Karam Yassin

John C Douglin

Igal G Rasin

Pietro G Santori

Björn Eriksson

Nicolas Bibent

Frédéric Jaouen

Simon Brandon

Dario R Dekel

The Effect of Membrane Thickness on AEMFC Performance: An Integrated Theoretical and Experimental Study

Keywords: Anion-exchange membrane, Anion-exchange membrane fuel cell, Water management, Modeling, Anode flooding, Cell performance. Graphical

 The critical effect of AEM thickness on AEMFC performance was investigated  Reduced membrane thickness significantly enhances AEMFC performance  Good agreement between model simulations and experimental results were obtained  Anode flooding and cathode dehydration can be suppressed by using thin AEM  AEM thickness has a strong impact on cathode hydration levels and ORR kinetics

Abstract

We present a comprehensive theoretical and experimental study of the effect of membrane thickness on the anion-exchange membrane (AEM) fuel cell (AEMFC) performance. AEMFC tests are carried out with several AEMs with thickness in the range of 5 -50 µm and assembled with a PtRu anode, and two different cathode catalysts (Pt/C or Fe-N-C). Dramatic improvements in cell performance are observed as the membrane thickness decreases, which is mainly attributed to reduced ohmic losses and enhanced water transport between the electrodes. The simulated cell performance obtained using our previously developed AEMFC model qualitatively and quantitatively explains the experimental results in the entire range of current densities (0-4 Acm - 2 ). Simulated results show that thinner membranes enhance water transport between the electrodes, mitigating the anode flooding, and resulting in increased local hydration in the membrane and cathode catalytic layer. These high hydration values enhance the anionic conductivity of the ionomeric materials, thereby improving cell performance. Furthermore, the enhanced water transport towards the cathode electrode provides sufficient water to participate in the oxygen reduction reaction, thus reducing the activation losses. Simulation modeling allows for a thorough understanding of cell behavior and aids in the development of the next generation of advanced AEMFCs.

Introduction

Over the last decade, considerable advances have been made in the development of highperformance anion-exchange membrane fuel cells (AEMFCs) [START_REF] Dekel | Review of cell performance in anion exchange membrane fuel cells[END_REF][START_REF] Gottesfeld | Anion exchange membrane fuel cells: Current status and remaining challenges[END_REF][START_REF] Ferriday | Alkaline fuel cell technology -A review[END_REF]. Affordable and abundant platinum group metal (PGM)-free catalysts for both the anode and cathode are being developed and integrated into the electrodes and cells, with promising results towards high-performance and PGM-free AEMFCs [START_REF] Douglin | A High-Temperature Anion-Exchange Membrane Fuel Cell with a Critical Raw Material-free Cathode[END_REF][START_REF] Adabi | Understanding how single-atom site density drives the performance and durability of PGM-free Fe e N e C cathodes in anion exchange membrane fuel cells d e[END_REF][START_REF] Pinsky | Metal nanoparticles entrapped in metal matrices[END_REF][START_REF] Biemolt | An Anion-Exchange Membrane Fuel Cell Containing Only Abundant and Affordable Materials[END_REF][START_REF] Truong | Platinum and platinum group metal-free catalysts for anion exchange membrane fuel cells[END_REF][START_REF] Negro | Toward Pt-Free Anion-Exchange Membrane Fuel Cells: Fe-Sn Carbon Nitride-Graphene Core-Shell Electrocatalysts for the Oxygen Reduction Reaction[END_REF][START_REF] Santori | High Performance FeNC and Mn-oxide/FeNC Layers for AEMFC Cathodes[END_REF][START_REF] Lilloja | Nitrogendoped carbide-derived carbon/carbon nanotube composites as cathode catalysts for anion exchange membrane fuel cell application[END_REF][START_REF] Xue | A highly-active, stable and low-cost platinum-free anode catalyst based on RuNi for hydroxide exchange membrane fuel cells[END_REF][START_REF] Sokka | Iron and cobalt containing electrospun carbon nanofibre-based cathode catalysts for anion exchange membrane fuel cell[END_REF][START_REF] Roy | Nickel-copper supported on a carbon black hydrogen oxidation catalyst integrated into an anion-exchange membrane fuel cell[END_REF][START_REF] Zion | Porphyrin Aerogel Catalysts for Oxygen Reduction Reaction in Anion-Exchange Membrane Fuel Cells[END_REF][START_REF] Zhu | Multi-scale study on bifunctional Co/Fe-N-C cathode catalyst layers with high active site density for the oxygen reduction reaction[END_REF]. In addition, an extensive variety of cost-effective anion exchange membranes (AEMs) have been synthesized and implemented in AEMFCs, demonstrating outstanding performance, competitive with the state-of-the-art of proton-exchange membrane fuel cells (PEMFCs) [START_REF] Hren | Alkaline membrane fuel cells: Anion exchange membranes and fuels[END_REF][START_REF] Biancolli | High-performance radiation grafted anionexchange membranes for fuel cell applications: Effects of irradiation conditions on ETFEbased membranes properties[END_REF][START_REF] Maurya | Rational design of polyaromatic ionomers for alkaline membrane fuel cells with >1 W cm-2 power density[END_REF][START_REF] Wang | Poly(aryl piperidinium) membranes and ionomers for hydroxide exchange membrane fuel cells[END_REF][START_REF] Kumari | Crosslinked quaternary phosphoniumfunctionalized poly(ether ether ketone) polymer-based anion-exchange membranes[END_REF][START_REF] Veh | Improving the water management in anion-exchange membrane fuel cells: via ultra-thin, directly deposited solid polymer electrolyte[END_REF][START_REF] Allushi | Highly conductive hydroxide exchange membranes containing fluorene-units tethered with dual pairs of quaternary piperidinium cations[END_REF][START_REF] Fan | Poly(bis-arylimidazoliums) possessing high hydroxide ion exchange capacity and high alkaline stability[END_REF][START_REF] Wang | Radiation-grafted anion-exchange membranes: the switch from low-to high-density polyethylene leads to remarkably enhanced fuel cell performance[END_REF][START_REF] Huang | Self-crosslinked blend alkaline anion exchange membranes with bi-continuous phase separated morphology to enhance ion conductivity[END_REF][START_REF] Mandal | Poly(norbornene) anion conductive membranes: homopolymer, block copolymer and random copolymer properties and performance[END_REF][START_REF] Omasta | Beyond catalysis and membranes: visualizing and solving the challenge of electrode water accumulation and flooding in AEMFCs[END_REF].

However, despite the encouraging progress in AEMFC performance, water management in the system remains a significant challenge for developing this technology [START_REF] Omasta | Beyond catalysis and membranes: Visualizing and solving the challenge of electrode water accumulation and flooding in AEMFCs[END_REF][START_REF] Omasta | Importance of balancing membrane and electrode water in anion exchange membrane fuel cells[END_REF][START_REF] Yassin | Quantifying the critical effect of water diffusivity in anion exchange membranes for fuel cell applications[END_REF][START_REF] Gerhardt | Along-the-Channel Impacts of Water Management and Carbon-Dioxide Contamination in Hydroxide-Exchange-Membrane Fuel Cells: A Modeling Study[END_REF][START_REF] Leonard | Asymmetric electrode ionomer for low relative humidity operation of anion exchange membrane fuel cells[END_REF][34][START_REF] Gutru | A comprehensive review on water management strategies and developments in anion exchange membrane fuel cells[END_REF][START_REF] Zhang | Self-adjusting anode catalyst layer for smart water management in anion exchange membrane fuel cells[END_REF]. During cell operation, specifically at high current density, deficient water supply and excessive water consumption severely reduce hydration levels at the cathode side [START_REF] Dekel | Steady state and transient simulation of anion exchange membrane fuel cells[END_REF]. Low hydration values decrease the hydroxide conductivity within the ionomeric materials of the membrane (near the cathode side) and the cathode catalytic layer (CL). This, in turn, causes higher ohmic losses in some regions of the membrane and in the cathode CL, which reduces cell performance [START_REF] Dekel | Steady state and transient simulation of anion exchange membrane fuel cells[END_REF].

Similarly, anode over-hydration and poor water removal may cause anode flooding, negatively affecting cell performance [START_REF] Omasta | Beyond catalysis and membranes: Visualizing and solving the challenge of electrode water accumulation and flooding in AEMFCs[END_REF][START_REF] Leonard | Asymmetric electrode ionomer for low relative humidity operation of anion exchange membrane fuel cells[END_REF][START_REF] Zhang | Self-adjusting anode catalyst layer for smart water management in anion exchange membrane fuel cells[END_REF]. Thus, to maximize AEMFC performance, a delicate water balance between the anode and cathode is necessary. While several strategies have been proposed to mitigate the simultaneous anode flooding and cathode drying [START_REF] Gutru | A comprehensive review on water management strategies and developments in anion exchange membrane fuel cells[END_REF], high AEMFC performance has mainly been achieved with the use of thinner AEMs, in the range of 5 -25 µm [START_REF] Mandal | Poly(norbornene) anion conductive membranes: homopolymer, block copolymer and random copolymer properties and performance[END_REF][START_REF] Mandal | The Importance of Water Transport in High Conductivity and High-Power Alkaline Fuel Cells[END_REF][START_REF] Wang | A high conductivity ultrathin anionexchange with 500+ h alkali stability for use in alkaline membrane fuel cells that can achieve 2 W cm-2 at 80 °C[END_REF].

The thickness of the AEM is indeed critical in determining the flux of water diffusing from anode to cathode during operation and thus critical in the distribution of water hydration levels perpendicular to the AEM plane. This directly affects the local anionic conductivity of the ionomeric materials (especially on or near the cathode side) and, thus, also the overall cell ohmic resistance. When a thinner AEM is used, greater water back-diffusion through the membrane is achieved, resulting in a more even through-plane water distribution (for otherwise unchanged conditions, such as a MEA, operating conditions, and current density) [START_REF] Saebea | Model based evaluation of alkaline anion exchange membrane fuel cells with water management[END_REF]. This decreases the tendency of the anode flooding, alleviates cathode catalyst layer dry-out, and also increases the hydration in the membrane locally on or near the cathode [START_REF] Oshiba | Improvement in the solid-state alkaline fuel cell performance through efficient water management strategies[END_REF] [START_REF] Veh | Improving the water management in anion-exchange membrane fuel cells: via ultra-thin, directly deposited solid polymer electrolyte[END_REF]. Despite the importance of AEM thickness on water management and, as a result, on AEMFC performance, limited studies have been dedicated to systematically investigate this effect on AEMFC behavior [START_REF] Veh | Improving the water management in anion-exchange membrane fuel cells: via ultra-thin, directly deposited solid polymer electrolyte[END_REF].

Modeling AEMFC performance has proven to be a critical tool for investigating and understanding the effect of design and operating parameters on the water and hydroxide transport across the cell [START_REF] Yassin | Quantifying the critical effect of water diffusivity in anion exchange membranes for fuel cell applications[END_REF][START_REF] Gerhardt | Along-the-Channel Impacts of Water Management and Carbon-Dioxide Contamination in Hydroxide-Exchange-Membrane Fuel Cells: A Modeling Study[END_REF]34,[START_REF] Dekel | Predicting performance stability of anion exchange membrane fuel cells[END_REF][START_REF] Yassin | A surprising relation between operating temperature and stability of anion exchange membrane fuel cells[END_REF][START_REF] Deng | Modeling of hydrogen alkaline membrane fuel cell with interfacial effect and water management optimization[END_REF][START_REF] Huo | Analytical modeling of liquid saturation jump effect for hydrogen alkaline anion membrane fuel cell[END_REF][START_REF] Huo | Experimental and analytical analysis of polarization and water transport behaviors of hydrogen alkaline membrane fuel cell[END_REF][START_REF] Machado | Three-dimensional agglomerate model of an anion exchange membrane fuel cell using air at the cathode -A parametric study[END_REF]. Generally, numerical models are validated by comparing the computed initial cell performance with the measured polarization curve. However, two or more polarization curves that are in good agreement with experimental data are required to predict cell performance accurately [START_REF] Tao | Parameter sensitivity examination and discussion of PEM fuel cell simulation model validation. Part I. Current status of modeling research and model development[END_REF]. In this study, we measure the performance of AEMFCs assembled with FAA-3 AEMs of different thicknesses, namely 5, 10, 20, 30, or 50 µm. Next, in order to quantitatively analyze the influence of these parameters on cell performance, we apply our AEMFC comprehensive model described in [START_REF] Dekel | Predicting performance stability of anion exchange membrane fuel cells[END_REF]. First, we show the comparison of six calculated polarization curves against the experimental data. Then, we discuss the effect of AEM thickness on hydration levels, ohmic resistance, water flux, and reactions distribution across the cell. Finally, we summarize the key findings and provide a guidelines for the next generation of AEM design for fuel cell technology.

Experimental

2.1

Catalyst, Electrode, and Ionomeric Materials

The electrode preparation and fuel cell testing were carried out at Technion for the measurements with Pt/C cathode, and at ICGM for measurements with the Fe-N-C cathode, with a generally similar procedure. The description is detailed separately below and in section 2.2 for these two measurement sets.

For the cells comprising a Pt/C cathode, PtRu/C anode catalyst (40% Pt and 20% Ru on carbon black, HiSPEC® 10000, with 2:1 Pt:Ru mass ratio supported on Vulcan XC 72R) and Pt/C cathode catalyst (40% Pt on carbon black, HiSPEC® 4000) were purchased from Alfa Aesar. For the cells comprising a Fe-N-C cathode, the anode catalyst was prepared from HiSPEC® 10000, by mixing it with the amount of uncatalyzed carbon black (Vulcan XC 72R) needed to decrease the PGM content to 40 wt% PtRu/C. The Fe-N-C cathode catalyst was prepared following a previously published procedure, as reported in detail in [START_REF] Adabi | Understanding how single-atom site density drives the performance and durability of PGM-free Fe e N e C cathodes in anion exchange membrane fuel cells d e[END_REF] [START_REF] Zitolo | Identification of catalytic sites for oxygen reduction in iron-and nitrogen-doped graphene materials[END_REF]. In brief, optimized amounts of anhydrous ferrous acetate, 1,10-phenanthroline and the metal-organic framework ZIF-8 were dry-mixed with planetary ball-milling, then pyrolyzed in flowing Ar at 1050°C, then pyrolyzed in flowing NH3 at 950°C. This catalyst was shown by spectroscopy and microscopy to comprise Fe mainly as atomically-dispersed Fe-N4 sites [START_REF] Adabi | Understanding how single-atom site density drives the performance and durability of PGM-free Fe e N e C cathodes in anion exchange membrane fuel cells d e[END_REF] [START_REF] Zitolo | Identification of catalytic sites for oxygen reduction in iron-and nitrogen-doped graphene materials[END_REF]. For electrode preparation, Toray carbon paper, TGP-H-060 with 5 wt% PTFE wet-proofing was purchased from FuelCellStore.

The AEMs used in this study lean on the FAA-3 chemistry, a commercial ionomeric material with anion-conducting properties, received from Fumatech GmbH (Germany). Five AEMs of different thicknesses were kindly provided by Fumatech in bromide form -FAA-3-05-rf, FAA-3-10, FAA-3-PE-20, FAA-3-30, and FAA-3-50. The last two digits correspond to the membrane thickness in µm. The anion-exchange ionomer used for the electrodes was a developmental material consisting of crosslinked polystyrene functionalized with trimethylamine, also supplied by Fumatech.

Membrane electrode assembly (MEA) fabrication

For the cells with a Pt/C cathode, the gas diffusion electrode method was employed to prepare the electrodes, following previously published procedures reported elsewhere [START_REF] Douglin | A high-temperature anion-exchange membrane fuel cell[END_REF] [START_REF] Wierzbicki | Are radicals formed during anion-exchange membrane fuel cell operation?[END_REF]. Three pairs of anodes and cathodes gas diffusion electrodes were loaded to 0.7 ± 0.05 mgPtRu cm -2 and 0.7 ± 0.05 mgPt cm -2 , respectively. Each pair of electrodes along with a 12.25 cm 2 piece of the AEM were immersed in an aqueous 1 M KOH solution for 1 h, with solution changes every 20 min, to convert it to the hydroxide form. The MEA was then assembled between two 5 cm 2 single-serpentine graphite flow field plates and the cell tightened using a 4.5 N m torque.

For the cells with PGM-free cathodes, the gas diffusion electrode method was also employed to prepare the electrodes, following previously published procedures reported elsewhere [START_REF] Santori | High Performance FeNC and Mn-oxide/FeNC Layers for AEMFC Cathodes[END_REF] [START_REF] Omasta | Beyond 1.0 W cm -2 Performance without Platinum: The Beginning of a New Era in Anion Exchange Membrane Fuel Cells[END_REF].

Three pairs of anode and cathode gas diffusion electrodes were loaded to 0.4 ± 0.02 mgPtRu cm - 2 and 0.8 ± 0.05 mgFeNC cm -2 , respectively. Each pair of electrodes along with a 9 cm 2 piece of the AEM were immersed in an aqueous 1 M KOH solution for 1 h, with solution changes every 20 min, to convert it to the hydroxide form. The MEA was then assembled in a 5 cm 2 single-cell fuel cell (Scribner), and the cell tightened using a 5.1 N m torque.

2.3

Anion-exchange membrane fuel cell testing All the AEMFCs were tested in an 850E Scribner Associates Fuel Cell test station. For the cells with the Pt/C cathode, the cell temperature was 60 °C with anode and cathode dewpoints of 60 o C, with H2 and O2 reactant gases at flow rates of 1 slpm without back-pressurization. Polarization curves were recorded potentiostatically by scanning from open-circuit voltage to 0.15 V at a scan rate of 5 mV s -1 to compare the beginning-of-life performances and mitigate against any degradation-related power losses.

For the cells with the Fe-N-C cathode, the cell temperature was 60 °C with anode and cathode dewpoints of 56 and 58 o C, respectively, with H2 and O2 reactant gases at flow rates of 1 slpm, and 2 bar back pressure applied at the cathode and no back pressure applied at the anode. Polarization curves were recorded potentiostatically by scanning from open-circuit voltage to 0.11 V, and then back up to OCV, at a moderately fast scan rate of 10 mV s -1 in order to mitigate against any degradation-related power losses. Only the OCV to 0.11 V scan is presented in the figures and used to compare the beginning-of-life performances.

Modeling approach

A one-dimensional time-dependent and isothermal model of AEMFC operation is employed in this study [START_REF] Dekel | Predicting performance stability of anion exchange membrane fuel cells[END_REF]. A five-layers MEA comprises a membrane, anode, and cathode CLs, and anode and cathode gas diffusion layers (GDLs), are included in the computational domain. The model takes into account mass transport (gas transport through GDLs and CLs as well as liquid water) across the MEA, electrochemical reactions (hydrogen oxidation and oxygen reduction reactions), and degradation kinetics of the ionomeric material within the membrane and the CLs. Details of the model system and equations as well as our numerical approach, which is primarily based on explicit finite differences, are fully described in [START_REF] Dekel | Steady state and transient simulation of anion exchange membrane fuel cells[END_REF][START_REF] Dekel | Predicting performance stability of anion exchange membrane fuel cells[END_REF]. AEM properties, design parameters, and operating conditions used for model validations are summarized in Tables 1 and2. as the AEM thickness is reduced from 30 to 5 µm. This enhancement is the result of a combination of decreased ohmic resistance and enhanced water transport through the membrane. The thinner membrane, associated with a higher gradient of water content between the electrodes, results in higher hydration levels in the membrane and on the cathode side. These high hydration values enhance the hydroxide conductivity (λ-dependent) of the membrane leading to better performance, as discussed in Ref. [START_REF] Yassin | Quantifying the critical effect of water diffusivity in anion exchange membranes for fuel cell applications[END_REF]. Furthermore, the enhanced water transport from anode to cathode with a thin membrane provides sufficient water to participate in the electrochemical reaction (especially the ORR) and thus reduces the activation losses. For example, at 0.8 V, a 5-times higher current density, from 0.1 to 0.5 A cm -2 , is observed when membrane thickness is reduced from 30 to 5 µm. Next, our AEMFC model [START_REF] Dekel | Predicting performance stability of anion exchange membrane fuel cells[END_REF] is used to simulate initial polarization curves in order to quantify the effect of membrane thickness on cell performance and increase the understanding of water management through the cell. Firstly, the model is validated against experimental data for AEMFC with a membrane thickness of 30 µm by tuning a single parameter for fitting the entire dataset, namely, the cathode electrochemical surface area. The remaining model parameters, shown in Tables 1 and2, were derived from literature review, technical data sheets (from Fumatech), and experimental measurements [START_REF] Marino | Hydroxide, halide and water transport in a model anion exchange membrane[END_REF][START_REF] Fortin | Highperformance alkaline water electrolysis using Aemion TM anion exchange membranes[END_REF][START_REF] Ion-Ebrasu | Graphene inclusion effect on anion-exchange membranes properties for alkaline water electrolyzers[END_REF][START_REF] Li | The Comparability of Pt to Pt-Ru in Catalyzing the Hydrogen Oxidation Reaction for Alkaline Polymer Electrolyte Fuel Cells Operated at 80 °C[END_REF]. Then, for validating all other AEMFC performance curves, the only parameter to update was AEM thickness. Fig. 1a depicts a good agreement between the simulated and experimental cell performances for all AEMs, over a broad range of current densities (from 0 to 3.5 A cm -2 ). The good agreement between model predictions and experimental data obtained after tuning only the electrochemical surface area provides strong confidence in the model validity and robustness. To further explore the effect of AEM thickness on cell performance, we tested additional AEMFCs with different MEAs and operating conditions. Fig. 2a shows experimental polarization curves of AEMFCs, with the dataset obtained using the Fe-N-C cathode, for different AEM thickness values (10, 30, and 50 µm). Similar to the case of Pt/C cathode, Fig. 2a shows a significant improvement in cell performance as AEM thickness decreases while, as in Fig. 1b, the Fig. 2b shows the effect on the power density curves. For example, at 0.6 and 0.8 V, we see an 8.5-fold increased current density, from 0.1 to 0.85 A cm -2 and from 0.04 to 0.34 A cm -2 , respectively, as the AEM thickness is reduced from 50 to 10 µm. However, when compared to the results obtained using the Pt/C cathode, the overall cell performance is lower. For instance, for the AEMFC made with the 30 µm AEM, we see a nearly 50 % reduction in limiting current density (at 0.2 V). Nevertheless, at 0.8

𝜂 , = 1 𝐼 𝑅 • 𝜙 𝑑𝑥 𝜙 is electric potential (V) 𝑅 is current generation (A m -3 ) [61][62 ]

Results and discussion

V, both cells with Pt/C and Fe-N-C cathode exhibit a similar current density of 0.1 A cm -2 . This is mainly due to the different character of catalysts used in the cathodes, as well as the different anode catalyst loadings and operating conditions used in the Fe-N-C-based cells. Figs. 1 and2 show that, despite using different cathode catalysts and operating conditions in the AEMFCs, membrane thickness remains a critical factor in determining the whole cell performance.

We then simulate the initial polarization curves of the cells with the Fe-N-C cathode and various AEM thickness values (10, 30, and 50 µm). Similar to the results in Fig. 1, using the cathode electrochemical surface area as a single fitting parameter, model validation is achieved against the AEMFC with a membrane thickness of 50 µm. The thicknesses of the different AEMs were then updated without additional fitting parameters. The simulated initial performance of all three cells, displayed in Fig. 2, agrees excellently with the experimental polarization curves. These validations further assure the model's confidence and accuracy for analysis and understanding of AEMFCs behavior. In Fig. 3, we analyze the effect of membrane thickness on cell performance for the case of AEMFCs with the Pt/C cathode (Fig. 1a). First, Fig. 3a shows hydration levels (λ = number of water molecules per hydroxide ion) across the cell for an AEMFC operated at a current density of 0.6 A cm -2 . The results show a significant reduction in the difference in hydration number values between the electrodes as the membrane thickness decreases from 30 to 5 μm; this is a direct result of an increase (with decreasing membrane thickness) in the dimensional gradient of  across the membrane 1 . The thin membrane thus results in more efficient water removal from the anode (lower λ), and supplies more water to the cathode side via diffusion, leading to increased λ values from 4.5 to 9, at the membrane-cathode interface, for a membrane thickness of 30 and 5 μm, respectively. The high hydration of the membrane, in turn, enhances the anionic conductivity of the ionomeric materials.

Next, in Fig. 3b we show the calculated water flux through the membrane, from the anode to the cathode (positive flux) as well as the consumption/production of water due to electrochemical reactions. The water flux for the three cells with Pt/C cathode (i.e., with AEMs of 5, 20, and 30 µm) is lower than the water consumption needed for the ORR at this current density, indicating that additional water must be supplied at the cathode through the gas phase from the cathode gas feed. Comparing the three cases, a reduced membrane thickness leads to an increase of the water flux through the membrane towards the cathode, resulting in higher cathode hydration levels, as depicted in Fig. 3a. On the other hand, as membrane thickness increases, the value of the water flux becomes very small compared to the rate of water production due to the HOR. This can quickly lead to anode flooding if additional water removal via the gas phase from the anode is not fast enough. Therefore, with a thinner AEM allowing higher water flux, the water balance is generally easier and the cell can be operated under higher current densities without anode flooding and cathode dry-out issues.

Next, in Fig. 3c, the IR (ohmic) drop is calculated as a function of current density, up to 0.8 A cm - 1 Notice that the dimensionless gradient in  appears (in Fig. 3) to decrease as membrane thickness is decreased. However, the dimensional gradient (which is proportional to the flux of water) actually increases with decreasing membrane thickness.

decay, depicted in Fig. 1a, can be attributed to voltage loss due to the membrane. The results show that the nonlinear nature of IR curves reflects the dependence of membrane resistance on current density. Decreasing membrane thickness, which is associated with enhanced water (Figs. 3a,b)

and hydroxide transport, significantly minimizes the ohmic drop across the membrane. For instance, at a current density of 0.6 A cm -2 , the IR drop with a membrane thickness of 30 μm contributes to around 160 mV of the total voltage loss, while a negligible IR drop is observed with the 5 μm membrane.

Finally, Fig. 3d demonstrates the local reaction rate across the anode and cathode CLs for the three cells, at a given current density of 0.6 A cm -2 . Besides the apparent effect of membrane thickness on the voltage losses due to transport through the membrane shown in Fig. 3c, it also has a substantial impact on the reaction distribution through the cathode CL. Thanks to the higher hydration levels in the cathode as membrane thickness decreases, a more uniform reaction distribution is observed in the cathode CL with the 5 µm AEM, indicating a better catalyst utilization. With thicker membranes (20 and 30 μm), however, the reaction is concentrated at the membrane-cathode CL interface, due to low water supply from the anode to the cathode, as can be seen in Fig. 3d. This implies a higher overpotential to drive the ORR (lower cathode performance).

Overall, the combined experimental and model results show that the effect of AEM thickness in AEMFC is complex, going well beyond the sole proportional increase of IR drop with thickness.

The AEM thickness deeply changes the water balance, modifying the gradient of ionic conductivity throughout the entire AEM (modifying therefore the IR drop in a non-proportional way with AEM thickness), but also through the cathode CL, impacting the catalyst utilization and cathode performance as well.

The effect of membrane thickness on the hydration levels is further studied and clarified in Fig. 4a. The calculated λ values at the membrane-cathode interface are plotted as iso-λ curves in a 2D color map, as a function of the AEM thickness and operating current density. From Fig. 4a it is evident that decreasing membrane thickness, for a constant current density, enhances the hydration level at the membrane-cathode interface. These enhanced hydration levels at the cathode CL also favor the AEMFC performance stability, as we previously described in [START_REF] Yassin | Quantifying the critical effect of water diffusivity in anion exchange membranes for fuel cell applications[END_REF][START_REF] Dekel | Predicting performance stability of anion exchange membrane fuel cells[END_REF][START_REF] Yassin | A surprising relation between operating temperature and stability of anion exchange membrane fuel cells[END_REF]. The calculated degradation rate as a function of lambda value is indicated on the bar graph in Fig. 4, showing an increase by several orders of magnitude of the degradation rate at the AEM-cathode interface when cathode dry-out takes place. In line with this, a recent experimental study with a Fe-N-C cathode suggests that the cell performance loss observed after 100 h operation takes place at the cathode CL, but leads essentially to ionomer degradation with little impact on the ORR activity of Fe-N-C [START_REF] Adabi | Understanding how single-atom site density drives the performance and durability of PGM-free Fe e N e C cathodes in anion exchange membrane fuel cells d e[END_REF]. Finally, the enhanced hydration levels shown in Fig. 4a cause an increase in the hydroxide conductivity in the membrane, increasing in turn the cell performance.

Conclusion

The anion-exchange membrane thickness is a critical parameter affecting the performance and stability of AEMFCs. In this study, we investigated experimentally and theoretically the effect of membrane thickness on AEMFC performance using both PGM-based and PGM-free cathode catalysts. Furthermore, we provide a comprehensive analysis of the performance of the PGMbased cathode catalyst using a 1D model of AEMFC. The high level of adequacy agreement between the experimental results and model predictions emphasizes the validity and robustness of our AEMFC model. The results show a significant enhancement in cell performance with the reduction of AEM thickness. This is mainly due to the enhanced water transport from the anode, through the AEM, towards the cathode electrode. Furthermore, thin AEM reduces the likelihood of anode flooding, increases local hydration in the membrane, and provides enough water to the cathode CL, resulting in higher hydration levels. As a result, the conductivity of the ionomeric material both in the membrane and in the cathode is enhanced. Overall, the quantitative and qualitative modeling results adequately explain water management in AEMFCs, demonstrating the importance of using thinner AEM for better water management in AEMFCs, ultimately leading to higher cell performance.

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.

  Water transport in the ionomeric phase. 𝐷 -diffusion coefficient of water in AEM xw -water mole fraction in ionomer phase -diffusion coefficient of water in AEM.

Fig. 1

 1 Fig. 1 compares the simulated performance of Pt/C cathode AEMFCs (solid lines) to experimental

Fig. 1 .

 1 Fig. 1. Comparison of simulated (solid curves) and experimental (dots) (a) initial polarization

Fig. 2 .

 2 Fig. 2. Comparison of simulated (solid curves) and experimental (dots) (a) initial polarization

Fig. 3 .

 3 Fig. 3. (a) Simulated λ profiles at a constant current density of 0.6 A cm -2 , (b) diffusive water flux

  Fig.4bshows the local hydroxide conductivity simulated in the membrane-cathode interface. Figs.4a,bshow the relative relationship between the current density and the AEM thickness to achieve the same degrees of hydration and hydroxide conductivity at the membrane-cathode interface, gaining further understanding of the influence of membrane thickness on the water management in AEMFCs.

Fig. 4 .

 4 Fig. 4. 2D contour plot (colormap) and AEMFC with Pt/C cathode operating curves (dash lines)

Table 1 .

 1 FAA-3 ionomeric materials and their properties, as used in model validations.

	Property				Value	Ref.
	Thickness				5-50 µm	Provided by
						Fumatech
	Conductivity				101 mS cm -1	[53]
	Water diffusivity			1.25•10 -11 •λ m 2 s -1	[54-56]
	AEM ion exchange capacity		1.43 mmol g -1	[53]
	Ionomer ion exchange capacity		2 mmol g -1	[57]
	Hydration number, λ ( ionomer/ air	𝜆 = 64.60𝑅𝐻 -104.90𝑅𝐻 + 64.35𝑅𝐻	[58]
	humidification equilibrium)		-11.05
	Maximum	𝜆,	𝜆	(	13	[58]
	corresponding to ionomer/water
	contact)

Table 2 .

 2 MEA design parameters and operating conditions.

	Design / operating			Value
	parameter
			Cells with Pt/C cathode	Cells with Fe-N-C Cathode
	Cell temperature		60 °C	60 °C
	Anode relative humidity (RH) 100 %	83 %
	Cathode RH		100 %	94 %
	Anode/Cathode	GDL	200 µm	200 µm
	thickness
	Anode back pressure		No backpressure	No backpressure
	Cathode back pressure		No backpressure	2 bars
	Polarization scan rate		5 mV s -1	10 mV s -1
	Anode catalyst loading		0.7 mgPtRu cm -2	0.4 mgPtRu cm -2
	Cathode catalyst loading		0.7 mgPt cm -2	0.8 mgFe-N-C cm -2
	Anode ionomer:carbon mass	0.4	0.4
	ratio
	Cathode ionomer:carbon mass	0.4	0.2
	ratio

Table . 3

 . Main model equations and parameters

	Parameter	Value/expression	Note	Ref.

Acknowledgments

This work was partially funded by the Nancy & Stephen Grand Technion Energy Program (GTEP); the European Union's Horizon 2020 research and innovation programme under grant agreement CREATE [721065], the French-Israeli Joint Research Project (PRC 2019-2021; MOST and CNRS) [PRC2347] (MOST grant No.3-15578); the Ministry of National Infrastructure, Energy and Water Resources of Israel through grant No. 3-16686 (219-11-135), the Mauerberger Foundation Fund (MFF); and the Israel Science Foundation through grant No. 1934/17. KY would also like to acknowledge the financial support of the Zvi Yanai fellowship of the Ministry of Science and technology. The authors would like to thank Dr. Michael Schuster and Dr. Bernd Bauer from Fumatech GmbH (Germany) for their generous supply of the membranes for this study.


N Recham

J-N Chotard

J-C Jumas

L Laffont

M Armand

J-M Tarascon

Ionothermal Synthesis of Li-Based Fluorophosphates Electrodes †

Fluoride-based materials have regained interest within the field of Li-ion batteries largely due to the advent of nanosizing which has transformed the insulating insertion compounds into attractive electrode materials. Herein we demonstrate the effectiveness of ionothermal synthesis in the preparation of nanometric LiFePO 4 F and LiTiPO 4 F phases with structures isotopic to tavorite LiFePO 4 (OH) at temperatures of only 260 °C, while temperatures of 600-700 °C are required to obtain coarse powders via the ceramic method. However, the redox-active phases, which are obtained in a high state of division, have lower redox voltages than LiFePO 4 despite the presence of fluorine. Additionally, LiTiPO 4 F shows staircase charge/discharge profiles with Ti 2/3þ and Ti 3/4þ couples. Though quite unusual in lithium intercalation oxides, a hint of a Li-driven Fe 3/4þ transition has been detected in LiFePO 4 F.

Introduction

Li-ion battery technology goes back to the early 1990s 1 with Sony's introduction of insertion compounds as electroactive materials for both positive and negative electrodes. Since then, advances made at both the electrode and electrolyte levels have been so tremendous that over the last 20 years their energy density has doubled to reach 200 W h/kg. Because of this high value and cycle life performance, Li-ion technology has conquered the portable electronic market and stands as a serious contender for electrical cars. However, for large-scale applications cost and safety issues must be overcome. 2 At the resources level, the cost is mainly determined by the element abundance; this is why materials based on sustainable 3d metal redox elements such as Mn (LiMn 2 O 4 , LiMn-PO 4 , Li 2 MnSiO 4 ), Ti (TiO 2 , Li 4 Ti 5 O 12 , Li 3 Ti 7 O 12 ), and Fe (LiFePO 4 , Li 2 FeSiO 4 ) are receiving growing interest with respect to Co-or Ni-based electrodes. 3 Ironically, most of the phosphate and silicate based materials, which were disregarded 25 years ago because of their insulating properties, have turned out to be some of the best materials for electrodes (e.g., LiFePO 4 ). 4 This was made possible by the chemist's ability to develop new synthetic approaches, not only to make nanopowders of these insulating phases but also to carbon "nanopaint" these powders so as to provide a good electronic wiring. [START_REF] Ravet | ECS Fall meeting[END_REF] This ability to turn insulating compounds into attractive electrode materials has rekindled interest in fluoridebased materials. Despite the fact that they are known to be poorer electronic conductors than their oxide counterparts, they display a greater potential owing to a larger ionicity of the M-F bond as compared to the M-O one.

Several fluoride based electrodes have resulted from this recent research, among which are FeF 3 , FeF 2 , and BiOF, 6 which react with Li via a conversion reaction mechanism, and Na 2 FePO 4 F, 7 Li 5 V(PO 4 ) 2 F 2 , 8 and LiVPO 4 F 9 in which the alkali metal can be reversibly extracted via a de/intercalation mechanism. The observation that the latter works at ≈ 4 V led Barker et al. 10 to synthesize most of the 3D-metal lithium fluorophosphate LiMPO 4 F phases (M = Fe, Ti, Co, Cr). The structures were then claimed, from undisclosed powder X-ray data, to be isostructural to Tavorite LiFePO 4 (OH), 11 but no electrochemical performances were mentioned.

These results give us the impetus to further explore the potential of LiMPO 4 F phases as potential electrode materials for Li-ion batteries. Fluoride based materials are generally made via solid-state reactions at high temperatures, and the LiMPO 4 F phases are no exception as they were made by reacting stoichiometric amounts of well selected precursors followed by several shake and bake steps 10 . The main drawback to ceramic synthesis, besides demanding a large amount of energy, is that it leads to highly polydisperse micrometer-sized powders, which results in poorly conducting materials. For this reason low temperature solvothermal/hydrothermal reactions conducted mainly in water are often preferred and have been used to great success in the preparation of electrochemically optimized LiFePO 4 . 12,13 However, our ability to conduct solvothermal reactions in ionic liquids 14 rather than in water allows us to preserve the nanosize and monodispersed character of the obtained powders while offering an unprecedented opportunity to expand the way in which inorganic materials can be synthesized. We recently utilized 15 our ionothermal synthetic process in the successful low temperature elaboration of Na 2 MPO 4 F (M = Fe, Mn). We describe herein how this process can be used to prepare, at mild temperatures, highly divided single-phased LiMPO 4 F (M = Ti, Fe) whose electrochemical performances will be described. In an effort to highlight the benefits of the ionothermal process over the ceramic process we have chosen to report each material's ceramic and ionothermal synthesis prior to presenting their electrochemical performances.

Results and Discussion

a. Ceramic Synthesis. The LiMPO 4 F (M = Ti, Fe) were at first synthesized via a ceramic route which consists of first ball milling, for 10 min in a SPEX-type miller, stoichiometric amounts of the suitable MF 3 and Li 3 PO 4 precursors to form 1 g of the desired LiMPO 4 F. Precursors were chosen based on the ease of the foreseen reaction (MF 3 þ Li 3 PO 4 f LiMPO 4 F þ 2LiF) owing to the strong lattice energy of LiF which acts as the driving force. The ball-milled powder was placed in either stainless steel (M = Ti) or platinum (M = Fe) containers, sealed under an argon atmosphere and heated at 700 °C for 24 h. The recovered powders were quickly rinsed with cold water to get rid of the formed LiF, washed with acetone, and oven-dried at 60 °C prior to being analyzed by X-ray diffraction. For both syntheses, the resulting powder turned out to almost be a single-phase with residual LiF despite the washes (Figure 1). All observed peaks could be indexed, except the peaks associated with LiF, in the previously reported space group P1 (2) with the following lattice parameters a for LiTiPO 4 F (Figure 1c). It is worth mentioning that attempts to synthesize LiFePO 4 F in a stainless steel tube were not as successful owing to the partial reduction of Fe 3þ by Fe, hence, the need to use Pt tubing. In contrast, we did not encounter this problem with LiTiPO 4 F as Ti (III) is stable when in contact with Fe 0z . b. Ionothermal Synthesis. To ensure the successful synthesis of single phased 3D-metal fluorophosphates, we designed our ionothermal synthesis around the chemical stability of the precursors and solvating properties of the ionic liquid with respect to those Li-and Fe-based precursors. Initially MF 3 and Li 3 PO 4 were chosen as the precursors for this reaction because of the success of these compounds in the aforementioned ceramic approach. While basic affinity and redox considerations were useful in narrowing potential ionic liquids, ultimately a survey of various ionic liquids had to be conducted starting with the most classical one, 1-ethyl-3-methylimadozolium bis-(trifluoromethanesulfonyl imide) (EMI-TFSI). Reactions to identify the optimum conditions for LiFePO 4 F were conducted as follows: To 1 g of equimolar mixture of FeF 3 and Li 3 PO 4 was added 5 mL of EMI-TFSI, the reaction was stirred for 20 min before being heated at 200 to 300 °C for 48 h, and finally the reaction was cooled down to room temperature. The recovered grey powders were (i) washed with 20 mL of acetone to get rid of traces in ionic liquid, (ii) rinsed in cold water to eliminate traces of LiF, which was expected to form during the synthesis, (iii) washed again with 20 mL of acetone, (iv) oven-dried at 60 °C, and (v) analyzed by X-ray. There was no evidence for the formation of LiFePO 4 F until 260 °C, which is also the temperature at which the TFSI -anion was found to undergo decomposition in the presence of FeF 3 with formation of elemental S (e.g., H 2 S smell) by an unknown mechanism, which caused the reduction of Fe 3þ in Fe 2þ and led to the formation of LiFePO 4 rather than LiFePO 4 F. To eschew this issue, we decided to substitute the TFSI -anion for the very robust trifluoromethanesulfonate (triflate), which was found to be stable with respect to FeF 3 . Using 1-butyl-3-methylimidazolium triflate and a synthetic temperature of 260 °C for 48 h, we succeeded in preparing single-phased LiFePO 4 F as deduced from XRD measurements (Figure 1b). Auspiciously the XRD Bragg peaks were broader than those found for the ceramic sample which is indicative of smaller particles; this was also confirmed by TEM measurements which show particle sizes as low as 20 nm (Figure 1f). This is in contrast with the ceramic process which produces particles ranging from 2 to 5 μm in size (Figure 1e).

A similar survey was conducted to determine the best ionic liquid for the synthesis of LiTiPO 4 F bearing in mind that (1) Ti 3þ is more stable than Fe 3þ so the TFSI -anion was not to be excluded and (2) that imidazolium starts to decompose in the presence of fluorides at temperatures near 280 °C probably due to the deprotonation of the acidic C2 proton. Therefore, we chose an ionic liquid protected at the C2 position with a methyl group and an OH (hydroxyl) group on the side chain to reduce the (12) Chen, J.; Wang, S.; Whittingham, M. S. reaction temperature via increasing the solubility of the precursors. The ionic liquid 1,2-dimethyl-3-(3-hydroxypropyl)imidazolium bis(trifluoromethane sulfonyl)imide was synthesized according to previous work. Using 5 mL of this ionic liquid in an autoclave, single-phased (Figure 1d) and nanosized LiTiPO 4 F powders (Figure 1h) were obtained after 48 h of heating at 260 °C while the ceramic process again produced micrometer size particles (Figure 1g).

At this point it is worth noting that, regardless of the type of ionic liquid used to carry out the synthesis, we have been able to obtain single-phased materials at temperatures as low as 260 °C as compared to the 700 °C necessary from a ceramic approach. This clearly demonstrates the great potential of ionothermal synthesis for inorganic materials once the proper ionic liquid has been identified. Further abounding about the positive attributes of ionothermal synthesis, we recall that early synthesis reports of tavorite (LiFePO 4 (F,OH) phases were calling for either a high temperature ceramic approach involving several (grinding þ heating) steps or a hydrothermal approach, which necessitated 3 days of reaction at 400 °C and under 150 MPa. 16 Although inherent to the use of ionic liquids and regardless of the type of materials we synthesized (LiFePO 4 F or LiTiPO 4 F) or the nature of the ionic liquid we used, we could at the end of each experiment recover the ionic liquid so that it could be reused. Our recovery protocol consists first in washing the reaction supernatant with a concentrated HCl solution to get rid of (LiFePO 4 F or LiTiPO 4 F) traces of powder which could remain after centrifugation. Afterward, the supernatant is washed with water to eliminate HCl traces and then diluted with dichloromethane (CH 2 Cl 2 ) to separate it from water. Finally dichloromethane is stripped with a rotary evaporator, leaving behind the regenerated ionic liquid which can be used for subsequent experiments.

The structure of these phases has been reported 10 to be isostructural with LiFePO 4 (OH) 11 , but no structural details have been given concerning the bond lengths or the Fe/Li environment. As such information is relevant to their electrochemical performances we decided to resolve their structure from powder data, as we could not grow single crystals. To perform such a study, the X-ray powder diffraction of each phase was recorded overnight on a Bruker D8 diffractometer using Cu KR or Co KR radiations (Cu λ 1 = 1.54053 A ˚; Cu λ 2 = 1.54431 A ˚, Co λ 1 = 1.78897 A ˚; Co λ 2 = 1.79285 A ˚) and equipped with a linear Vantec counter. The powder patterns were indexed (DICVOL04 17 ) in a triclinic cell, (space group P1). The structures were then refined using the Rietveld method with the Fullprof program, 18 starting with atomic coordinates found by FOX. 19 X-ray powder patterns, crystallographic data, and selected interatomic distances are shown in Figure 2a-e and Tables 1 and2, respectively. This structure is built from transition metal centered octahedra MO 4 F 2 (M=Fe, Ti) linked together by fluorine vertices in the trans position forming chains along the c axis. In both structures, two crystallographically independent metal sites are found in the unit cell (labelled Fe1 and Fe2 for LiFePO 4 F (Figure 2a, inset) and Ti1 and Ti2 for LiTiPO 4 F (Figure 2e, inset). Consequently, the MO 4 F 2 chains along the c axis constitute an alternation of M1 and M2 centered octahedra. These octahedra are slightly distorted with angles in the range 90.6-98.2°(Table 2). The 57 Fe M€ ossbauer spectrum for LiFePO 4 F, as shown by an isomer shift of 0.43 mm/s vs R-Fe, is characteristic of Fe 3þ with therefore a barely asymmetric quadrupole doublet, which confirms slightly different Fe sites. The robustness of the fitting was improved by introducing two Fe sites with slight different QS and IS (Figure 4a, Table 3) which is in good agreement with the structural findings. Regarding the nature and thus the size of the metal, those chains are more or less tilted, the Fe1-F-Fe2 angle being 130.8°compared to the Ti1-F-Ti2 angle of 134.7°(Figure 2a,e insets). Each oxygen from the equatorial plane of the octahedron is common to a PO 4 tetrahedron bridging the MO 4 F 2 leading to the formation of a three-dimensional framework and delimiting three tunnels along the [100],

[010], and [101] directions where the Li þ ions lie on the edges (Figure 2b-g). Regardless of the metal (Ti or Fe), the tunnels along the [100] (Fe: 4.20 A ˚Â 2.64 A ˚; Ti 4.17 A ˚Â 2.59 A ˚) and [010] (Fe: 4.02 A ˚Â 2.59 A ˚; Ti 4.06 A ˚Â 2.66 A ˚) directions are larger than the third one along the [101] (Fe: 1.97 A ˚Â 2.94 A ˚; Ti 2.05 A ˚Â 2.90 A ˚) and constitute probably the main diffusion path for lithium. Along that line, the appreciable ionic conductivity reported for the LiMgSO 4 F (σ = 5 Â 10 -8 S 3 cm -1 at 200 °C and E a = 0.94 eV), which adopts the same tavorite-type structure, does not come as a total surprise. 20 c. Electrochemistry. Turning to the electrochemical performances, the samples were investigated in Li-half Swagelok cells assembled in an argon drybox. Prior to being used as positive electrodes, the active materials were ball milled with a 15% weight ratio of SP carbon. The cells were loaded with 6-8 mg/cm 2 of active material as the positive electrode and 1 cm 2 Li disks as the negative electrode and were cycled in LiPF 6 -based EC-DMC electrolyte between 1.5 and 4.5 V (Figure 3b). For reasons of concision the data will only be reported for the ionic liquid samples; this data shows similar voltage-composition profiles to those of their ceramic counterparts but far better kinetics (e.g., less polarization between charge and discharge), owing to their highly divided states. The voltage composition curve of LiTiPO 4 F shows a staircase variation with two pseudo-plateaus located near 2.9 and 1.7 V. These plateaus correspond to the oxidation of Ti 3þ into Ti 4þ and the reduction of Ti 3þ into Ti 2þ when departing from x = 1 in charge or discharge, respectively. This process is fully reversible, and Li/LiTiPO 4 F can cycle over a Δx of 1 leading to a sustained reversible capacity of 150 mA h/g (Figure 3b, inset). At this point it should be recalled that a similar staircase voltage curve centered near x = 1 with two plateaus at 4.2 V (V þ3 S V 4þ ) and 1.8 V (V þ3 S V þ2 ) was reported for the isostructural LiVPO 4 F phase 8 . However, essential differences are noted (1) on the amplitude of these plateaus, which are limited to about x = 0.5 on either side in our case while they were shown to extend up to x = 0.8 for the V-based phase, and (2) on the larger voltage hysteresis for the pseudo plateaus crossover.

The LiFePO 4 F phase was found to be electrochemically active as well (Figure 3a) with a redox voltage centered around 3 V and a sustained reversible capacity of about 0.85 Li (e.g., 150 mA h/g) per unit formula. Nevertheless, the curve is atypical with the presence of a staircase discharge for LiTiPO 4 F that seems to be absent from the subsequent charge while it resurfaces in the second discharge. This is related to structural transitions with the appearance of a low temperature phase having limited kinetics with respect to the first one, so that upon oxidation its potential approaches that of the first one, hence hindering its straight appearance in the charging voltage profile. Our discharge cut off voltage may sound arbitrary, but it is not, because it is the result of studying the reversibility of several Li/LiFePO 4 F cells as a function of their discharge voltages. It was found that extending the discharge below 1.4 V triggered a conversion-type reaction leading to the presence of Fe metal as determined by M€ ossbauer measurements (not shown here).

Additionally, we explored the feasibility of electrochemically oxidizing LiFePO 4 F, which implies the rarely encountered oxidation of Fe 3þ to Fe 4þ . Surprisingly, we observed (Figure 3a inset), although limited, some reversible capacity (Δx ≈ 0.1) suggesting the feasibility of such an option. To tests this hypothesis M€ ossbauer spectra were recorded for the oxidized (x ≈ -0.1) and partially reduced (x ≈ 0.80) sample. For the oxidized sample Li 0.9 FePO 4 F, the M€ ossbauer spectrum shows a slight asymmetric doublet (Figure 4c). Successful fitting of the spectrum has been realized with two sub-spectra: the first one which dominates, with hyperfine parameters δ = 0.43 mm/s and Δ = 1.13 mm/s, is characteristic of pristine LiFePO 4 F while the second one, with a weak contribution (7%) to the total absorption and hyperfine parameters δ = 0.13 mm/s and Δ = 0.17 mm/s, reveals a change in the oxidation state of iron. Such values of isomeric shifts (IS) in the range of -0.15 to 0.15 mm/s are typical of the Fe 4þ state, [21][22][23] so we can attribute these second subspectra to oxidized Fe 4þ . The reduced (x ≈ 0.80) sample shows the typical spectra for Fe 3þ as a pristine material, which accounts for 30% of the total absorption, together with two quadrupole doublets having similar isomer shift parameters (δ = 1.25 mm/s, Δ = 2.71 mm/s; δ = 1.24 mm/s, Δ = 2.07 mm/s) but a different contribution (49 and 20%, respectively), indicative of two non-equivalent Fe 2þ sites (Figure 4b, Table 3). The structural changes associated with Li insertion/deinsertion into both LiTiPO 4 F and LiFePO 4 F phases have been studied by in situ XRD measurements. The details of this study, together with an in situ M€ ossbauer study of the Fe-based phase, have revealed the appearance of structural phase changes for each staircase voltage change.

Obviously, these phases do not compare positively with respect to LiFePO 4 , which has become the benchmark material for the battery community for two reasons: (1) they display staircase voltage composition dependence as opposed to the single insertion voltage displayed by LiFePO 4 and are preferred for applications though charge monitoring is more complex, and (2) they show a lower redox voltage leading to lower energy density. This last point is somewhat worth a thought as it reminds us that F-based compounds do not systematically imply an increase in potential. It does for isotypic compounds only, as witnessed by the increase in the redox voltage when moving from the LiFePO 4 (OH) phase (2.6 V) [START_REF] Groguenec | Presented at the ACS Meeting[END_REF] to LiFe-PO 4 F (3.0 V).

Conclusions

In summary this work has further confirmed the benefits of ionothermal over ceramic synthesis in preparing highly divided LiFePO 4 F and LiTiPO 4 F materials, although the synthesized materials fall short of meeting LiFePO 4 performances. Therefore, their rich crystal chemistry provides a valuable asset in the hunt for better positive electrode materials.

  = 5.1551(3) A ˚, b = 5.3044(3) A ˚, c = 7.2612(4)A ˚, R = 107.356(5)°, β = 107.855(6)°, γ = 98.618(5)°, and V = 173.91(2) A ˚3 for LiFePO 4 F (Figure 1a) and a = 5.1991(2) A ˚, b = 5.3139(2) A ˚, c = 7.2428(3) A ˚, R = 106.975(3)°, β = 108.262(4)°, γ = 97.655(4)°, and V = 176.10(2) A ˚3

Figure 1 .

 1 Figure 1. XRD powder patterns for the LiFePO 4 F and LiTiPO 4 F phases prepared by the ceramic (a, c) and ionothermal (b, d) processes, respectively, together with their corresponding morphology as deduced by SEM (e, g) and TEM (f, h).

Figure 2 .

 2 Figure 2. X-ray powder patterns of LiFePO 4 F (a) and LiTiPO 4 F (e). Experimental data (red dots), calculated pattern (black line), Bragg positions (green ticks), and difference curve (blue line) are shown. In both diagrams, first Bragg ticks correspond to the main phase LiMPO 4 F (M = Fe, Ti) while the second corresponds to LiF. Projections of MO 4 F 2 octahedra chains (M = Fe, Ti) and resulting angles M1-F-M2. M1 and M2 centered octahedra are represented in red and blue in insets of Figures 3a,e, respectively. LiMPO 4 F structure projections along (b and f) the [100] direction, (c and h) the [010] direction, and (d and g) the [101] direction. Tunnels along those three directions are highlighted with green dashed lines as well as inner distances. MO 4 F 2 octahedra are represented in blue and red while PO 4 tetrahedra are in purple. For clarity, lithium atoms in the tunnels at the left are omitted.

  x FePO 4 F 0.43(1) 1.13(1) 0.29(5) 93 Fe(III) x ≈ 0.1 0.13(3) 0.17(4) 0.29(5) 7 Fe(IV) a δ = isomer shift vs R-Fe, Δ = quadrupole splitting, Γ = line width.

Figure 3 .

 3 Figure 3. Voltage-composition curve for (a) Li/LiFePO 4 F and (b) Li/ LiTiPO 4 F cells using ionothermally made positive electrodes, together with their corresponding cycling capacity retention as inset. The cells were cycled at a C/15 rate. The voltage-composition curve for a LiFePO 4 F, which has been started toward oxidation, under the same C/15 current conditions, is shown as an inset in (a).

Figure 4 .

 4 Figure 4. 57 Fe M€ ossbauer spectra of LiFePO 4 F (a), Li 1þx FePO 4 F (x ≈ 0.8) (b) and Li 1-x FePO 4 F (x ≈ 0.1) (c).

Table 1 .

 1 Lattice Parameters, Cell Volumes, and Atomic Coordinates for LiFePO 4 F and LiTiPO 4 F a The given reliability factors, *R p and *R wp , are relatively high for the Fe-based compound owing to the large background in our collected data that we have not corrected for.

			LiFePO 4 F
	a (A ˚)		b (A ˚)		c (A ˚)	volume (A ˚3)
	5.1551(3)	5.3044(3)	7.2612(4)	173.91(2)
	R		β		γ
	107.357(5)°107.855(6)°98.618(5)°χ
	2		*R p		*R wp
	1.77		28.0		18.0
		Wyckoff
	atom	position	x	y	z	occupation
	Li1	2i	0.38(5)	0.68(5)	0.83(4)	0.5
	Li2	2i	0.23(4)	0.60(4)	0.73(5)	0.5
	Fe1	1a	0(-)	0 ( -)	0 ( -)	1
	Fe2	1b	0(-)	0 ( -)	1/2(-)	1
	P	2 i	0.325(4)	0.649(4)	0.251(3)	1
	O1	2i	0.621(6)	0.756(5)	0.411(4)	1
	O2	2i	0.111(6)	0.662(5)	0.360(4)	1
	O3	2i	0.318(6)	0.354(6)	0.140(4)	1
	O4	2i	0.259(6)	0.793(5)	0.08(4)	1
	F	2 i	0.121(5)	0.912(4)	0.759(4)	1
			LiTiPO 4 F
	a (A ˚)		b (A ˚)		c (A ˚)	volume (A ˚3)
	5.1991(2)	5.3139(2)	7.2428(3)	176.10(2)
	R		β		γ
	106.975(3)°108.262(4)°97.655(4)°χ
	2		R p		R wp
	2.23		15.7		12.0
		Wyckoff
	atom	position	x	y	z	occupation
	Li1	2i	0.278(17)	0.627(13)	0.781(11)	0.5
	Li2	2i	0.232(14)	0.584(15)	0.609(11)	0.5
	Ti1	1a	0(-)	0 ( -)	0 ( -)	1
	Ti2	1b	0(-)		0 ( -)	1/2(-)	1
	P	2 i	0.3238(15) 0.6372(13) 0.2481(10) 1
	O1	2i	0.626(2)	0.7579(19) 0.3980(19) 1
	O2	2i	0.124(3)	0.6670(19) 0.373(2)	1
	O3	2i	0.330(2)	0.351(2)	0.1560(19) 1
	O4	2i	0.259(2)	0.784(2)	0.0996(17) 1
	F	2 i	0.1074(19) 0.9171(16) 0.7627(14) 1

a (16) Genkina, E. A.; Kabalov, Y. K.; Maksimov, B. A.; Mel 0 nikov, O. K. Kristallografiya 1984, 29, 50. (17) Boultif, A.; Lou€ er, D. J. Appl. Crystallogr. 2004, 37, 724. (18) Rodriguez-Carvajal, J. Physica B 1993, 192, 55. (19) Favre-Nicolin, V.; Cerny, R. J. Appl. Crystallogr. 2002, 35, 734 (see also: http://objcryst.sourceforge.net).

Table 2 .

 2 Selected Interatomic Distances (A ˚) and Angles (deg) in LiFePO 4 F and LiTiPO 4 F

	LiFePO 4 F			LiTiPO 4 F
	Fe1 O1(Â2)	1.97(3) Ti1 O1(Â2)	1.965(9)
	O2(Â2)	2.04(3)		O2(Â2)	2.029(13)
	F(Â2)	2.01(3)		F(Â2)	2.004(10)
	F-Fe1-O1	90.6°F-Ti1-O1	90.5°F
	-Fe1-O2	90.6°F-Ti1-O2	92.8°O
	1-Fe1-O2	93.0°O1-Ti1-O2	93.0°F
	e2 O3(Â2)	2.04(3) Ti2 O3(Â2)	2.098(9)
	O4(Â2)	1.93(3)		O4(Â2)	1.984(11)
	F(Â2)	1.98(3)		F(Â2)	1.92(11)
	F-Fe2-O3	91.6°F-Ti2-O3	90.5°F
	-Fe2-O4	98.2°F-Ti2-O4	94.0°O
	3-Fe2-O4	91.3°O3-Ti2-O4	90.5°P
	O1	1.50(4) P	O1	1.521(11)
	O2	1.54(4)		O2	1.574(19)
	O3	1.52(4)		O3	1.481(12)
	O4	1.58(4)		O4	1.487(15)
	Li1 O1	2.4(3)	Li1 O1	2.29(7)
	O2	2.6(2)		O2	2.17(6)
	O3	1.6(3)		O3	1.92(7)
	O4	2.1(3)		O4	2.25(8)
	F	2.0(3)		F	1.89(8)
	Li2 O1	2.2(2)	Li2 O1	2.04(8)
	O2	1.9(2)		O2	1.82(9)
	O2	2.7(4)		O2	2.19(8)
	O3	2.2(2)		O3	2.28(7)
	O4	2.4(3)		F	2.10(8)
	F	1.8(2)

Table 3 .

 3 57 Fe M€ ossbauer Hyperfine Parameters of LiFePO 4 F, Li 1þx FePO 4 F, and Li 1-x FePO 4 F a

† Accepted as part of the 2010 "Materials Chemistry of Energy Conversion Special Issue".*Corresponding author. E-mail: jean-marie.tarascon@sc.u-picardie.fr.

Acknowledgment. We would like to thank M. Courty for having performed TGA/DSC measurements not reported in the paper, K. Djellab for his assistance in taking SEM/ TEM pictures, Elena M. Arroyo for sharing preliminary calculations on these materials, and members of the ALISTORE-ERI nanopositive team for fruitful discussions.


domain: sdv chim

text:
Nandy Canavy

Antoine Rouilly

Christophe Drouet

Carole Thouron

Maëlenn Aufray
email: maelenn.aufray@ensiacet.fr

Influence of Flax Fibers on Epoxide-Amine Composites: Energetics of Interphase Formation

Keywords: composites, flax fibers, epoxide-amine matrices, mixing calorimetry, thermal properties, enthalpy, interphases

Flax fibers are particularly relevant in composite fabrication due to natural availability and mechanical properties close to glass fibers. We explore flax fiber-containing epoxy-amine (DGEBA/DETA) polymers with wide industrial applicability. Flax fibers impact the glass transition temperature (DSC), with a Tg drop of 67°C at 30 wt.% loading. For deeper insight, we develop here an original mixing calorimetry approach to follow in situ DGEBA/fibers and DETA/fibers interphases. DGEBA does not interact quantitatively with flax fibers while DETA/fibers interaction is significantly exothermic and directly related to fibers content. DETA/water interaction only accounts for 25% of total heat effect. Experiments run by contacting DETA with each of the main components of flax fibers (cellulose, hemicelluloses, lignin) point out systematic exothermic effects. Considering their mean proportion in flax fibers, the sum of enthalpy contributions agrees well with the overall effect measured with entire fibers.

The main contribution (58%) arises from cellulose. XRD analyses and literature considerations allow concluding on the direct implication of DETA molecules with cellulose chains, resulting in Tg decrease. This work clarifies the chemical role of flax fibers in DGEBA/DETA thermosets, down to the molecular-scale, and highlights that mixing calorimetry is a powerful tool to follow interphase formation in situ.

Introduction

Bio-based products are increasingly attracting attention worldwide, with the objective to develop ever-more ecological and sustainable added-value materials, for a large number of applications [1-3]. In this context, plant fibers are particularly well suited to enter the composition of composite materials according to their morphology and mechanical properties, with an expected view to allow for mass decrease, lower cost and easier recycling for the manufacture of diversified eco-friendly materials. This is nowadays still not the case as it is often necessary to put at least equivalent or higher masses of flax fabric to obtain equivalent properties to those of glass/thermoset composites. Also, the price of natural fibers is still high due to the huge demand from the garment textile industry but it is expected to decrease in the future [4]. Among such natural sources, hemp, bamboo and flax (linseed) fibers are particularly appealing due to their availability and mechanical inner strength [4, 5]. Therefore, they have been the object of various studies in the field of composites. Flax fibers, arising from the Linum usitatissimum plant, may be seen for example as a good alternative to glass fibers, due to their similar mechanical properties, eco-friendly character and lower cost [5-7]. Such fibers may therefore be incorporated in [1] inorganic or organic matrices for improving their mechanical behavior.

Among organic matrices, epoxy-amine mixtures have demonstrated their relevance for a wealth of applicative purposes such as sealants, paints and coatings in diverse strategic industrial sectors such as ground transportation, aeronautics, construction and others. [2,7] They represent indeed an important class of high performance thermosetting polymer precursors.

Their major assets are a high Young modulus, resistance to corrosion, thermal stability and low shrinkage under curing. Epoxy thermosets are often obtained by polymerization of epoxide prepolymers with (poly)amine hardeners, ultimately leading to epoxy-amine networks. Such monomers may be associated to bio-sources fibers with the objective to modulate the physicochemical/mechanical properties of the final product. [2,8] In such composites, the interfacial domain between the fibers and the matrix is thus undoubtedly an important factor in the final characteristics of the material. In fact, since polymerization occurs directly in the presence of the fibers, the interaction of the latter with the epoxide and amine reactants could play a major role in tailoring the final material properties, e.g. mechanical, but also possibly on its durability upon ageing.

Even though this interphase between fibers and epoxide-amine systems likely has a direct impact on the composite properties, it did not receive specific attention to this date, and the type of interaction involved is still unclear. The chemical nature of the natural fibers considered, as well as their humidity could however play an essential role in the interphase formation. [9] Our recent observations indicated that the contact of a common polyamine hardener, namely diethylenetriamine (DETA), with flax fibers led to a shift of the amine-to-epoxide mole ratio r for the Bisphenol A diglycidyl ether (DGEBA)/DETA system. [10] These findings thus highlighted the role of the DETA/flax fibers interaction in the polymerization reaction. It is now interesting to determine the origin of these effects and to quantify them. Better understanding these phenomena could indeed allow controlling these effects and potentially taking advantage of them for further tailoring material properties. Also, because of their hydrophilicity that tends to prevent good adhesion between natural fibers and the matrix, it is common to carry out a treatment on the fibers to overcome this issue, typically involving a "sizing" agents. [11] However, considering the involvement of chemical products and dedicated steps to undergo this treatment, these types of steps remain questionable from an ecological point of view. Better apprehending the role of humidity and fiber chemistry on their interaction with the matrix may help unveil new approaches for fibers handling/engineering.

In this paper, we aimed at exploring in details such flax fibers -amine interphase (using DETA as polyamine), especially by developing an isothermal mixing microcalorimetry strategy to follow in situ interphase formation and quantify the related interaction energetics, considering a DGEBA/DETA epoxide-amine model matrix for composite manufacturing. DSC experiments confirmed the role played by flax fibers on such matrices thermal properties. To go further, the influence of the chemical components of natural flax fibers (cellulose, hemicelluloses and lignin) as well of water on the interphase properties was explored based on mixing calorimetry and X-ray diffraction (XRD) analyses.

Experimental section

Materials

Flax fibers, in the form of a 2/2 Serge fabric with a surfacic weight of 360 g m -2 , were obtained from Depestele (Le Bocasse, France). The constitutive fibers were previously coated with a sizing agent used to facilitate their weaving. The fabric was stored in a climatic chamber at 25°C, 60% HR (relative humidity) until use. To facilitate experiments, it was cut into sections of 5 mm x 3 mm.

To evaluate the effect of several parameters, the fibers were used as such or after specific drying or washing treatments. Drying was operated in an oven preset to 103°C for 24h and the dried fibers were then stored in a desiccator before analysis. The so-called "washing" treatment was undergone to eliminate the sizing agent associated to the fibers. To this aim, the latter were rinsed with ultrapure water several times in a tube until the water became transparent (elimination of the sizing agent leading to a yellowish coloration). Then, the washed fibers were dried at room temperature for one day.

In order to examine one by one the three main components of flax fibers, cellulose (Sigma Aldrich, USA), hemicellulose from corn (Limagrain, France) and lignin from wheat (CIMV, France) were used, in the form of powders.

The epoxide monomer used was Bisphenol A diglycidyl ether (DGEBA), reference DER 332 from Sigma Aldrich (molecular weight 348.52 g mol -1 ), with a functionality of 2 (i.e. two epoxide terminal groups per DGEBA molecule). The selected amine hardener, acting as a curing agent for the DGEBA epoxide compound, was diethylenetriamine (DETA) from Sigma Aldrich (molecular weight 103.17 g mol -1 ) with a functionality of 5 protons, arising from 2 primary amines and 1 secondary amine. Figure S1 (Supporting Information) reports the chemical formulas of these two monomers and the calculation relationship for the stoichiometry ratio r, taking into account the amine and epoxide functionalities. In order to target complete reticulation of the epoxy matrix, stoichiometric proportions of reactants (r = 1) were used throughout this work for the preparation of DGEBA/DETA samples. This ratio was indeed selected to offer the most favorable conditions of analysis, as it allows the final thermoset polymer to reach the maximal glass transition temperature, Tg∞. Polymerization of the DGEBA/DETA mixtures was undergone at 140°C for 1h, after a first period at room temperature for 3h, starting from 10 g DGEBA and 1.18 g DETA. In the case of DGEBA/DETA/flax fibers composites, the procedure followed is depicted in Figure 1, starting from 3 g DGEBA, 0.35 g DETA and 1.5 g fibers (~ 30 wt.% fiber content).

Mixing calorimetry

The interaction between DGEBA and DETA monomers and flax fibers was followed, and quantified from an enthalpy point of view, using isothermal mixing microcalorimetry (C80 Setaram, France). The measuring calorimeter cell made of Hastelloy® consists in two compartments separated by a Teflon® membrane. The calorimeter internal temperature was set to either 25°C (common temperature for composite preparation) or 50°C (to increase DGEBA fluidity and facilitate infusion processes), [12] as specified in the text. After temperature equilibration and start of baseline recording, the membrane is punctured thanks to the lowering of a piercing helix-shaped device inside the cell, thus allowing contact between the upper and lower compartment contents. In this work, the upper compartment typically contained liquid DETA or DGEBA (1.5 g) while the lower compartment contained the flax fiber sections or their individual chemical components as powders in proportions indicated in the text, typically up to 0.4 g. Upon contact, the corresponding heat flow is measured via a thermopile (3D sensor technology surrounding the measurement cell) composed of 19 superimposed thermo-elements, each thermo-element containing 38 thermocouples. The heat flow measurement is made relatively to an identical but empty reference cell (containing only the membrane) included in a second parallel thermopile, connected in opposition to the measurement one. Based on the calibration of the calorimeter (heat flow resolution: 0.10 µW, temperature accuracy: ± 0.1°C), integration of the heat flow signal over time, using the Calisto® software, leads to the enthalpy of the reaction/interaction studied. Potassium chloride (KCl, Merck) dissolution in water was used as regular checks for calibration. After each experiment, a blank experiment was carried out after return of the baseline to its stabilized level, to take into account the eventual asymmetry of the calorimeter setup. The obtained ΔHblank enthalpy value was then subtracted from the raw measured enthalpy to determine the corrected reaction enthalpy that will be considered (ΔHR).

Note that, in this work, all the given ∆HR values refer to interaction of the several types of compounds (fibers, water, cellulose, hemicelluloses or lignin) with a fixed mass of 1.5 g of DGEBA or DETA, respectively.

Physicochemical characterization

Differential scanning calorimetry (DSC) was used for the determination of glass transition temperature (Tg) of the various epoxide-amine-based networks prepared in this work, as indicated in the text. DSC runs were acquired on a Netzsch 204 apparatus, and data processing was performed using the Proteus® software. Calibration was checked with indium fusion. The samples were placed in pierced aluminum crucibles and the mass was measured with the accuracy of ± 0.1 mg. Three runs were programmed, from -70 °C to 180 °C, typically at a heating rate of 50 K min -1 . Between two cycles, an isothermal step of 5 min was integrated. This increase in heating rate was intended to facilitate determination of Tg. Note that the Tg values reported in this manuscript refer to the "midpoint" ones obtained at the second run.

X-ray diffraction (XRD) was used to analyze flax fibers before and after contact with DETA.

To this aim, XRD patterns were acquired using an Equinox 100 diffractometer (INEL, France) equipped with a cobalt anticathode (λCo =1.78897 Å). The whole 2θ range was recorded simultaneously using a curved detector, for 2h acquisitions.

Scanning electron microscopy (SEM) images were recorded on a FEI Quanta 450 microscope operated at 10 kV and under low vacuum (typically around 100 Pa of water vapor).

Results and discussion

One of the key parameters describing a thermosetting polymer network is its glass transition temperature (Tg). The Tg value measured for the polymerized DGEBA/DETA "reference" system at stoichiometry was measured to 146 ± 5 °C (midpoint). In contrast, the addition of flax fibers (30 wt.%) prior to polymerization drastically decreased this value down to 79 ± 5 °C (Figure 2a). This decrease of 67°C is evidence of a clear interaction between the polymeric network and the flax fibers, but it is impossible to know if the interaction between the hardener and the fibers or the resin with the fibers is physical or chemical as the fibers may absorb the monomers as by a sponge-like effect and/or react with them. Indeed, in thermosets, Tg depends on the crosslinking density as chain mobility is constrained by the network structure. [13] The occurrence of a chemical reaction seems corroborated by analyses performed in a previous study, evidencing a shift in the stoichiometric ratio r between amine and epoxide functions available for polymerization. [10] Interestingly, repetition of this DSC experiment with washed fibers -to eliminate the sizing agent as described in the experimental section -led to a very similar Tg drop to 74 ± 5 °C (Figure 2b); therefore the observed Tg decrease cannot be assigned to the sizing agent itself. Apart from the latter, water is also known to be naturally associated to flax fibers, with a mean content of 8 wt.%, [5] and could play a role in the observed phenomenon. Effects of water addition on epoxide-amine polymerization were indeed reported, at least on the kinetic parameters: [9,14] water added to the system up to 2-3 wt.% was typically shown to accelerate the polymerization reaction, probably by helping the opening of the epoxide endgroups, thus promoting the reaction with amines. [9] Beyond this amount, water was shown to play instead a negative role on epoxide-amine polymerization by forming a secondary phase with dissolved amine, making it less prone to react with epoxide groups. To check the possible influence of the water content included in raw fibers, a DSC analysis was carried out on a standard). SEM observations allowed checking the retained general shape and morphology of the fibers despite this treatment without noticeable degradation (Figure 3). In this case, the Tg value was lowered to 114 ± 5 °C (Figure 2b). This intermediate Tg value between the reference and the composite formed with raw (humid) fibers indicates that water indeed plays a role in the Tg drop, but that water alone cannot explain the drastic modification of Tg measured for raw fibers. Therefore, the organic components of the fibers likely impact the quality of the DGEBA/DETA polymer network upon curing. These observations are likely related to interfacial processes occurring between the monomer(s) and the fibers and/or water. In order to shed light on these observations, to explore underlying mechanisms, and to verify if the fibers-monomers interaction is either chemical or physical (simple differential adsorption), we developed -for the first time on such fiber-reinforced composites -an in situ analysis approach based on mixing microcalorimetry. This strategy was inspired from our recently reported works developed on other systems, essentially to follow DETA-metal interactions, which unveiled the relevance of the mixing microcalorimetry tool for exploring interphase formations in situ. [15] In a first stage, we investigated the possible interaction between the DGEBA epoxide monomer and flax fibers (1.5 g DGEBA, 0.1 g fibers). The measured enthalpy of interaction, at the typical temperature of 50°C, was found to be negligible (∆HR DGEBA/fibers = +29 mJ corresponding to +0.29 J g -1 of fibers) by being lower than the absolute value of the blank experiment (159 mJ). Absence of a detectable heat of interaction in these conditions strongly suggests that DGEBA does not interact significantly with the fibers. This result is interesting as it is counter-intuitive regarding the possible chemical reactivity of epoxide cycles e.g. in the presence of OH-exposing organic compounds, [16] and thus highlights the relevance of direct calorimetry measurements.

In a second stage, we examined the interaction between DETA and flax fibers. A typical example of evolution of the heat flow versus time is depicted in Figure 4a (50°C, 1.5 g DETA, 0.1 g fibers), and Figure S2 reports as additional illustration the full thermogram obtained at 25°C. The effect of a variation of fibers content (1.5 g DETA, fiber mass from 0 to 0.4 g) was followed at both 25 and 50°C, as reported in Figure 4b. In this case, in contrast to DGEBA, a noticeable exothermic phenomenon was systematically recorded, with a measured ∆HR typically ranging between 0 and -65 J. This led to ∆HR DETA/fibers = -210 ± 2 J g -1 of fibers at 25°C and ∆HR DETA/fibers = -180 ± 5 J g -1 of fibers at 50°C. These findings thus suggest the existence of significant interaction between DETA molecules and flax fibers; and this interaction likely plays a key role in the Tg decrease evidenced previously by DSC. Since flax fibers contain, however, a non-negligible amount of water (mean of 8 wt.%, [5] as illustrated by TG analysis on Figure S3), the potential impact of the presence of water in the ∆HR values measured deserves specific attention. To this aim, series of experiments were carried out, at both 25 and 50°C, by simply mixing DETA with increasing amounts of water (diamond symbols, plotted versus the water/DETA mole ratio, Figure 5). Such DETA/water interaction was found to be exothermic. If recalculated per gram of water, this led to the very similar enthalpy values of ∆HR DETA/water = -632 ± 11 J g -1 of water at 25°C and -636 ± 13 J g -1 of water at 50°C. Considering the mean 8 wt.% humidity of the fibers, the equivalent mass of water associated to the fibers was determined to evaluate the contribution of the DETA/water interaction in the global exothermic effect observed in the DETA/fibers system. Results have been added in Figure 5 (circle symbols). For both temperatures, leading to rather similar outcomes, comparison of the DETA/fibers and DETA/water systems allowed us to evaluate a factor of nearly 4 (precisely 4.2 at 25°C and 3.5 at 50°C) between the two sets of data;

DETA/fibers being systematically the most exothermic system. Therefore, water molecules associated to the fibers in their natural (humid) state may only explain about one-fourth of the global exothermic effect detected with the humid fibers, thus confirming the DSC results and showing that the (organic) fibers themselves interact strongly with DETA. In order to explore further the underlying mechanism of interaction between the organic components of the flax fibers and DETA, we have performed calorimetry experiments by selecting, one by one, each of the three main components of flax, namely cellulose (representing 61-71 wt.% of the total mass), hemicelluloses (18.6-20.6 wt.%) and lignin (1.8-5.7 wt.%). [17] Indeed, in short, the flax fiber structure may be described as a hemicelluloses/lignin matrix reinforced by embedded cellulosic microfibrils. Several analyses were conducted here using 0.2 g of each component separately and 1.5 g of DETA. Since all trends identified in this study were systematically observed at both calorimeter temperatures, these additional experiments were run at the temperature of 50°C that is particularly suited for infusion processes for composite fabrication. Every component of the fibers was found to react significantly and in an exothermic way with DETA. The ΔHR value found for each organic moiety (per gram) was wt.% hemicelluloses, 3.7 wt.% lignin), it was then possible to evaluate the enthalpy contribution of each of these three substances, as well as that of water (8 wt.%) to calculate a global ∆HR calc value corresponding to the composition of the natural (humid) fibers (Figure 6 and Table 1).

The resulting calculation, ∆HR calc = -203 ± 7 J g -1 , is in rather good agreement (within 12%)

with the experimental value obtained with the fibers, ∆HR DETA/fibers = -180 ± 5 J g -1 , reported above from the analysis of Figure 4b. The difference observed here remains limited, and may arise from several factors: i) in actual fibers the three components are organized in a threedimensional structure providing some degree of confinement (thus limiting somewhat accessibility to the liquid phase), ii) some degree of hydration may be associated with the organic moieties (easy equilibration with atmospheric humidity once exposed to air) and iii)

slight variations may exist in the composition of the organic moieties compared to their counterparts in actual flax fibers. Nevertheless, these findings show a good general correspondence in terms of global enthalpy of interaction, thus validating that the exothermic peak observed with flax fibers may be decomposed into four main components related (per order of importance) to cellulose, water, hemicelluloses and lignin. The existence of a heat of interaction upon contacting the DETA triamine with H2O molecules appears rather unsurprising, as at least hydrogen bonds may be expected to form. In contrast, interaction of DETA with the lignocellulosic components was less obvious. Finally, from these mixing calorimetry results, it may be inferred that the main contribution (58%) in the DETA/fibers exothermic interaction concerns the cellulose component of the fibers. This discovery was not especially expected, as cellulose is the most crystalline organic moiety contained in flax fibers. The existence of an interaction between cellulose and epoxy-amine systems has been reported in the literature by measuring curing heat-flow curves by DSC. [START_REF] Sottos | Studies in Polymer Science[END_REF] In the present work, we highlight the direct interaction between cellulose (alone or as a component of flax fibers) and the DETA polyamine, and our findings underline again the high potential of (poly)amines such as DETA to chemically interact with a variety of substrates, as we showed previously in the case of metal (hydr)oxidized surfaces. [15] Additional calorimetry experiments were carried out here in order to reach deeper understanding of the DETA/cellulose interaction.

In this view, a first hypothesis could lie in the highly alkaline character of DETA, which may provide a local concentration of OH -ions upon interaction with water molecules enclosed in the organic moiety. To test this hypothesis, cellulose was subjected to a sodium hydroxide solution in replacement for DETA in the upper compartment of the calorimetry setup. A NaOH concentration of 0.1 M was selected so as to approach the pH reached when 1.5 g DETA is contacted with few drops of water (pH ~ 13). Measurement with the mixing calorimeter showed that cellulose (0.2 g) reacted exothermically with sodium hydroxide 0.1 M (1.5 g) with an interaction enthalpy ΔHR NaOH/cellulose = -18 J g -1 of cellulose. Such an effect can likely be assigned to partial dissolution of cellulose under alkaline pH. [19] However, this value is 10 times lower than the one obtained for cellulose reaction with DETA; therefore interaction with the only ions OH -cannot explain the high exothermic enthalpy contribution measured with DETA.

Instead, a special role of the DETA molecule itself needs to be implied.

The authors concluded that, upon contact, a rearrangement of the chains occurred, with the separation of adjacent cellulose chains from the crystalline regions and DETA intercalation toward cross-linking to form a cellulose-DETA complex, involving the secondary amine as well as one primary amine (the other amino group being suggested to form additional hydrogen bonds). Cellulose being poly-hydroxylated, the involvement of -OH groups in this type of interaction may be postulated. To identify similar modifications when subjecting flax fibers to DETA, we performed X-ray diffraction analyses (XRD). To this aim, few sections of flax fibers were soaked in DETA for 10 min. The sample was then analyzed by XRD (cobalt source) and compared to the XRD pattern of initial flax fibers, as shown in Figure 7. Clear differences may be detected between the two patterns. The pattern of initial flax fibers (red upper curve) shows a main peak around 2θ = 26.5° and a doublet around 17 and 19° (also see Figure S4 for peak positioning). After contact with DETA (green lower curve), the main peak was shifted toward lower angles, the doublet disappeared, and a new peak appeared at low angles around 7°. These results strongly remind those obtained by Segal and Loeb on cellulose/DETA interaction, [20] as evidenced graphically in Figure 7 by the superimposed black-and-white patterns (note that Segal and Loeb used a copper source for XRD analyses, which was taken into account for the superimposition). While the main features discussed above for initial flax (upper red curve) are indicative of a high amount of cellulose polymorph I, [21] the pattern obtained after contact with DETA becomes closer to the crystallographic features of the DETA-cellulose complex as described by Segal and Loeb. [20] Based on this XRD analysis, it thus confirms that at least the cellulosic component of flax fibers undergo a chemical interaction with DETA, likely involving cellulose-DETA complexes as reported by Segal and Loeb. Since this transformation directly involves DETA molecules, it may thus explain the significant decrease of Tg observed by DSC after polymerization, by considering that part of the DETA amino functions are associated to the organic component(s) of flax fibers and will not readily react with epoxide groups. The black-and-white superimposed inlet is reproduced with permission, [20] 1960, Wiley). Note: the two maxima at 22.5 and 14.7° from this reference paper, relative to an XRD copper source, become respectively positioned at 26.2 and 17.1 in relation to the cobalt source used here Since this transformation directly involves DETA molecules, it may thus explain the significant decrease of Tg observed by DSC after polymerization, by considering that part of the DETA amino functions are associated the organic component(s) of flax fibers and will not readily react with epoxide groups. Such a decrease of Tg could be found useful, as for the use of plasticizers. On the other hand, if this decrease is not suitable, an increase of initial amount of hardener in the reacting mixture (leading to an excess of initial amine functions) may offset the lack of reactive functions, as described previously [10].

Conclusion

In this work, by way of DSC analyses, we showed the direct effect of flax fibers on the thermal properties of epoxy-amine systems, and in particular a drop of their glass transition temperature Tg. This effect was then scrutinized in details by mixing microcalorimetry to determine whether the resin or the hardener was chemically modified by flax fibers or simply absorbed by them. We first evidenced that the DGEBA component did not interact quantitatively with flax fibers. In contrast, DETA/fibers interaction led to a significant exothermic effect, which was quantified and directly linked to the fibers content. Thanks to additional calorimetry experiments involving each of the three main components of flax fibers -namely cellulose, hemicellulose and lignin -as well as water, we then showed that each of these moieties reacted exothermically with DETA. Taking into account the proportions of each of these compounds in flax fibers, the sum of enthalpy effects was then found in good accordance with the overall exothermic effect measured with the whole fibers. The main contribution was that of cellulose, responsible for 58% of the global heat effect, despite its

Table of contents text:

We develop a mixing calorimetry approach to follow in situ interphase formation between flax fibers and DGEBA or DETA. DETA/fibers interaction is particularly exothermic. Each

Figure 1 .

 1 Figure 1. Procedure for DGEBA/DETA/flax fibers composite preparation

Figure 2 .Figure 3 .

 23 Figure 2. a) Evidence of Tg downshift for the DGEBA/DETA system in the presence of flax fibers (30 wt.%), b) Measured Tg values for the DGEBA/DETA system in different conditions: reference and in the presence of flax fibers (raw, washed, or dried at 103°C / 24 h)

Figure 4 .

 4 Figure 4. a) Example of thermogram (mixing calorimetry) obtained by contacting DETA with flax fibers, at 50°C (the raw value ∆H leads to the interaction enthalpy ∆HR after subtraction of the blank), b) evolution of ∆HR DETA/flax fibers versus fiber mass, at 25°C and 50°C

Figure 5 .

 5 Figure 5. Evolution of ∆HR for the DETA/flax fibers and DETA/water systems upon increasing water content, at 25°C and 50°C

  the mean proportion of each component in flax(i.e. 66 wt.% cellulose, 19.6

Figure 6 .

 6 Figure 6. Schematic overview of enthalpy contributions (exothermic, absolute values) of each component of flax fibers in interaction with 1.5 g DETA (50°C), in comparison to 1 g of full fibers

Figure 7 .

 7 Figure 7. XRD patterns of initial and DETA-soaked flax fibers. The black-and-white superimposed inlet is reproduced with permission,[20] 1960, Wiley). Note: the two maxima at 22.5 and 14.7° from this reference paper, relative to an XRD copper source, become respectively positioned at 26.2 and 17.1 in relation to the cobalt source used here

  partial crystallized state. With the help of XRD analyses and other considerations from cellulose literature, we finally concluded on the direct implication of DETA molecules with cellulosic chains in the formation of DETA/cellulose complexes, thus limiting the amount of "reactive" amino functions of DETA in view of subsequent polymerization with DGEBA, hence explaining the decreased Tg values. This work sheds light on the chemical role of flax fibers in epoxy-amine thermosets, down to the molecular-scale, and highlights the relevance of the mixing calorimetry approach to follow in situ interphase formation. INP-ENSIACET and the CIRIMAT Laboratory (UMR CNRS 5085) since February 2008. Her expertise lays in adherence measurements and the investigation of chemical reactions at interfaces leading to interphase formation. She has developed in the lab mechanical test setups as a reproducible 3-point bending test. M. Aufray has more than twenty years of experience in the field of polymers and practical adhesion. In addition, she recurrently works at the interface of several teams, as her research area is highly transversal.

  component of flax fibers leads to exothermic effects, cellulose giving the most prominent enthalpy contribution. DETA interacts with cellulosic chains, which limits the amount of "reactive" amino functions for polymerization and decreases the Tg value. N. Canavy, A. Rouilly, C. Drouet, C. Thouron, M. Aufray* Influence of Flax Fibers on Epoxide-Amine Composites: Energetics of Interphase Formation ToC figure:

Acknowledgements

The authors wish to thank Zineb Fritah for support to calorimetry and Mathieu Boutin for humidity check by thermogravimetry.

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Revised: ((will be filled in by the editorial staff))

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Conflict of interest

Antoine Rouilly is Assistant Professor at the University of Toulouse in the Laboratory of Agro-Industrial Chemistry (UMR Toulouse INP/INRAE 1010). His field of expertise concerns biopolymer processing mainly for non-food applications and the role of water in the physicalchemical properties of biopolymers. He has been working on various biopolymers such as protein (storage and structural), starch and cellulose but also on complex systems such as lignocellulosic fibers or agricultural by-products. He is since 2012 the leader of a team of his lab dedicated to agro-materials and biopolymer processing. • Cellulose is responsible for 58% of the heat effect despite its crystallized state (83 characters)


Annabelle Larmure

J M Brayer

Mathieu Chanis

Eric Lucot

J C Monnet

Thierry Castel

Marjorie Ubertosi

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  Taking into account soils and climate change in assessing the production potential of a legume crop of interest: pea Taking into account soils and climate change in assessing the production potential of a legume crop of interest: pea Larmure A. 1 , Brayer J. 1 , Chanis M. 1 , Lucot E. 2 , Monnet J-C. 2 , Castel T. 3 , Ubertosi M. 1

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UMR 6282 Biogéosciences CNRS/ Université Bourgogne Franche-Comté -Institut Agro Dijon, France Pea (Pisum sativum L) is an annual seed legume cultivated for its high protein content and for its environmental services (e.g. symbiotic fixation of atmospheric N 2 ) Processing of "SOIL" semantic and cartographic data necessary to simulate pea production potential at an agricultural scale in the Burgundy Franche-Comté county (France)

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Climate change and the increasing pressure on resources are challenging the mode of agricultural production. To meet the challenge of more autonomous, efficient and sustainable protein production, it is necessary to acquire knowledge on the evolution of the protein crops production potential.

Map of soil mapping units (SMU) in