the scattering of each layer. The FAULTS software has been developed for this purpose [START_REF] Casas-Cabanas | a New Program for Refinement of Powder Diffraction Patterns from Layered Structures[END_REF][START_REF] Casas-Cabanas | A Program for Refinement of Structures with Extended Defects[END_REF] with high reliability for several types of materials. 64-67 FAULTS models the diffractogram by considering a probabilistic assembly of layers along defined stacking vectors. We define the structure from A, B, C and D layers as building blocks stacked along the [100] β vector, with probabilities α ji to transition from a layer j to i (where j and i are A, B, C or D). The probabilities of transition relate to the probabilities of existence P i through Equations 2 and 3 below. [START_REF] Treacy | How DIFFaX Works[END_REF] All symmetry elements but the identity are dropped as periodicity is no longer mandatory.

The atomic coordinates are listed in Table S1. As in Rietveld refinement, FAULTS can also account for size broadening, occupations, Debye-Waller factors, atomic positions and unit cell variations, among others. This approach allows to describe bidimensional defects and disorder. Since the transition probabilities can be refined, we can then quantify the amount of each layer.

𝑃𝑃 𝑖𝑖 = � 𝑃𝑃 𝑗𝑗 𝛼𝛼 𝑗𝑗𝑖𝑖 𝑁𝑁 𝑗𝑗=1 (2)
� 𝑃𝑃 𝑖𝑖 = 1 𝑁𝑁 𝑖𝑖=1 (3)
For a first set of simulations and for the sake of simplicity, only two sets of binary combinations were considered, that is A/B and B/C, as they suffice to explore all binary outcomes (see Figure 4b).

Transition probabilities were initially fixed to be equal for every arrival layer -α Ai = α Bi or α Bi = α Ci for i = A or B and B or C, respectively-and thus equal to the existence probability (α ji = P i ). Increments of 0.1 in existence probabilities were probed. The results can be seen in Figure S3a for the A/B mixing system for comparison purposes (Figure S3c), where an analogous trend was observed. These first simulations were crucial to fix a plausible initial input model for the following refinements. The complete values of probabilities employed can be found in Table S2.

In order to reproduce the diffuse scattering and reach an overall satisfactory fit of the experimental pattern, we had to include all four layers and asymmetrical transition probabilities (α ji ≠ α ij ) throughout the FAULTS refinement. The final refinement is depicted in Figure 5a for a restricted 2θ range for clarity. And the crystallographic parameters summarized in Table 1. The recovered transition probabilities are summarized in the Α ji tensor in Table S3, and a representation of the complete 2θ range can be found in Figure S4. The powder diffractogram is very well reproduced, achieving figures of merits of R wp = 7.3 % and χ 2 = 18.70. Ergo, our results confirm that α-FeB is an intergrowth of CrBtype and β-FeB structures stacked faults along the [100] β axis. This quantitative refinement permits to calculate the existence probabilities for the four types of layers: P A = 0.20, P B = 0.33, P C = 0.33 and P D = 0.14. Noteworthy, these values highly deviate from those initially proposed by Kanaizuka. [START_REF] Kanaizuka | Phase Diagram of Pseudobinary CrB-MnB and MnB-FeB Systems: Crystal Structure of the Low-Temperature Modification of FeB[END_REF] Rietveld refinement with a phase mixture of the CrB and β-FeB as model was also performed for comparison (Figure 5b). The difference in the modelling between Rietveld and FAULTS is evident and shows that α-FeB is an intimate intergrowth of the two structure types and not simply a phase mixture. Then the β-FeB cluster fraction is evaluated as the sum of the existence probabilities of pure β-FeB combinations plus half of the existence probabilities of β-FeB/CrB combinations. The CrB cluster fraction is calculated analogously. In the end, we evaluate β-FeB and CrB-type fractions at 0.57 and 0.43, respectively.

We further extended the analysis to assess the distribution of β-FeB and CrB domain sizes along the stacking direction. To do so, we adapted a model reported by Serrano-Sevillano and collaborators. 64 We generated a stochastic sequence of 100,000 layers following the transition probabilities α ji obtained from the FAULTS refinement, in order to obtain statistically meaningful values of average domain size. For enhanced accuracy, the calculation is repeated for 20 sequences generated independently. We obtained average domain sizes of 2.5 and 2.1 nm for the β-FeB and CrB domains, respectively. These values are consistent with the thickness of the contrasted domains observed by TEM (Figure 3).

All in all, from the different models studied to explain the features of the α-FeB diffraction pattern (Table 1), the nanoscale intergrowth of α-FeB and CrB-type structures stands out for its consistency with experimental results.

α-FeB: new reagent towards Fe 2 AlB 2

The lack of insights into its structure has hindered the study of the reactivity of α-FeB, while β-FeB is already known as precursor towards Fe 2 AlB 2 (Figures 1a-b and 6b), [START_REF] Liu | Rapid Synthesis and Characterization of a Nanolaminated Fe2AlB2 Compound[END_REF][START_REF] Bennett | Magnetic and Magnetocaloric Properties of Fe2AlB2 Synthesized by Single-Step Reactive Hot Pressing[END_REF]

Figure 1 .

 1 Figure 1. (a) Structure of β-FeB projected along the ab plane, main structural features are highlighted: BFe 6 trigonal prisms, B-chains and trigonal chains. The latter are the linear rectangular face-sharing arrangement of the BFe 6 trigonal prisms. (b) Structure of β-FeB projected along the ac plane showing the edge-sharing piling of the trigonal chain. (c) Structure of FeB in the CrB-type structure, suggested by Barinov et al., 47 projected along the (c) bc and (d) ab planes. BFe 6 trigonal prisms in the CrB structure-type share both the rectangular and triangular faces to form bidimensional Fe 2 B 2 sheets.

FeCl 2 ( 1 )Figure 2 .

 212 Figure 2. Powder XRD diffractograms of α-and β-FeB products synthesized in molten salts at 600 and 1100 °C, respectively, and washed with methanol. The expected Bragg positions for the β-FeB structure are marked with green stripes. Miller indexes marked in red correspond to the characteristic absences of β-FeB reflections in α-FeB. Inset: Magnification of the diffraction patterns in the framed region, characteristic absences in α-FeB are marked with red dots, Fe 2 B traces are marked with *.

Figure 3 .

 3 Figure 3. TEM micrographs of α-FeB. Inset in (a) shows the corresponding size distribution histogram.(b) A higher magnification of a particle highlights streaks with varying contrasts and an average thickness of ~ 2 nm. These features are characteristic of the α-FeB structure.[START_REF] Rades | Size and Crystallinity Dependence of Magnetism in Nanoscale Iron Boride, α-FeB[END_REF]

) Å, b = 2 .

 2 951(18) Å, c = 4.062(2) Å. This model is very close to the reported β-FeB structure (S.G. Pnma, a = 5.495 Å, b = 2.941 Å, c = 4.057 Å), with the main difference being a c glide plane substituting the m mirror perpendicular to the b axis. The glide plane imposes the extra systematic absence of h0l: l = 2n, which in turn forbids the (101), (201) and (301) reflections. To further account for the remaining experimental absences, we attempted to solve the structure in this space group by direct methods, but could not reach any plausible result. Alternatively, we tested the assumption that the atoms would lie in the same positions than in β-FeB given the similarity of the cell metrics and the fact that α-FeB transitions to β-FeB under thermal treatment. In turn, placing Fe and B atoms in the original β-FeB positions does extinguish the remaining (111) β and (211) β peaks. However, that is also true for the experimentally observed (011) β reflection, which discards this atomic positioning. Refinement in general positions does not yield a plausible model nor does it reproduce the diffraction pattern (best fit shown in Figure S2b with crystallographic parameters in

  Figure 4b indicates the resulting structures from the binary mixtures of these layers stacked along the [100] β axis. Kanaizuka 48 then computed the interferences of the scattered waves from each layer by considering their probabilities of existence P A = 0.55 and P B = P C = P D = 0.15, which could account for the change of relative intensities of the (101) β , (201) β , (111) β , (211) βand (301) β reflections and for the two distinct crystallographic Fe positions identified by[START_REF] Altomare | EXPO Software for Solving Crystal Structures by Powder Diffraction Data: Methods and Application[END_REF] Fe-Mössbauer spectra in α-FeB. However, the calculated diffractogram was still far from reproducing the experimental profile.[START_REF] Kanaizuka | Phase Diagram of Pseudobinary CrB-MnB and MnB-FeB Systems: Crystal Structure of the Low-Temperature Modification of FeB[END_REF][START_REF] Rades | Synthese Und Charakterisierung von Nanopartikeln Im System Eisen-Bor[END_REF]

Figure 4 .

 4 Figure 4. (a) The four possible layer-types according to Kanaizuka's model. (b) Possible binary combinations outcome. (c) Visualization of binary outcomes.[START_REF] Kanaizuka | Phase Diagram of Pseudobinary CrB-MnB and MnB-FeB Systems: Crystal Structure of the Low-Temperature Modification of FeB[END_REF]

5 (

 5 and S3b for the A/B and B/C binary combinations, respectively. One can readily observe that some combinations significantly reduce the intensities of the reflections missing in α-FeB. The case where P B = P C = 0.marked in red in FigureS3b) stands out as all the (101) β , (201) β , (111) β , (211) β and (301) β reflections vanish, in agreement with the experimental diagram. This strongly supports the hypothesis of CrBtype/β-FeB intergrowth by stacking faults along the [100] β axis, although the diffuse scattering is still not well modelled. Further simulations (FigureS3d) were carried out by fixing the existence probability P B = P C = 0.5 and by varying the transition probabilities α ji , under the simplification hypothesis that transition probabilities are symmetric (α ji = α ij ). The main effect observed is the modification of the diffuse scattering profile with α BB , which at the limit of α BB = α CC = 0 and α BC = α CB = 1 transforms in the characteristic Bragg peaks of the CrB-type structure. Similar α ji screening simulations were carried out

Figure 5 .

 5 Figure 5. Different refinement approaches to model α-FeB. (a) FAULTS refinement with the 4-layers model proposed by Kanaizuka. 69 (b) Rietveld refinement of a phase mixture β-FeB/CrB-type structure.

  our structure resolution shows that α-FeB is an intergrowth made of nano-domains of β-FeB and CrB-type structures. The CrB-type component of this intergrowth presents a clear structural relationship with Fe 2 AlB 2 (Figure 6), which could ensure insertion of Al atoms to form Fe 2 AlB 2 , and thus easier conversion to this MAB phase. Overall, comparing the reactivity of α-FeB and β-FeB could highlight interesting effects of nanoscale domains and of crystal defects. We have then evaluated the possibility to use α-FeB in the synthesis of Fe 2 AlB 2 .

Figure 6 .Figure 7 .

 67 Figure 6. Structural relationship between the (a) CrB-type structure and (b) Fe 2 AlB 2 .

TOC

Table 1 .

 1 Summary of crystallographic parameters for the refinement of all the models studied for α-FeB. In the case of the phase mixture model, only the β-FeB structure parameters are shown. Pnma).[START_REF] Barinov | Structure and Magnetic Properties of the α-FeB Phase Obtained by Mechanical Working[END_REF] The Cmcm and the Pnma space groups are linked by a groupsubgroup relation, for which the axis in Cmcm relate to those of Pnma as: a Cmcm = b Pnma , b Cmcm = c Pnma and c Cmcm = a Pnma . This allows to compare the lattice parameters, which remain almost unaltered for b Pnma (2.9408 Å vs. 2.954 Å in a Cmcm ), while c Pnma expands from 4.0477 Å to 7.564 Å in b Cmcm , as in the latter it corresponds to the stacking direction of the trigonal layers. This expansion is compensated by the contraction of the a Pnma by approximately the same amount, which in turn yields similar cell volumes: V Pnma = 65.414 Å vs. V Cmcm = 65.982 Å. The increase in symmetry would then yield the CrBtype structure, shown in Figure

	Hypothesis:	(i)	(ii)		(iii)
	Model:	Anti-site mixing	S.G. Cmcm	S.G. Pnca	Phase mixture	Intergrowth
	a (Å)	5.4954(2)	2.9460(4)	5.4944(4)	5.49312(3)	5.4989(2)
	b (Å)	2.9492(3)	7.6100(3)	2.9491(2)	2.94817(4)	2.9505(3)
	c (Å)	4.0622(2)	2.9491(2)	4.0614(3)	4.06097(5)	4.0641(2)
	Volume (Å 3 )	65.961(8)	66.130(10)	65.808(8)	65.766(13)	65.961(8)
	R wp (%)	52.8	49.9	33.3	35.1	7.4
	χ 2	604.50	549.40	245.10	267.34	18.72

Hypothesis (ii): Crystallization in the CrB-type structure (S.G. Cmcm) Barinov et al. proposed that the higher symmetry space group Cmcm under the lattice parameters a = 2.954 Å, b = 7.564 Å, c = 2.953 Å was responsible for imposing systematic absences in α-FeB compared to the ideal β-FeB structure (S.G.

Table 1 )

 1 . No other proposed S.G. by EXPO2014 could account for the observed and missing reflections of α-FeB. Consequently, we

discarded hypothesis (ii).

Acknowledgements

The authors acknowledge the French Ministry for Research and the Doctoral School ED397 for funding of F.I. This work has been supported by the European Research Council (ERC) Consolidator Grant GENESIS under the European Union's Horizon 2020 research and innovation program (grant agreement n° 864850). In situ XRD was performed at ESRF under the proposal MA-5188 and received financial support from the ESRF.

excess ensures fast delivery of Al atoms, while the excess of Al versus FeB fastens its incorporation, as previously reported. [START_REF] Liu | Rapid Synthesis and Characterization of a Nanolaminated Fe2AlB2 Compound[END_REF][START_REF] Bennett | Magnetic and Magnetocaloric Properties of Fe2AlB2 Synthesized by Single-Step Reactive Hot Pressing[END_REF][START_REF] Dey | Synthesis of Nanolayered Ternary Borides Powders (MAB Phases) by Sustainable Molten Salt Shielded Synthesis/Sintering (MS3) Process[END_REF] In order to favor the Al incorporation, we have increased the surface of the FeB particles exposed to the reaction medium by developing an impregnation protocol. The particles were dispersed in methanol and then abruptly mixed with previously ground a LiCl/KCl eutectic salt mixture in order to avoid re-aggregation. If this step is by-passed, FeB is retrieved in the final product, up to a 20 % weight fraction.

Both α and β-FeB were reacted following the same protocol, by heating at 1000°C for 1 h. The XRD diagrams of the two resulting powders (Figure S5) exhibit reflections of Fe 2 AlB 2 as major phase.

However, α-FeB yielded a significantly higher purity of 61 wt. % versus only 43 wt. % when β-FeB was employed. This higher purity of the MAB phase can be linked to the insertion reaction that was expected from the structural similarity between the CrB-type component and the Fe 2 AlB 2 product. To confirm this point, we assessed whether α-FeB transitioned through the high temperature β phase before formation of Fe 2 AlB 2 . We employed synchrotron-based in situ X-ray diffraction to probe this reaction mechanism, by acquiring an XRD pattern every minute in transmission mode through the reaction medium with a composition identical to the lab-scale synthesis. The diagrams are plotted against the reaction time as a heatmap in shows an identical reaction sequence, at similar temperatures. This suggests that incorporation of Al atoms into the β-FeB nano-domains in α-FeB is a rate-limiting step for the overall reaction, which is consistent with a larger structural rearrangement required to transform into the Fe 2 AlB 2 structure. The difference of grain sizes could also be at the origin of the reactivity difference, although this is unlikely as all FeB is consumed in both cases, suggesting that Al diffusion does note limit the reaction.

Consequently, α-FeB could replace advantageously β-FeB as a precursor to produce Fe 2 AlB 2 , since it ensures higher Fe 2 AlB 2 purity and can be produced at 400 °C against 1100 °C for β-FeB, thus ensuring lower energy consumption.


Chenghao Xin

Zhelu Hu

Zhuoqun Fang

Mahima Chaudhary

Hengyang Xiang

Xiangzhen Xu

Lionel Aigouy

Zhuoying Chen
email: zhuoying.chen@espci.fr

Flexible and wearable plasmonic-enabled organic/inorganic hybrid photothermoelectric generators

Keywords: photothermoelectric conversion, organic/nanoparticle hybrids, light harvesting, plasmonic nanoparticles, photothermal effects

HAL is

Introduction

Thermoelectric (TE) devices based on solution-processed conjugated organic and organic-inorganic hybrid materials represent currently an emerging system for near-ambient energy harvesting applications (< 200 °C). [1][2][3][4][5][6] While many previous examples demonstrated only TE conversion, [7][8][9][10][11] new application needs have emerged recently requiring photothermoelectric conversions, 5,[12][13][14] aiming to perform light harvesting 12,14 and in some case positioning for device integration into multifunctional smart textile. 5 By comparison photovoltaic technologies, photothermoelectric generators can offer their unique functionality in both situations with light (e.g. by light-induced temperature difference) 5,12,15 or without light (e.g. by environment-induced temperature difference). 5,16,17 Concerning light harvesting, organic solar TE generators face similar challenges as inorganic TE counterparts to obtain sufficient temperature difference across the device under illumination. On classical inorganic TE modules, a few mitigation strategies have been proposed: For example, on Bi 2 Te 3 TE elements, commercial multilayer thin-film light absorbers have been applied to boost the temperature difference of the device under solar illumination. 18 Similar photothermal effect can be also well achieved by plasmonic nanoparticles: 19,20 A layer of metal nanoparticle/polymer hybrid beads, 21 or simply a layer of silver nanoparticles, 22 was recently applied on a commercial inorganic TE module. Nevertheless, on flexible organic and organicinorganic hybrid solar TE generators, there is hardly an attempt to apply such a plasmonic-induced photothermal concept. Indeed, in contrast to common solar concentration techniques involving the integration of mirrors to generate sufficient temperature difference across the device, 12,14 the photothermal concept gains advantages in the simplicity of device fabrication and installation favorable for the light-weight and flexible characteristics of the final devices. A few studies were reported recently applying nonplasmonic photothermal materials on organic/hybrid TE devices but exhibiting only modest performances: For example, applying a MoS 2 /polyurethane photothermal composites on the hot ends of a tellurium-doped PEDOT:PSS TE generator, He et al achieved a voltage output of 1.48 mV under sun light. 15 Recently, towards the application of smart textile, on polypropylene nonwoven fabrics, Zhang et al. applied a polypyrrole as the photothermal layer to fabricate flexible polypyrrole/PEDOT:Tos-based photothermoelectric generators capable to generate 0.677 mV and 0.0636 nW under sun light when worn on arm. 5 How to achieve flexible photothermoelectric generators with more significant voltage and power output remain currently a hot topic of research.

In this work, for the first time, carefully designed solution-processed colloidal plasmonic gold nanoparticles (Au NPs) are integrated onto organic-inorganic hybrid TE generators to achieve solutionprocessed and flexible photothermoelectric generators on flexible polyethylene terephthalate (PET) substrates and on polypropylene nonwoven fabrics for solar light harvesting without complex solar concentration techniques. Various batches of plasmonic photothermal Au NPs were synthesized by colloidal routes and a "visible NP ink" was formulated in order to absorb the majority of the visible portion of the solar spectrum. In order to maximize the device performance, we first investigated three different types of n-type flexible TE materials applicable on PET substrates and compared their performances when coupled with p-type PEDOT:PSS. The photothermal layer, the visible NP ink, was then deposited onto the hot ends of optimized hybrid TE generators to complete the final device structure. Under simulated AM 1.5G 1-sun solar illumination, we investigated the achieved temperature difference (T) measured between the hot and cold electrodes and the maximum voltage and power output generated by the final device by comparison to the control device without photothermal NPs.

Optimized photothermoelectric devices were further fabricated onto polypropylene nonwoven fabrics on which we explored the device characteristics under natural sun light as the device worn on arm.

Results and discussion

The Au NP photothermal layer

The morphologies of three batches of CTAB-capped Au NPs are shown in Figure 1 (a), (b), and (c), namely batch a, b, and c, respectively. The elongated morphology of these Au NPs can allow a facile tuning of their longitudinal localized surface plasmon resonance (LSPR) absorption. These three batches thus exhibit a distinctive solution optical absorption peak at  = 745 nm, 652 nm, and 554 nm, respectively (Figure 1(d) and Table S1 of supplementary information (SI)). In order to absorb the majority of the visible portion of solar spectrum ( from  = 500 nm to 850 nm), a visible Au NP ink was formulated by combining these three batches of NP solutions according to a certain proportion (described in the experimental section in SI). The solution-state optical absorption of this NP ink and the thin-film absorption when this ink was deposited and dried onto a PET substrate are shown as the green and purple curves in Figure 1(d). The major part of the visible-absorbing capability of this NP ink was still maintained after being deposited and dried onto PET. In contrast to the metallic nature of bulk gold, revealed by standard four-point probe measurements, a film of  2.5-µm-thickness formed CTABcapped Au NPs deposited onto a PET substrate is however resistive with an average conductivity of 1.5×10 -5 S/cm (Figure S2(a) in SI). This is likely due to the insulating NP ligands, CTAB, with a 16carbon alkyl chain in its chemical structure. Such an Au NP film was then applied below as the photothermal layer on the hot electrodes of the photothermoelectric generators, with its resistive nature preventing electrical shorting between electrodes after its deposition. S1 with an average aspect-ratio (AR = average length/average diameter) of 3.9 (a), 2.3 (b), and 1.9 (c), respectively, namely batch a, b, and c. (d) The optical extinction measured on these three NP solutions, the visible NP ink formed by a combination of these three batches, and the NP layer after the NP ink was deposited and dried onto a PET substrate.

The optimization of the hybrid TE devices

In order to optimize the flexible TE devices to be applied below for photothermoelectric generators, three different n-type materials was explored here to couple with the p-type PEDOT:PSS, which is currently still one of the state-of-the-art p-type printable TE materials compatible with flexible substrates. 4 The n-type materials investigated here, possible to be deposited onto flexible PET substrates and requiring no or only mild (⩽ 150 °C) heat treatments after deposition, are orthorhombic -Ag 2 Se, MoS 2 nanosheet powders, and commercial indium-doped tin oxide (ITO) films, respectively. Both -Ag 2 Se and ITO films have been recently reported to be promising n-type flexible TE materials exhibiting room-temperature power factors ranging from a few hundreds to 1700 µW m -1 K -2 for -Ag 2 Se 23 and a few tens of µW m -1 K -2 for ITO 24 (power factor PF = S 2 , where S is the Seebeck coefficient, is the electrical conductivity). MoS 2 was also shown to be an interesting TE materials with room-temperature n-type transport behavior. [25][26][27][28] But its physical property is highly dependent on the sample orientation (due to its anisotropic nature), 25 number of layers (or bulk), 26 doping conditions, 25 and the sample preparation techniques (e.g. exfoliated 29 or synthesized from chemistry routes 27,28 ).

The XRD spectrum and the microscopic morphology revealed by scanning electron microscopy (SEM) of the orthorhombic -Ag 2 Se film synthesized in this work are shown in Figure 2 The XRD and morphology characterizations of the MoS 2 synthesized in this work are shown in Figure S3 of SI. XRD confirmed the crystallinity of both materials, with the majority of the Ag 2 Se film corresponds well to the orthorhombic -Ag 2 Se phase (JCPDS no. 00-024-1041) together with the existence of a small amount of unreacted Se, which is not uncommon for -Ag 2 Se films without high temperature annealing 23 (not compatible to the flexible PET substrates). In order to measure the Seebeck coefficient, "uni-leg" structure devices were fabricated on PET substrates, where the hot ends of each stripe was connected to the cold ends by conductive silver paste. The Seebeck coefficient near room temperature (T  55 K), extracted from the slope of the linear fit of the Seebeck voltage (V) per leg generated under different steady temperature difference (T) between hot and cold ends, 7,8,30,31 was measured to be 19.8 µV K -1 for the p-type PEDOT:PSS stripes (Fig. 2(c)). With the same method, among the three n-type materials investigated, the -Ag 2 Se film exhibited a largest Seebeck coefficient reaching -129.2 µV K -1 (Fig. 2(c) and Table 1). By a 4-point probe method, the room-temperature conductivity () of the p-type PEDOT:PSS stripe and the three n-type materials were measured and listed in Table 1 (and Figure S4). With the measured Seebeck coefficient and the electrical conductivity, one can determine the room-temperature power factor of these p and n-type stripes, among which -Ag 2 Se exhibited a largest power factor of 802 µW m -1 K -1 , followed by PEDOT:PSS with a power factor of 43 µW m -1 K -1 . Various TE devices were then fabricated on PET substrates, formed by 6 pairs of p/n legs of PEDOT:PSS together with one of the three above-mentioned n-type materials. The measured TE characteristics are shown in Fig. 2(d) and Fig. S3(c). In agreement with the power factor evaluation, under an identical ∆T of 50 K, among the different devices fabricated, the most optimized TE device was found to be the one based on PEDOT:PSS/ -Ag 2 Se exhibiting a maximum power output of 310 nW.

Hybrid photothermoelectric generators on PET substrates

We then applied the visible Au NP ink together with the most optimized TE configuration found above, with a p/n structure of p-type PEDOT:PSS and n-type -Ag 2 Se stripes, to build photothermoelectric generators on flexible PET substrates. These devices, with and without NPs (the control condition), were then tested under a half-covered condition (Fig. 3(a)), by exposing their hot ends to simulated solar AM1.5G illumination operated at 1sun condition (100 mW cm -2 ) whereas the cold ends were kept in dark. Here, the simulated solar illumination thus served as an indirect heat source. After one minute of illumination, a stable (steadystate) T was generated between the hot and cold electrodes of the devices (Fig. S1). As shown in Fig. NPs) as more and more of them were deposited. A saturation of power output was found at four NP coatings, likely due to a trade-off between NP absorption and thermal transport over NP film thickness.

The device characteristics obtained on the optimized NP-decorated hybrid photothermoelectric generators are clearly much higher than those measured on the control device, where only a voltage output of 10 mV and a maximum power output of 11.3 nW were measured. The essential reason behind the 13-fold higher power output of the NP-coated device than the control device lies in the higher temperature difference (T) built up between the hot and cold ends of the devices under 1-Sun illumination thanks to the photothermal effect of the Au NPs: a steady-state T of 36.7 K was measured on the hybrid device with four NP coatings under 1-sun illumination, whereas a T of 11.8 K was measured on the control device under identical measurement conditions (Fig. S1).

The dynamic voltage output characteristics, where the 1-sun illumination was modulated ON and OFF, of the hybrid photothermoelectric generators with and without NPs are shown in Figure 3(e). The rise time rise of both devices, defined as the time necessary for the device to reach 80% of its maximum voltage output, was measured to be 17s and 10s, respectively for the NP-decorated device and the control device without NP. The slightly longer rise time measured on the NP-hybrid device by comparison to the control device is coherent with the additional processes involved in NP-hybrid devices in terms of photo-absorption, photothermal conversion, and the heat transfer processes from the photothermal NPs to the hot ends of the devices upon illumination.

The maximum power output of 146 nW measured on the current flexible NP-hybrid devices under 1sun illumination is remarkable compared to previous works: For example, previous demonstrated solar TE generators based on p-type single-wall CNTs and n-type polyethylenimine-doped single-wall CNTs generated a power output of 180 nW under 2-sun illumination. 14 While higher output power can be achieved by connected multiple modules in series or in parallel and/or applying sun light concentration lenses, previously reported single module of PEDOT:PSS-based solar TE generators, when tested under 1-sun illumination and a similar half-covered testing condition, generated only power outputs ranging from 11 nW to maximum 110 nW depending on the sample crystallinity and the doping conditions. 7 Recent solar TE generators based on p-type PEDOT:PSS and n-type Nickle metallic thin film achieved only 9.8 nW under 1-sun illumination. 32 The current hybrid photothermoelectric devices thus represent an interesting strategy to boost the power output without complex light concentration lenses. In addition to the device architecture shown in Fig. 3(a), we have also experimented another device configuration where the photothermal Au NPs incorporated directedly inside the TE elements (results and discussions in SI, Fig. S5 andS6). Despite the different optimizations performed, the maximum power generated under 1-sun illumination was found to be modest ( 14 nW) on these devices. These results nevertheless indicated that it is more advantageous to place the photothermal Au NPs on top of the hot-electrode side of the device (as in Fig. 3(a)) instead of incorporating the NPs inside the TE materials.

Wearable hybrid photothermoelectric generators on fabrics

Motivated by the growing interests on the application of flexible TE devices for wearable smart textile, based on the optimized device architecture achieved above (Fig. 3), we further fabricated functional and wearable photothermoelectric generators based on hybrid photothermal Au-NP-coated PEDOT:PSS/Ag 2 Se on polypropylene nonwoven fabrics (Figure 4). Under sunshine and ambient condition at 25 °C, a voltage output of 24 mV was measured on the device (Fig. 4(b)), while only 0.1 mV was measured on the same device without sunshine (indoors). By comparison to the device fabricated on PET substrates (Fig. 3), the slight lower voltage output of the current device fabricated on fabrics likely originates from the different surface properties between the two types of substrate and the resultant different NP distribution on substrate, together with the different light intensity received on device under the natural ambient sun light by comparison to the simulated 1-sun illumination.

Nevertheless, the voltage output (24 mV) measured on the current device when worn on arm and under natural sun illumination is significantly higher than previous reports on wearable photothermoelectric generators: For example, nanogenerators based on MoS 2 /PU photothermal film on PET substrates generated only a voltage output of 1.48 mV under outdoor sun light illumination. 15 Polypyrrole-enabled PEDOT:Tos-based photothermoelectric generators built on polypropylene nonwoven fabrics generated only 0.677 mV under natural sun light with device worn on arm. 5 There were also flexible and wearable TE devices based on p/n Sb 2 Te 3 /Bi 2 Te 3 fabricated on silk 33 and uni-leg-type MoS 2 -CNTs fabricated on cotton. 17 But no sun light harvest has been reported on these devices built on silk or cotton. Considering the on-going developments obtained on ultra-low power electronic circuits 34,[START_REF] Nebel | Low Power Design in Deep Submicron Electronics[END_REF] and the possibility to combine wearable electronics with the micro-electromechanical system (MEMS) technology, [START_REF] Ciuti | [END_REF]37 the current wearable photothermoelectric generators offer bright perspectives to harvest energy for multifunctional smart textiles.

Conclusions

In summary, in this work, through a series of synthesis and optimizations, we successfully fabricated Formulation of Au NP ink. The above-mentioned different batches of CTAB-capped Au NPs were then mixed in an optimized proportion, in order to absorb the visible portion of solar spectrum ( from  = 500 nm to 850 nm, referred below as "visible NP ink"): The batch a, b, and c of NP solutions were first diluted to a certain concentration (5.7 mg/mL for a, 5.3 mg/mL for b, and 2.2 mg/mL for c) and then mixed together according to the volume ratio a:b:c = 7%:18%:75% before applied.

Cysteamine-capped Au NPs. Plasmonic Au NPs passivated by short-chain ligands of cysteamine were synthesized according to previous reports. 3,4 In brief, 400 uL of cysteamine (213 mM) was mixed with 40 ml of HAuCl 4 (1.42 mM) solution under gentle stirring in dark. After that, 10 uL of NaBH 4 (10 mM) solution, freshly prepared in an ice-water bath, was added into the above solution under vigorous stirring for 10 min, resulting a color change of the solution from the brownish to clear wine red. Mild stirring was continued for another 30 minutes. Finally, the solution was stored in dark in the refrigerator (4 °C) before use.

Fabrication of hybrid photothermoelectric generators

Deposition of p-type PEDOT:PSS legs. 1.9 mL of aqueous PEDOT:PSS (PH1000, Ossila) was first doped by adding 0.1 mL dimethyl sulfoxide (DMSO) corresponding to 5 vol% of doping according to previous reports. 5 The p-type PEDOT:PSS legs were fabricated by drop-casting the doped PEDOT:PSS solution through a screen mask onto the unmasked areas of a clean substrate of PET or polypropylene nonwoven fabrics to form multiple 2 cm by 0.3 cm PEDOT:PSS strips, followed by drying on a hotplate at 130 °C for 15 minutes under ambient conditions and mask removal. These legs on PET were subsequently immersed in ethylene glycol (EG) for 120 minutes according to previous reports 5,6 to optimize the thermoelectric figure-of-merit (ZT), followed by a rinse of DI water and drying on a hotplate at 120ºC for 15 minutes. The EG treatment step was omitted for the legs deposited on polypropylene nonwoven fabrics.

Synthesis and deposition of n-type

Ag 2 Se legs. The Ag 2 Se ink was prepared according to Mallick et al. 7 Briefly, 1.22 g of Ag and 0.45 g of Se powders were mixed in 1.81 mL toluene together with 0.1 g of polystyrene. This was followed by an addition of a 2 wt% adhesive solution, formed by string 0.2 g of ethyl cellulose, 8 ml of terpineol, and 0.5 ml of titanium isopropoxide together until forming a homogenous solution. The ink was then stirred for at least 48 h at 1200 rpm at room temperature. To form n-type Ag 2 Se legs on substrates, the above-mentioned Ag 2 Se ink was printed on either the PET substrate or polypropylene nonwoven fabrics by doctor blading followed by drying on a hotplate at 150 °C for 10 minutes.

Synthesis and deposition of n-type MoS 2 legs. Molybdenum disulfide MoS 2 nanosheet powders

were synthesized by a modified hydrothermal method similar to that reported by Lun et al. 8 Briefly, 0.3 g of MoO 3 and 1.8 g of Na 2 S•9H 2 O were first mixed into 30 mL of DI water under magnetic stirring.

After they were totally dissolved, 2.8 ml of 0.6 M hydrochloric acid solution was then added into the mixture. The mixture was then transferred into a PTFE lined stainless steel autoclave and heated at 260 °C for 12 h. After 12 hours of heat treatment, the autoclave was allowed to cooled down to room temperature. The resultant black solids was collected, filtered, and washed with DI water. The solids were then dried under vacuum at 80°C for 24 h. In order to deposit MoS 2 legs on PET substrates, 0.1g of MoS 2 powder and 0.01g of Poly(vinyl alcohol) (PVA, average Mw = 13000 -23000 g/mol) were mixed into 1 ml of DI water. This solution was then used to deposit MoS 2 legs on PET substrates by the same doctor blade method as for Ag 2 Se legs.

Deposition of n-type

Indium-doped tin oxide (ITO) legs. ITO-coated PET substrates were obtained from Sigma Aldrich. 3M adhesive tapes were applied as a mask to cover certain area of the ITO-PET substrate in order to define multiple ITO strips (legs). On the un-covered area, the ITO layer was etched by diluted hydrochloric acid solution ( 3.4 wt%) and together with zinc powders, followed by DI-water rinsing and adhesive tape removal.

Fabrication of p/n-type photothermoelectric generators.

To obtain p/n-type photothermoelectric generators, p-type PEDOT:PSS legs were first deposited onto a PET substrate or a piece of polypropylene nonwoven fabric as described above. N-type Ag 2 Se legs were then deposited as described above onto the same substrate. These p and n legs were then connected together by conductive silver paint (RS components, 123-9911). After the silver electrodes were completely the abovementioned CTAB-capped Au NP visible ink was deposited onto the "hot ends" of the device (i.e. electrodes that will be exposed under illumination) followed by complete drying off the solvent (water).

This NP ink deposition was repeated multiple times (referred as numbers of NP coatings below). The "control" devices refer to the devices without any deposition of Au NPs. Finally, these legs were connected in series by conductive silver paint.

Fabrication of cysteamine-Au

Characterizations

The optical absorbance of different Au NP solutions and NP films on PET substrates were measured by a Jasco V770 UV-Visible/near-IR spectrometer. SEM characterizations were performed by a FEI Magellan 400 system with a standard field emission gun source. SEM observations were carried out either directly on a complete NP-decorated photothermoelectric generator built on a PET substrate, or on an Ag 2 Se film deposited on a silicon substrate. XRD spectra was measured by a PANalytical X'Pert X-ray diffractometer using Cu-K radiation. Glass substrates were used for XRD measurements. TEM samples were prepared by drop-casting the Au NP solution onto a carbon-coated copper TEM grid followed by solvent drying. TEM images were obtained by a JEOL 2010 TEM microscope operated at 200 kV. The electrical conductivity was measured at the room temperature (1 atm, 25-27 °C) by a standard four-point probe method (Ossila Four-Point Probe system with built-in source measurement unit). On each p-or n-type materials involved in this work, their Seebeck coefficients were measured in an argon-filled glovebox at room temperature by a home-built setup according to previously reported methods 9,10 on a device composed of 8 legs of this type of material deposited on a PET substrate and connected in series with their hot ends connected to cold ends. The Seebeck coefficients were extracted from the slope of the linear fit of the output voltage (V) generated under the different steady temperature difference (T) measured between the hot and cold end of the sample via the formula , where N is the number of legs, and S is the Seebeck coefficient. Concerning the characterizations of the hybrid photothermoelectric generators on PET substrates, devices were measured by exposing half of them (with their hot ends) to simulated solar AM1.5G illumination operated at 1-sun condition (100 mW cm -2 ) whereas the other half of the devices were kept in dark.

Under illumination, the temperature difference (T) between the hot and cold ends of the device was monitored over time by digital thermometers (RS PRO RS42). After reaching a steady-state T, I-V sweeps was performed on the device by a Keithley 2612B source measurement unit (SMU). Steady-state

T was achieved after one minute of light exposure (Figure S1 of supporting information (SI)).

Response over time characteristics were measured by a Keithley 2612B SMU (for voltage outputs) or by a digital thermometer (for light-induced T) when the hot side of the device was exposed or covered from simulated 1-sun illumination. For the characterizations of the photothermoelectric generators on polypropylene nonwoven fabric, devices were wear on arm with half of it exposed to natural sun light at ambient condition at room temperature ( 25 °C) with the generated voltage measured by a digital multimeter.

Supporting figures

Figure 1 .

 1 Figure 1. (a), (b), and (c) TEM images of the three batches of plasmonic Au NPs listed in TableS1 with

  (a) and (b).

Figure 2 .

 2 Figure 2. (a) The XRD spectrum and (b) the SEM image of the orthorhombic -Ag 2 Se film synthesized

Figure 3 (

 3 photothermoelectric generators on flexible PET substrates. Figure3(a) exhibits the device structure of

  3(d), by comparison to the control device, devices with their hot ends coated by Au NPs systematically exhibit higher voltage and power output under illumination. The most-optimized condition was found on the hybrid device with four NP coatings (red symbols, Fig.3(d)), where a steady-state voltage output of 34 mV and a maximum power output of 146.4 nW were measured. The increase of device performance over the first four times of NP coating indicates the beneficial effect of the photothermal agents (i.e. Au

Figure 3 .

 3 Figure 3. (a) Schematic describing the device architeture and (b) the optical image of the flexible Au-

Figure 4 .

 4 Figure 4. Pictures showing the voltage output of the photothermal Au-NP-coated PEDOT:PSS/Ag 2 Se

  NP-doped PEDOT:PSS uni-leg type photothermoelectric generators. Uni-leg-type photothermoelectric generators were also fabricated based on eight PEDOT:PSS legs with cysteamine-capped Au NPs doped inside the PEDOT:PSS layer. Briefly, 1 to 10 ml of the above-mentioned cysteamine-capped Au NP solution was precipitated by centrifugation with solvent discarded. The precipitate was then redispersed into 1 ml of PEDOT:PSS solution with the help of ultrasonic agitation. The resultant cysteamine-Au NP-doped PEDOT:PSS solution was applied by the same manner described above to form multiple Au-doped PEDOT:PSS legs onto PET substrates.

Figure S1 .

 S1 Figure S1. Temperature difference between the hot and cold side of the devices ( T) measured over time, with the 1-sun illumination (100 mW cm -2 ) modulated ON and OFF, measured on a hybrid Au-

Figure S2 .

 S2 Figure S2. Sheet resistance (black filled symbols) and conductivity (red empty symbols) of the thin

Figure S3 .

 S3 Figure S3. (a) TEM image and (b) XRD pattern of the MoS 2 powder synthesized. (c) Steady-state

Figure S4 .

 S4 Figure S4. Sheet resistance (black empty symbols) and conductivity (red filled symbols) of a stripe of

Figure S5 .

 S5 Figure S5. Optical absorbance cysteamine-capped Au NPs solution in water. Insets show the image of

Figure S6 .

 S6 Figure S6. (a) Schematic describing the device structure of the flexible "uni-leg"-type cysteamine-Au

Table 1 .

 1 The room-temperature electrical conductivity ( ), Seebeck coefficient (S), and power factor measured on the different materials investigated in this work.

	Material	PEDOT:PSS	Ag 2 Se	ITO	MoS 2
	(S cm -1 )	1107 ± 18	481 ± 30	1050 ± 0.5	11 ± 0.2
	S (µV K -1 )	19.8 ± 0.8	-129.2 ± 2.2 -14.6 ± 0.3	-32.3 ± 0.51
	Power factor	43.4	802.9	22.4	1.1
	(µW m -1 K -2 )

Acknowledgement

The authors acknowledge the PROCES project (ANR-17-CE05-0028-01) and the China Scholarship Council (CSC) for the PhD scholarships of C.Xin and Z. Hu. This project (PhD thesis of M. Chaudhary) has received funding from the European Union's Horizon 2020 research and innovation program under the Marie Skłodowska-Curie grant agreement No 754387. Appendix A. Supplementary data

photothermoelectric generators achieved also offer bright perspectives for the developments of smart textiles capable of energy harvesting.

Credit author statement

C. Xin: synthesis, device fabrication and characterizations, data analysis, figure and table preparation.

Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Supplementary data to this article can be found online.

REFERENCES Supplementary Information

18

(CTAB) solution at room temperature under magnetic stirring. Then, freshly prepared 0.6 ml of 0.01 M sodium borohydride (NaBH 4 ) solution was injected into the above solution, followed by vigorously stirring at 1200 rpm for 2 minutes. After the color of this solution changed from yellow to brown, it was kept unstirred for aging at room temperature for 30 minutes.

The growth of elongated Au NPs or nanorods was then achieved by the following procedure: In an 100-ml-volumn glass vessel, 1.8 g of CTAB and 0.3086 g of sodium oleate (NaOL) was added. Then, 50 ml of water at 50 ℃ was added into this glass vessel to dissolve the mixture under magnetic stirring.

After the solution was cooled down to 30 ℃, different amounts (listed in Table S1) of 4 mM silver nitrate (AgNO 3 ) solution was added. The whole solution was then kept unstirred for 15 minutes at 30 ℃.

Separately, 50 ml of 1mM HAuCl 4 solution was prepared and poured into the above-mentioned solution under magnetic stirring at 700 rpm for 90 minutes until the solution turned colorless. Different amounts of 12.1 mM HCl solution were then added into this solution under magnetic stirring (at 400 rpm) for 15 minutes. Then, 250 μl of 0.064 M ascorbic acid (AA) solution was added as well under stirring for 30s to form the growth solution. Next, different amounts of Au nanoparticle seed solutions prepared above were added into the growth solution under vigorously stirring at 1200 rpm for 30s. Finally, the whole solution was kept unstirred over night for 12h at 30℃ for NP growth. The final product was isolated by centrifugation at 6000 rpm for 30 minutes and decantation. The obtained NPs were then re-dispersed in deionized water. The centrifugation/decantation process was repeated twice in order to get rid of the leftover reactants and impurities. The final Au NPs were re-dispersed in deionized (DI) water.