Huu-Dat Nguyen

Thi Khanh

Ly Nguyen

Emilie Planes

Jacques Jestin

Lionel Procar

Sandrine Lyonnard
email: sandrine.lyonnard@cea.fr

Cristina Iojoiu
email: cristina.iojoiu@lepmi.grenoble-inp.fr

Tailoring the Proton Conductivity and Microstructure of Block Copolymers by Counter Cation-Selective Membrane Fabrication

Keywords: Counter cation, membrane morphology control, Lewis acid strength, block copolymers, enhance membrane conductivity, perfluorosulfonic acid, SANS, polymer electrolyte

Here we report a simple-but-effective method to control the membrane morphology and transport properties of aromatic multi-block copolymers bearing perfluorosulfonic functions, via casting with different counter-cations. Five monovalent cations with different sizes, polarity, hydrophobicity, i.e., H + , Li + , K + , Cs + , TEA + , and one double-valence cation, i.e., Ca 2+ , were selected for manufacturing block copolymer membranes. We show that the counter-cation has a strong impact on the superstructure long-range order by acting as either block-separator, either block-compatibilizer, therefore tuning the thermodynamics of the self-assembly process. Hence, by selecting the cations, highly ordered or completely disordered phase-separated block morphologies can be created. The effect of counter-cation nature on the morphology is strongly reflected on the proton conductivity of acidified membranes. At 25 °C and 10% relative humidity, the acidified TEA + -cast membranes are 22 times less conductive than the acidified Cs + -cast ones. By combining microscopy and neutron scattering techniques, we reveal the direct correlation between enhanced functional properties and quality of membrane microstructure directed by the nature of cations with beneficial characteristics.

Our findings highlight the role and importance of cation selection to tailor the functional properties of multiblock ionomers applicable as solid electrolytes for energy conversion devices.

Introduction

Block copolymers are among the most studied class of polymers towards innovative applications due to their exceptional properties and modularity towards a broad range of applications. In the last decades, they were largely studied as polymer electrolytes for advanced electrochemical energy conversion and storage systems such as proton-exchange membrane fuel cells (PEMFCs), [START_REF] Ghassemi | Multiblock Sulfonated-Fluorinated Poly(Arylene Ether)s for a Proton Exchange Membrane Fuel Cell[END_REF][START_REF] Yu | Synthesis and Characterization of Sulfonated-Fluorinated, Hydrophilic-Hydrophobic Multiblock Copolymers for Proton Exchange Membranes[END_REF][START_REF] Li | A New Class of Highly-Conducting Polymer Electrolyte Membranes: Aromatic ABA Triblock Copolymers[END_REF] anionexchange membrane fuel cells (AEMFCs), [START_REF] Weiber | Anion Conducting Multiblock Poly(Arylene Ether Sulfone)s Containing Hydrophilic Segments Densely Functionalized with Quaternary Ammonium Groups[END_REF][START_REF] Liu | Anion Conducting Multiblock Copolymers with Multiple Head-Groups[END_REF] direct methanol fuel cells (DMFCs), [START_REF] Li | Controlled Disulfonated Poly(Arylene Ether Sulfone) Multiblock Copolymers for Direct Methanol Fuel Cells[END_REF][START_REF] Xue | Novel Proton Exchange Membranes Based on Sulfonated-Phosphonated Poly (p-Phenylene-Co-Aryl Ether Ketone) Terpolymers with Microblock Structures for Passive Direct Methanol Fuel Cells[END_REF] redox flow batteries (RFBs), [START_REF] Hong | Structural Effect of the Hydrophobic Block on the Chemical Stability of Ion-Conducting Multiblock Copolymers for Flow Battery[END_REF] and lithium batteries (LIBs), [START_REF] Porcarelli | Single-Ion Block Copoly(Ionic Liquid)S As Electrolytes for All-Solid State Lithium Batteries[END_REF][START_REF] Nguyen | Nanostructured Multi-Block Copolymer Single-Ion Conductors for Safer High-Performance Lithium Batteries[END_REF] etc., owing to their ability to achieve essential and antagonist properties, i.e., excellent thermomechanical strength and high ionic conductivity. In these materials, the nanostructured morphology derived from hydrophilic (ionic)-hydrophobic phase separation is the key to control the desired functional properties. [START_REF] Nguyen | Controlling Microstructure-Transport Interplay in Highly Phase-Separated Perfluorosulfonated Aromatic Multiblock Ionomers via Molecular Architecture Design[END_REF][START_REF] Li | Enhancement of Proton Transport by Nanochannels in Comb-Shaped Copoly(Arylene Ether Sulfone)S[END_REF][START_REF] Shin | Hydrocarbon-Based Polymer Electrolyte Membranes: Importance of Morphology on Ion Transport and Membrane Stability[END_REF] Indeed, the formation of an effective nanoscale ionic-channel network allowing interconnection of high mobility ionic pathways boosts ion transport across the membrane without affecting the mechanical strength. [START_REF] Li | Ion Transport by Nanochannels in Ion-Containing Aromatic Copolymers[END_REF][START_REF] Peckham | Structure-Morphology-Property Relationships of Non-Perfluorinated Proton-Conducting Membranes[END_REF][START_REF] He | Nanostructured Ion-Exchange Membranes for Fuel Cells: Recent Advances and Perspectives[END_REF] The welldefined nanophase separation originates from the thermo-dynamical incompatibility between the two constitutive blocks. [START_REF] Bates | Block Copolymer Thermodynamics: Theory And Experiment[END_REF][START_REF] Mai | Self-Assembly of Block Copolymers[END_REF] Therefore, adequate molecular design and cast process could be combined to optimize self-assembly towards enhanced performances of the manufactured membrane for their application in real devices.

Cation-conducting aromatic block copolymers were largely investigated in view of energy applications due to their high thermal/electrochemical stability, low cost, ease of synthesis and modification towards desirable properties, etc. [START_REF] Nguyen | Nanostructured Multi-Block Copolymer Single-Ion Conductors for Safer High-Performance Lithium Batteries[END_REF][START_REF] Park | Sulfonated Hydrocarbon Membranes for Medium-Temperature and Low-Humidity Proton Exchange Membrane Fuel Cells (PEMFCs)[END_REF][START_REF] Zhang | Recent Development of Polymer Electrolyte Membranes for Fuel Cells[END_REF][START_REF] Varcoe | Anion-Exchange Membranes in Electrochemical Energy Systems[END_REF][START_REF] Li | Ion Exchange Membranes for Vanadium Redox Flow Battery (VRB) Applications[END_REF] The nanophase separation of ionic aromatic block copolymers, however, is hindered by (i) their highly rigid nature, (ii) the multi-block structure (iii) the relatively high polydispersity index (PDI) related to the synthesis process (i.e., PDI close to 2 by polycondensation). To overcome these hurdles, much efforts were dedicated to design block structures with optimized block size, [START_REF] Nguyen | Controlling Microstructure-Transport Interplay in Highly Phase-Separated Perfluorosulfonated Aromatic Multiblock Ionomers via Molecular Architecture Design[END_REF][START_REF] Lee | Effects of Block Length and Solution-Casting Conditions on the Final Morphology and Properties of Disulfonated Poly(Arylene Ether Sulfone) Multiblock Copolymer Films for Proton Exchange Membranes[END_REF][START_REF] Assumma | Effects of Block Length and Membrane Processing Conditions on the Morphology and Properties of Perfluorosulfonated Poly(Arylene Ether Sulfone) Multiblock Copolymer Membranes for PEMFC[END_REF][START_REF] Fan | The Effect of Block Length upon Structure, Physical Properties, and Transport within a Series of Sulfonated Poly(Arylene Ether Sulfone)S[END_REF] or by introducing flexible perfluoroalkyl spacers. [START_REF] Mikami | Synthesis and Properties of Multiblock Copoly(Arylene Ether)s Containing Superacid Groups for Fuel Cell Membranes[END_REF][START_REF] Assumma | Synthesis of Partially Fluorinated Poly(Arylene Ether Sulfone) Multiblock Copolymers Bearing Perfluorosulfonic Functions[END_REF][START_REF] Nguyen | Perfluorosulfonyl Imide versus Perfluorosulfonic Acid Ionomers in Proton-Exchange Membrane Fuel Cells at Low Relative Humidity[END_REF][START_REF] Lee | Fuel Cell Catalyst Layers and Membrane-Electrode Assemblies Containing Multiblock Poly(Arylene Ether Sulfones) Bearing Perfluorosulfonic Acid Side Chains[END_REF] In the latter case, the perfluoroalkyl spacer simultaneously improves (i) the formation of ionic clusters by increasing the phase separation between ionic side chains and rigid polymer main chains, and (ii) the acidity of acidic function by transforming aryl sulfonic acid (pKa=1) into perfluorosulfonic acid (pKa=6) with higher proton-dissociation ability. [START_REF] Kreuer | On the Development of Proton Conducting Polymer Membranes for Hydrogen and Methanol Fuel Cells[END_REF] Recently, post-synthesis strategies such as solvent selectivity [START_REF] Lee | Effects of Block Length and Solution-Casting Conditions on the Final Morphology and Properties of Disulfonated Poly(Arylene Ether Sulfone) Multiblock Copolymer Films for Proton Exchange Membranes[END_REF][START_REF] Assumma | Effects of Block Length and Membrane Processing Conditions on the Morphology and Properties of Perfluorosulfonated Poly(Arylene Ether Sulfone) Multiblock Copolymer Membranes for PEMFC[END_REF] for film cast, thermal treatments, [START_REF] Assumma | Effects of Block Length and Membrane Processing Conditions on the Morphology and Properties of Perfluorosulfonated Poly(Arylene Ether Sulfone) Multiblock Copolymer Membranes for PEMFC[END_REF][START_REF] Lee | Solvent-Assisted Thermal Annealing of Disulfonated Poly(Arylene Ether Sulfone) Random Copolymers for Low Humidity Polymer Electrolyte Membrane Fuel Cells[END_REF][START_REF] Nguyen | Aromatic Copolymer/Nafion Blends Outperforming the Corresponding Pristine Ionomers[END_REF] crosslinking, [START_REF] Lee | Highly Conductive and Durable Poly(Arylene Ether Sulfone) Anion Exchange Membrane with End-Group Cross-Linking[END_REF][START_REF] Kim | Cross-Linked Sulfonated Poly(Arylene Ether Sulfone) Containing a Flexible and Hydrophobic Bishydroxy Perfluoropolyether Cross-Linker for High-Performance Proton Exchange Membrane[END_REF][START_REF] Kim | End-Group Cross-Linked Sulfonated Poly(Arylene Ether Sulfone) via Thiol-Ene Click Reaction for High-Performance Proton Exchange Membrane[END_REF][START_REF] Kallitsis | Crosslinked Wholly Aromatic Polyether Membranes Based on Quinoline Derivatives and Their Application in High Temperature Polymer Electrolyte Membrane Fuel Cells[END_REF] or applying electric field during film-cast process, [START_REF] Liu | Tailoring the Structure of S-PEEK/PDMS Proton Conductive Membranes through Applied Electric Fields[END_REF][START_REF] Zhao | Electric Field Processing to Control the Structure of Titanium Oxide/Sulfonated Poly (Ether Ether Ketone) Hybrid Proton Exchange Membranes[END_REF][START_REF] Wang | Electric-Field-Aligned Functionalized-Layered Double Hydroxide/Polyphenyl Ether Composite Membrane for Ion Transport[END_REF] etc., have received much attention as possible ways to control membrane microstructure.

Here, we report an easy, simple and highly efficient method to tailor the block copolymer morphology and, consequently, to enhance the functional properties of proton conducting membranes based on aromatic ionomers, by selecting the counter-cations used during membrane manufacturing.

Results and Discussion

The influence of five monocharged cations M + with different radius, Lewis acid strength (LAS), hydrophobicity and organic/inorganic nature, namely H + , Li + , K + , Cs + , TEA + (Figure 1a and Table S1), on membrane morphology and proton transport properties was systematically studied. A doublecharged cation, i.e., Ca 2+ , is also considered. The ionomer is a multiblock copoly(arylene ether sulfone) bearing perfluorosulfonic acid functions (named "In", Figure 1b-c), synthesized as described elsewhere. [START_REF] Assumma | Synthesis of Partially Fluorinated Poly(Arylene Ether Sulfone) Multiblock Copolymers Bearing Perfluorosulfonic Functions[END_REF] The molecular weight (MW) of both the FPES and PES blocks was fixed to 10000 g mol - 1 . The MW of the ionic side chain was not taken into account, and the hydrophobic/hydrophilic block ratio is the same in all samples tested in this study. To obtain In polymers with different countercations, the acidic side chains were exchanged using 2M aqueous solution of corresponding metalchloride (Li + , Ca 2+ , K + , Cs + ) salts, or neutralized with triethylamine. The M z+ -form ionomers were cast from the non-selective solvent DMAc. [START_REF] Assumma | Effects of Block Length and Membrane Processing Conditions on the Morphology and Properties of Perfluorosulfonated Poly(Arylene Ether Sulfone) Multiblock Copolymer Membranes for PEMFC[END_REF] The resulting membranes (labeled In-M z+ , Figure 1b) were thermal annealed and then acidified to obtain proton conducting membranes (labeled In-M z+ /H + , Figure 1c). For all samples, the ion-exchange capacities (IEC) measured by titration were found to be very close (1.271.30 meq. H + /g, Table S1), indicating that most counter-cations were exchanged by protons in the acidified membranes. This ensures that the transport properties of In-M z+ /H + can be impacted only by the morphology and microstructure of the membranes. By changing the countercation, we tune the anion-cation interactions, inter-chains interactions, ionic block polarity and we expect to vary the compatibility between the two blocks, which are key factors determining the final morphology and microstructure of the membranes. We report the comprehensive analysis of these cation-dependent effects by combining proton conductivity, dynamic vapor sorption (DVS), dynamical mechanical analysis (DMA), small-angle neutron scattering (SANS) and field-emission electron microscopy (FE-SEM) characterizations, bringing insights on how to best optimize membrane transport properties by controlling the membrane microstructure/morphology through counter-cation selective manufacturing process.

Cation Influence on Membrane Morphology and Transport Properties

The alpha relaxation temperatures (T) measured by DMA give information on the polymer chain dynamics in the salt-form (In-M z+ ) compounds (Figure 2a-b). The results show that T decreases linearly on increasing the size of alkali monovalent cations, which is in agreement with the results on Nafion reported by Bas et al. [START_REF] Moukheiber | Infrared and Thermal Behaviour of Proton Exchange Membrane (PEM) after Cationic Contamination[END_REF][START_REF] Bas | Key Counter Ion Parameters Governing Polluted Nafion Membrane Properties[END_REF] Despite a very different cation size, the organic and proton countercation membranes, In-TEA + and In-H + , exhibit close T values, lower than those of other monovalent cations, while the T of In-Ca 2+ membranes is much higher, proving much stronger inter-chain interactions due to the divalent nature of calcium ion. The DMA results therefore clearly evidence that the nature of the cation has a direct impact on the strengths of inter-chain interactions as well as on the phase separation. Note that In-Ca 2+ , In-Cs + , In-K + and In-Li + present a small peak around 220 °C corresponding to the relaxation of the hydrophobic phase, which is also observed in DMA curves of the acidified samples (Figure S1). [START_REF] Assumma | Effects of Block Length and Membrane Processing Conditions on the Morphology and Properties of Perfluorosulfonated Poly(Arylene Ether Sulfone) Multiblock Copolymer Membranes for PEMFC[END_REF][START_REF] Assumma | Synthesis of Partially Fluorinated Poly(Arylene Ether Sulfone) Multiblock Copolymers Bearing Perfluorosulfonic Functions[END_REF] The influence of counter-cation on acidified membrane water uptake and transport properties was further evaluated by measuring the vapor sorption and proton conductivity of acidified, In-M z+ /H + , at 25 °C as a function of relative humidity (RH) (Figure 2cd). The sorption isotherms of all In-M z+ /H + membranes, representing the number of water molecules per ionic group (labeled as ) as a function of the relative humidity (RH), present similar shapes which are typical for these ionomers. [START_REF] Nguyen | Perfluorosulfonyl Imide versus Perfluorosulfonic Acid Ionomers in Proton-Exchange Membrane Fuel Cells at Low Relative Humidity[END_REF] However,

In-H + /H + and In-TEA + /H + show slightly lower hydration at high RH (Figure 2c), e.g., 7 water molecules per ionic group at 95% RH against 8-9 for all other compounds. In general, the proton conductivity of ionomer membranes strongly depends on the water uptake. Yet, the moderate reduction in hydration cannot explain the large variations of proton conductivity. Thus, conduction of the acidified membranes is found to vary significantly depending on the counter-cation selected for manufacturing. The membrane cast with Cs + cation shows the highest proton conductivity, while the membranes cast from TEA + and H + show the lowest proton conductivity. At 95% RH, for instance, In-Cs + /H + is around 3.5 times more conductive than In-H + /H + and In-TEA + /H + , an effect that is amplified at lower hydration (Figure 2d). Indeed, at low relative humidity (i.e., 10%) where the water uptake is similar for all compounds (~1.5-2), differences in conductivity by more than a factor 20 are found amongst the various cation-selective membranes.

Figure 2. a) DMA curves of salt-form In-M z+ ionomer bearing different counter cations. b) Relaxation temperature (T) of ionic domains of In-M z+ as a function of cation size. c) Sorption isotherms of acidified membranes expressed as the number of water molecules per ionic group () as a function of relative humidity (DVS results) and d) proton conductivities of acidified membranes measured at 25 °C as a function of the relative humidity. Variations in proton conductivity of acidified membranes measured at 25 °C and 10% RH as a function of LAS (e) and cation radius (f). Dashed lines are guide

to the eyes, allowing to evidence the anomalous behavior of Ca 2+ and TEA + compounds.

Interestingly, we observe that proton conductivity increases with the radius of cast monovalent inorganic cations, i.e., H + <Li + <K + <Cs + and this trend was observed over the whole temperature range (Figure S2). Globally, the proton conductivity of the acidified membranes is strongly impacted by the nature of cation used for membrane casting. Danyliv et al. [START_REF] Danyliv | Highly Phase Separated Aromatic Ionomers Bearing Perfluorosulfonic Acids by Bottom-up Synthesis: Effect of Cation on Membrane Morphology and Functional Properties[END_REF] reported a conductivity two time lower for a random copolymer cast with counter-cation K + than Li + but to our knowledge, there is any reported study on the impact of counter-cation on the proton conductivity of membranes based on block copolymers. Figures 2e and Figure 2f represent the proton conductivities measured at 25 °C at 10% RH, plotted in function of the cation LAS and radius, respectively. These representations are helpful to identify possible correlations between the cation characteristics and the functional property of the obtained acidified membranes. In fact, we observe that all cations seem to follow a general trend, i.e., the proton conductivity decreases with increasing LAS and increases with increasing cation radius, indicating the impact of both parameters on the realization of efficient conduction in the ionomers.

However, there are two noticeable exceptions for In-Ca 2+ /H + and In-TEA + /H + . Similar tendency was also observed for the conductivity measured at 60 °C, 95% RH, and at 25 °C, 60% RH as a function of LAS (Figure S3a) and cation radius (Figure S3b). These materials seem to have a peculiar behavior: i) despite the largest cation size, the conductivity of In-TEA + /H + is much lower than that of In-Cs + /H + sample, an effect that could be explained by the polarity and hydrophobicity of TEA + , which may limit the formation and interconnection of conducting pathways, ii) by following the conductivity trend, it seems that, with equivalent cation size, the membrane cast with double-charge cation is more conductive than the one cast from mono-charge cation, which may relate to the stronger inter-chains interactions, as seen from the DMA results, and, consequently, a better organization of ionic domains.

Hence, as both IEC and water uptake are comparable in all our membranes, it can be supposed that the huge differences in conductivity must result from cation-induced modifications of ionomer nanostructure and ionic domains tortuosity.

Morphology-Property Interplay

To evaluate this hypothesis, the morphologies of the acidified membranes were investigated by microscopy and scattering techniques, which both probe the typical sizes of phase-separated domains as well as the degree of long-range order in the materials. On Figure 3 we represent the FESEM images of dried membranes (Figure 3a), together with the corresponding SANS profiles (Figure 3b) of the same materials in dry (top) and hydrated (20, bottom) states. The FESEM images reveal obvious differences on membrane microstructure as a function of countercation, as highlighted by ordering these data from best organized (top) to less organized (bottom) microstructures. The membrane cast from hydrophobic organic TEA + cation shows the worst phase separation (orange). The other membranes cast from monovalent cations display a phase-separated microstructure with hydrophobic FPES domains (dark regions) dispersed into to a continuous hydrophilic psPES phase (white regions) (see Figure 1bc for the chemical structure of FPES and psPES blocks). On increasing monovalent cation size (H + <Li + <K + <Cs + ), the long-range ordering of phase-separated nanodomains improves. While In-H + /H + is poorly organized (grey), the hydrophobic domains adopt locally a hexagonal packing for the biggest cation Cs + (brownish). The divalent Ca 2+ cation leads to a similar microstructure to that of Cs + , but with even more regular and better-defined hexagonal architecture (green). A statistical analysis of the FESEM images allowed to quantify both the mean diameter of PFES block domains (black nodules) and their polydispersity obtained from the full width at half maximum (FWHM) of the size distribution (Figure S4a). The nodule size follows a linear trend with cation size (Figure S5), except for In-H + /H + where the FPES size is bigger and the distribution much broader than for all other compounds. These results are confirmed by the 1D SANS profiles vs scattering vector Q shown on Figure 3b, which contain the usual characteristic features of the block copolymers. [START_REF] Nguyen | Controlling Microstructure-Transport Interplay in Highly Phase-Separated Perfluorosulfonated Aromatic Multiblock Ionomers via Molecular Architecture Design[END_REF][START_REF] Assumma | Effects of Block Length and Membrane Processing Conditions on the Morphology and Properties of Perfluorosulfonated Poly(Arylene Ether Sulfone) Multiblock Copolymer Membranes for PEMFC[END_REF] The low-Q correlation peak is due to long-range ordering between the hydrophilic psPES and hydrophobic FPES blocks, while the high-Q correlation peak (so-called ionomer peak) reflects the ionic domains organization, i.e., nanoscale phase separation between the "ps" side chains and PES polymer main chains. Moreover, the high-Q intensity decay as Q -4 (so-called Porod's law) indicates the existence of sharp interfaces between the ionic network and PES chains. Note that the ionomer peak is revealed by the presence of water that swells the ionic domains and induces neutron contrast, while the block peak can be detected in both dry and hydrated states. [START_REF] Nguyen | Controlling Microstructure-Transport Interplay in Highly Phase-Separated Perfluorosulfonated Aromatic Multiblock Ionomers via Molecular Architecture Design[END_REF] The peak positions Qblock and Qionic give the mean separation distance between block aggregates and ionic domains, dblock and dionic defined as d=2/Qblock/ionic, respectively, while the relative widths Qblock/Qblock and Qiono/Qiono indicate the degree of order in the materials.

The sharper the peak, the better organized is the material at the corresponding length scale, i.e., typically > 35 nm for block superstructure and 25 nanometers for the ionic phase. On Figure 3, we have classified the intensity-shifted 1D SANS scattering profiles following the FESEM increasing order. The visual inspection of the block peak shape confirms the decreasing large-scale order (increasing low-Q peak broadening) from Ca 2+ , Cs + , K + , Li + , H + to TEA + , in excellent agreement with the microscopy results. The In-TEA + /H + membrane has no block peak in dry state, and a broad, poorly defined one in hydrated state. The In-H + /H + compound is almost as disordered and, moreover, the block peak is shifted towards small Q values, reflecting the increased block separation distance compatible with the greatest size and broadest distribution of the hydrophobic nodules in this membrane (Figure S5). The best organization of block superstructure is found for the In-Ca 2+/ H + and In-Cs + /H + membranes, for which the SANS data exhibit an additional shoulder peak (in medium-Q region 0.0390.042 Å -1 , corresponding to correlation distances of 150160 Å) compatible with an hexagonal arrangement as depicted in the corresponding FESEM images. These behaviors are better visualized on Figure S4b-c andFigure S4d-e where the block peak and ionomer peak regions are highlighted, respectively. The large-scale separation distance between hydrophobic blocks is similar in all membranes (Figure S4b, dashed line), except for In-H + /H + , as noticed before. The same data plotted using normalized I/Ipeakmax and (Q-Qblock) unit axis allow to observe directly the continuous increase in block peak width from Ca 2+ to TEA + (black arrow on Figure S4c). Similarly, we see that the ionic separation distances are almost unchanged, despite the utilization of cations with different sizes during membrane casting (Figure S4d, dashed line). Looking back to the literature on sulfonated polystyrene, [START_REF] Zhou | Reconciling STEM and X-Ray Scattering Data to Determine the Nanoscale Ionic Aggregate Morphology in Sulfonated Polystyrene Ionomers[END_REF][START_REF] Jeon | Dynamic Mechanical Properties and Morphology of Sulfonated Polystyrene Ionomers Neutralized with Mixtures of Various Cations[END_REF] Nafion, [START_REF] Shi | Structure-Transport Relationship of Perfluorosulfonic-Acid Membranes in Different Cationic Forms[END_REF][START_REF] Peng | The Ion and Water Transport Properties of K+ and Na+ Form Perfluorosulfonic Acid Polymer[END_REF][START_REF] Peng | Transport Properties of Perfluorosulfonate Membranes Ion Exchanged with Cations[END_REF] and random aromatic copolymers bearing perfluorosulfonic acid [START_REF] Danyliv | Highly Phase Separated Aromatic Ionomers Bearing Perfluorosulfonic Acids by Bottom-up Synthesis: Effect of Cation on Membrane Morphology and Functional Properties[END_REF] we know that ion aggregation depends on the counter-cation size. Zhou et al. [START_REF] Zhou | Reconciling STEM and X-Ray Scattering Data to Determine the Nanoscale Ionic Aggregate Morphology in Sulfonated Polystyrene Ionomers[END_REF] found that the cation concentration in ionic aggregates with monovalent cations is 23 times higher than that of ionomers neutralized with divalent cations. Surprisingly, opposite behavior was reported for Nafion ® N212 membrane, in which the SAXS intensity of both ionic peak and matrix-knee decreased with increasing cation size. [START_REF] Shi | Structure-Transport Relationship of Perfluorosulfonic-Acid Membranes in Different Cationic Forms[END_REF] In those cases, however, the counter cations still remained inside the membrane, which is not the case here as our ionomers were completely acidified before characterization. Danyliv et al. [START_REF] Danyliv | Highly Phase Separated Aromatic Ionomers Bearing Perfluorosulfonic Acids by Bottom-up Synthesis: Effect of Cation on Membrane Morphology and Functional Properties[END_REF] reported the hydrophilic domains in the acidified random copolymer membranes obtained from Li + counter-cation are more homogeneous than those from K + , for the last two types of the ionic clusters was evidence: one of a very high organization and the other of worse organization. In this work, we observe that the ionic peak widths are moderately affected by the nature of the casting counter cation (Figure S4e, black arrow), with In-TEA + /H + membrane still exhibiting the poorer nanoscale phase separation. Therefore, quite surprisingly, the choice of the cation has a limited impact on the local ionic structure.

Figure 4. Mean block separation distances (blue circles. Open symbols: dry and closed symbols: hydrated) and hydrophobic FPES cluster sizes (red triangles) determined from the Gaussian peak position of SANS data and SEM distribution, respectively, plotted in function of the cation radius (a) and LAS (b). FWHM of SANS (Qblock/Qblock) and SEM data in function of the cation radius (c) and LAS (d). Dashed lines are guide for the eyes, indicating the moderate variations in sizes/distances with cation features (a, b), and the master curve behaviors of order parameter (c, d). Noticeable anomalies

are indicated by colored squares, e.g., H + , TEA + and Ca 2+ compounds deviate from the average behavior.

By fitting the FESEM size distribution and SANS peaks with Gaussian functions (Figure S6), we can now examine the characteristic values extracted from microscopy and scattering techniques in more details. Figure 4 displays the mean FPES cluster sizes SFPES (SEM results, red triangles) and block separation distances dblock (SANS in dry and hydrated states, empty and full blue circles, respectively) in function of the cation radius (Figure 4a) and LAS (Figure 4b). First, we observe a strong correlation between the two experimental methods employed to quantify the evolution of phase separation as a function of counter-cation. The differences in absolute values obtained by SANS and FESEM in dry state are attributed to distinct definitions: mean diameter of hydrophobic block clusters in the case of FESEM (red triangles, ranging from 200 to 350 Å typically) and separation distance between two scattering objects in SANS (blue circles). It seems that the superstructure spacing dblock is not much dependent on the nature of the casting counter-cation, except for H + , a behavior that might not be ascribed to size effect only, as H + is very close to Li + . The trend is that dblock and SFPES slightly increase on increasing cation radius. Upon increasing hydration for all types of ionomers, values of dblock increase, as expected, up to around 400 Å at = 20, except for the In-TEA + (500 Å) and, particularly, for the H + ionomer which shows much larger distances (from 700 Å in dry state to 900 Å in hydrated state). By subtracting the cluster dimension from the separation distance, one can estimate the size of the ionic block regions (Sionic defined as dblock-SFPES, Figure S5, green squares). Values are found in the range of 130 Å for dried membranes and 200 Å for hydrated ionomers. Again, the anomalous behavior of In-H + /H + is clear. Note also that Cs + , which is the largest and most hydrophobic alkaline cation, yields the smallest ionic block domains size. Regarding the degree of order in the materials, the evolution of FWHM of FESEM distributions and relative FWHM of scattering peaks in function of LAS and cation radius nicely correlate (Figure 4c-d). The behavior of all cations follows a trend (symbolized by dashed lines), except for Ca 2+ and TEA + . Interestingly, In-H + /H + sample does not deviate from the general trend regarding the level of organization. On increasing the cation radius, there is a clear diminution of the FWHM, indicating that a more homogenous and better-defined nanostructuration is formed when bigger cations were used during membrane casting. On increasing LAS, in contrast, we see an almost linear increase of FWHM values, showing that higher polarizability is not favorable to obtain ordered block morphologies. The deviation of In-Ca 2+ /H + and In-TEA + /H + from the main trend is quite insightful, as it shows that there is a direct impact of both size (TEA + is so big that it is out of the main curve) and bonding (divalent is out of the trend). Moreover, the TEA + cation is organic in nature and hydrophobic. Since our ionomers were completely acidified before characterization, the changes in FWHM from SANS and SEM measurements is exclusively linked to the membrane microstructure.

To summarize, we have fully characterized the multiscale structure of the In-M z+ /H + membranes and we can conclude that there is a relatively limited impact of cation selection on the nanoscale organization of ionic domains within the ionic blocks, while it has a strong influence on the structural arrangements at larger scale where polymer blocks segregate. Therefore, it can be supposed that the counter-cation plays an important role on the thermodynamic interactions between the copolymer and solvent, and between the two copolymer blocks. These observations allow to assess the prominent role of the counter cation in the self-assembly process by acting as an effective facilitator agent (Ca 2+ , Cs + ), a fair facilitator agent (K + , Li + ) or a preventer agent (H + , TEA + ), yielding either high-quality or poorquality phase-separated morphologies. Cations therefore can help in separating the two blocks of the polymers or, on the contrary, attenuate their by-design incompatibility. These facilitating or preventing abilities towards forming highly ordered block superstructures may arise from different properties.

TEA + is the largest counter cation, but also the more hydrophobic, organic in nature, therefore it can act as a compatibilizer between the hydrophilic and hydrophobic blocks, leading to poor separation.

Note that TEA + hydrophobicity also slightly degrades the local organization of ionic functions separated from the polymer main chains to form the ionic clusters. Regarding membrane cast from H + , the poor phase separation and bigger block size could arise from the formation of a partially crosslinked structure during thermal annealing at 150 °C, due to the dehydration between two intermolecular perfluorosulfonic groups. [START_REF] Assumma | Synthesis of Partially Fluorinated Poly(Arylene Ether Sulfone) Multiblock Copolymers Bearing Perfluorosulfonic Functions[END_REF] This behavior can be evidenced from DMA curves of all acidified membranes by an increase of storage modulus above 165 °C (Figure S1). However, the amount of acidic function lost after thermal annealing is very small as this sample has similar IEC than the others and was completely soluble in DMAc. Regarding Ca 2+ , we observed that the In-Ca 2+ /H + membrane exhibits a lower block size and a more homogenous distribution than its K + analogue, although K + is in the same row as Ca 2+ in the periodic table and has a bigger size. Hence, we may hypothesize that the divalent nature of the ion is a key property towards block segregation and, consequently, induces the excellent behavior of In-Ca 2+ /H + membrane. and the order parameter Qblock/Qblock (d). Dashed lines are guide to the eyes, showing the master curve behavior followed by all cation-types cast membranes when using the structural features as key monitoring parameters for transport efficiency Considering the in-depth structural characterization, we can revisit our understanding of the functional behavior of the polymers, e.g., their efficiency as proton conductors. On Figure 5a-b, we show the variations of proton conductivity at 30 °C and 95% RH as a function of the cation LAS (Figure 5a) and radius (Figure 5b), e.g., the usual representation we already adopted in Figure 2e-f (at 25 °C, 10% RH) and Figure S3 (at 60 °C, 95% RH, and at 25 °C, 60% RH), and we compare it to a direct structure-transport correlative representation produced by plotting the same conductivity data in function of relevant structural parameters (Figure 5c-d). As such, we select the mean ionic block domain size, Sionic (Figure 5c), which quantifies the mean extension of the connected phase where protons are transported, and the order parameter Qblock/Qblock (Figure 5d) that characterizes the quality of hydrophobic block/hydrophilic block phase separation. We see that the conductivity continuously increases with decreasing ionic domain size and increasing long-range structural order (or decreasing FWHM).

Although Ca 2+ and TEA + based materials have a peculiar behavior when considering cation characteristics as the monitoring parameter (Figure 5a-b), this is not any more the case when we rather consider structural parameters as key factors influencing proton transport. In fact, both anomalous compounds now fall on the general curve materialized by the dashed lines in Figures 5c-d. It seems that the conductivity of In-Ca 2+ /H + also follows the monovalent-cation trend. The structure-property interplay confirms the assumption about better defined and more regular ionic blocks domains via using larger cation sizes, which results in more homogeneous ps-PES domains with enhanced longrange connectivity, which is critical to improve the proton transport through the membrane. As discussed earlier, the lower the hydration level of the membrane, the higher the conductivity gap among the samples. Therefore, greatest differences in membrane proton conductivity were observed at reduced RH (Figure 2e-f). If we analyze the 10%-RH conductivity results using the dried membranes structural parameters, we obtain similar trends (Figure S7), showing the robustness of the conclusions.

Conclusions

In summary, the counter-cation used during membrane manufacturing has a strong impact on both the morphology and the functional properties of acidified aromatic ionomer membranes. The cation influence is more pronounced on the large-scale phase separation between the hydrophilic and hydrophobic blocks (40 nm) than on the small-scale phase separation between the ionic side chain and the polymer main chain (35 nm) where they have direct interaction. Indeed, the characteristics of the counter cations, i.e., their size, charge, organic vs inorganic nature, hydrophobicity and Lewis acid strength, play a key role in the self-assembly process by tuning the constitutive copolymer blocks incompatibility, therefore aiding or preventing efficient block segregation. The organization of the block superstructure improves on increasing the size and the charge of counter cations, while it degrades with increasing Lewis acid strength. Interestingly, the membranes cast with the smallest (H + ) and largest (TEA + ) counter cation have similar morphology and proton-transport properties. Both have quite disordered nanostructures due to the poor phase separation between the hydrophilic and hydrophobic blocks which results in much lower proton conductivity than the materials cast with alkaline counter cations. The acidified membrane cast with Cs + cation shows the highest proton conductivity due to the high-quality phase separated block superstructure formed by the regular packing of well-defined hydrophobic objects. The divalent Ca 2+ cation is also acting as an efficient structuring agent yielding excellent functional performances to the ionomer. With this work, we showed a simple-but-effective method to control the morphology and functional properties of aromatic ionomers for advanced applications.

Experimental

Membrane Preparation

The mixture of 1.0 g ionomer powder in 14 ml of DMAc was (i) stirred at 60 C for 24 h, (ii) left under vacuum for eliminating gas bubbles, (iii) centrifuged at 5000 rpm for removing solid impurities to obtain homogeneous solution. The solution was cast onto a glass substrate using an Elcometer 4340 Automatic Film Applicator. The "as-cast" membranes, obtained after drying at 60 °C in an oven for 24 h, were submitted to a thermal annealing by following a procedure described elsewhere. [START_REF] Assumma | Effects of Block Length and Membrane Processing Conditions on the Morphology and Properties of Perfluorosulfonated Poly(Arylene Ether Sulfone) Multiblock Copolymer Membranes for PEMFC[END_REF][START_REF] Nguyen | Aromatic Copolymer/Nafion Blends Outperforming the Corresponding Pristine Ionomers[END_REF] The obtained membranes was immersed in distilled water for 24 h to remove solvent residue and acidified two times using 2 M aqueous HCl solution, followed by washing several times with distilled water to remove acid trace. The In-H + membrane cast directly from ionomer in acidic form was also realized with the same condition for comparison.

Characterization

The IEC was measured by acid-base titration in organic medium as described elsewhere. [START_REF] Iojoiu | Chemical and Physicochemical Characterizations of Ionomers[END_REF] The DMA was performed on Q800 system (TA Instruments) with a heating rate of 2 °C/min in temperature range (i) from 30 to 270 °C for acidified membranes, and (ii) from 30 to 150400 °C for salt-form membranes depending of the Tg of the ps-PES hydrophilic block. The DVS was determined at 25 °C from 0 to 90% by steps of 10% using a DVS IGASORP (Hiden Isochema, UK). The samples (50 mg) were previously dried in the instrument by exposing to 0% RH at 80 °C for 8 h. The FESEM was analyzed with a back scattered detector (very sensitive to atomic number of atoms present in the material) by a ZEISS Ultra 55 device on dried lead-salt membranes. To increase the contrast between hydrophilic and hydrophobic domains, the proton in acidified membranes was exchanged by Pb 2+ cation. The dark regions correspond to hydrophobic domains whereas the hydrophilic domains are clear on these images. ImageJ "open source" software was used for the image analysis: a tailored procedure, which iterates a binarization, a watershed segmentation and a particle analysis, furnished the size distributions of FPES blocks. For one sample, three images with two magnifications were post-treated to further increase the image contrast and analyze the size distribution of the hydrophobic and hydrophilic domains. The SANS analysis was performed on D22 spectrometer (Institut Laue-Langevin, Grenoble, France) and on PAXY spectrometer (Laboratoire Léon-Brillouin, Paris, France). All measurements were conducted on fully-hydrated acidified membranes using experimental set-up and data reduction as described elsewhere. [START_REF] Assumma | Effects of Block Length and Membrane Processing Conditions on the Morphology and Properties of Perfluorosulfonated Poly(Arylene Ether Sulfone) Multiblock Copolymer Membranes for PEMFC[END_REF][START_REF] Assumma | Synthesis of Partially Fluorinated Poly(Arylene Ether Sulfone) Multiblock Copolymers Bearing Perfluorosulfonic Functions[END_REF] The through-plane proton conductivity () were obtained from a lab-made conductivity cell 11 using a Material Mates 7260 analyzer from 5 Hz to 10 MHz. The acidified ionomer films (60-80 m in thickness) were sandwiched between two stainless steel electrodes (blocking electrodes). The conductivity measurements were carried out (i) at 25 °C as a function of RH from 10 to 90% with step of 10% and (ii) at 95% as a function of cell temperature from 30 to 90 °C with step of 10 °C. The samples were stabilized at measuring condition for 5 h prior to measurement and Zview software was used to analyze and fit the data. The equivalent circuit and fitting parameters were depicted elsewhere. [START_REF] Nguyen | Controlling Microstructure-Transport Interplay in Highly Phase-Separated Perfluorosulfonated Aromatic Multiblock Ionomers via Molecular Architecture Design[END_REF] The conductivity was calculated by Equation 1:

𝜎 = 𝐿 𝑅 × 𝑆 (𝟏)
where L is sample thickness, S is upper electrode surface area, and R is bulk membrane resistance.

Figure 1 .

 1 Figure 1. (a) Cation selection for membrane manufacturing and their main properties, e.g. nature, valence, size and lewis acid strength (LAS): H + (black), Li + (blue), K + (red), Cs + (pink), TEA + (orange) and Ca 2+ (green). Chemical structure of In ionomer in salt form after casting, In-M z+ (b) and acidic form after acidification, In-M z+/ H + (c). The ionomer is composed of a hydrophobic fluorinated

Figure 3 .

 3 Figure 3. a) FESEM micrographs of dried acidified ionomer films as a function of the cation used during manufacturing, presented from lowest (bottom, H + and TEA + ) to highest (top, Ca 2+ and Cs + ) ordered structures. The insets show the segregation between hydrophobic (black) and hydrophilic (white) regions. In the case of Ca 2+ and Cs + , a local hexagonal arrangement is noticed. To enhance the contrast, protons of acidified membranes were exchanged by Pb 2+ cations. b) Corresponding 1D

Figure 5 .

 5 Figure 5. Proton conductivity at 30 °C and 95% RH of the various acidified ionomers as a function of the LAS (a) and the cation radius (b), showing the anomalous behavior of Ca 2+ and TEA + -cast

Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENT

The authors acknowledge the Institut Laue-Langevin (ILL, Grenoble, France) and Laboratoire Léon Brillouin (LLB, CEA Saclay, France) for providing the neutron facilities used in this work via proposal numbers 9-11-1709 (DOI: doi:10.5291/ILL-DATA.9-11-1709) and 150, respectively. The authors acknowledge the financial support from the French National Research Agency within the NSPEM project (ANR-16-CE05-0016), and the Centre of Excellence of Multifunctional Architectured Materials "CEMAM" AN-10-LABX-44-01).


domain: chim phys phys.cond.cm-ms

text:
Dr Tommaso Nicolini
email: tommaso.nicolini@enscbp.fr

Jokubas Surgailis

Dr Achilleas Savva

Dr Alberto D Scaccabarozzi

Dr Rana Nakar

Dr Damien Thuau

Dr Guillaume Wantz

Dr Lee J Richter

Dr Olivier Dautel

Georges Hadziioannou

Natalie Stingelin
email: natalie.stingelin@gatech.edu

King Abdullah

Prof G Hadziioannou

Prof. N N Stingelin

A Low-Swelling Polymeric Mixed Conductor Operating in Aqueous Electrolytes

Keywords: mixed conduction, organic electrochemical transistors, hydrophilic conjugated polymers, poly

ou non, émanant des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

Importantly, the initial P3HHT film thickness is essentially restored after de-doping while p(g2T-TT) remains substantially swollen. We, thus, expand the materials-design toolbox for the creation of lowswelling soft mixed conductors with tailored properties and applications in bioelectronics and beyond.

The demand for soft mixed conductors, e.g., for use in batteries [1] and organic bioelectronics device platforms [2] has led to numerous strategies for the rational chemical design of materials able to transduce ionic into electronic currents and vice versa [3]. One important feature that these materials must possess is the ability to take up ions. However, efficient ion uptake/flow is often accompanied by an intrinsic volumetric expansion of the active material, rendered more pronounced in case excess solvent is also incorporated into the material during ion uptake. Such an expansion/shrinkage can compromise mechanical integrity during cycling, limiting device stability as was shown for batteries [4] as well as for bioelectronics devices [5,6] and pronounced volumetric changes associated with ion insertion/intercalation have been recognized as one of the major limitations in lithium ion-based technologies [7].

Chemical design may assist in modulating the interplay between swelling-induced enhanced ion dynamics and undesirable effects on stability. In the bioelectronics area, targeted structures typically contain conjugated backbones to provide for good electronic charge transport; while flexible, often polar, side chains are used to assist with ion diffusion, to introduce solubility/processability and to aid with the materials' compatibility with biological environments. More specifically, one promising approach has relied on functionalizing conducting [8] and semiconducting [9,[START_REF] Giovannitti | Proc. Natl. Acad. Sci[END_REF][START_REF] Giovannitti | [END_REF] polymer backbones with hydrophilic substituents, such as glycolated side chains. This functionalization has proven to be extremely versatile, leading to materials that allowed successful fabrication of p-type [START_REF] Giovannitti | Proc. Natl. Acad. Sci[END_REF], n-type and ambipolar [START_REF] Giovannitti | [END_REF] organic electrochemical transistors (OECT)s. However, oligo(ethylene glycol) (and similar) side chains render materials purification difficult; they also can lead to structural disorder such as backbone torsionboth effects that can limit electron/hole transport [12]. Moreover, introduction of oligo(ethylene glycol) side chains frequently results in high sensitivity of important thermal phase transitions to environmental conditions (e.g., humidity). To give an example, the glass transition temperature, T g , of poly[3tetra(ethylene glycol)thiophene]a material comprised of a polythiophene backbone and oligo(ethylene glycol) side chainsis notably affected by exposure to a humid environment [13]. In turn, the torsional backbone order and, hence, optoelectronic features significantly vary with temperature [13]. Structural [14] and mechanical properties also will be affected, rendering systematic establishment of structure/property relations intricate. In this communication, we demonstrate that in poly [3-(6-hydroxy)hexyl thiophene] (P3HHT; see inset in Figure 1a for chemical structure), hydroxylated alkyl side chains introduce mixed conduction when operated with aqueous electrolytes, as work on random copolymers of poly(3-hexyl thiophene) (P3HT) and P3HHT already indicated [15], but without excessive swelling and/or unwanted plasticizing effects observed in materials with oligoethylene-glycol side chains [13,16,17,18]. The synthetic procedure and chemical characterization of the P3HHT are given in the Supplementary Information (Scheme S1, Figure S1-8).

In order to assess the swelling behavior of P3HHT upon exposure to aqueous electrolytes, we first characterized the volumetric change upon continuously flowing an aqueous solution of 0.1 mol/L KCl over P3HHT films, making a comparison to a reference sample in air (ambient relative humidity of (30 to 40) %) using Quartz Crystal Microbalance with Dissipation monitoring (QCM-D). Initially, no electrical bias was applied to obtain insights into the "passive" swelling behavior of P3HHT (Figure 1b). We employed (92 ± 4) nm thick films, and deduced changes in mass uptake from the measured frequencies, f, utilizing the Sauerbrey equation and converted these into thickness changes (see Supplementary Information for more details). KCl electrolyte solutions were used for these measurements as this allowed relatively straightforward comparison with literature data [19].

Reassuringly, based on the deduced mass uptake of the P3HHT film/KCl electrolyte systems, we infer a very minor thickness change (+2.4 %) of the polymer and, in turn, a nearly negligible volumetric change, assuming that the cross-sectional area is not affected by swelling in the configuration used for this material because we used films 'pinned' to a substrate (Figure 1c). This minute change is in contrast to the behavior

of poly(2-(3,3′-bis(2-(2-(2-methoxyethoxy)ethoxy)ethoxy)-[2,2'-bithiophen]-5-yl)thieno[3,2-b]thiophene) (p(g2T-TT))
, for which a volumetric (thickness) change of at least (+10 to +15) % was calculated under these conditions (see Figure 1b, QCM-D section of SI for multiple harmonics data, Figure S9, and modeling approach). The contrast is even stronger with respect to cross-linked poly(3,4ethylenedioxythiophene):polystyrene sulfonate (PEDOT:PSS), which swells by 85 % in the same scenario [20]. Since P3HHT displays a similar surface hydrophilicity as, e.g., PEDOT:PSS, deduced from water contact-angle measurements (Figure S10), our initial "passive" swelling data suggests that P3HHT may feature a lower bulk hydrophilicity compared to PEDOT:PSSand, likely, compared to p(g2T-TT) and other conjugated polymers with highly polar side-chain functionalization. While comparison with materials other than p(g2T-TT) is rendered challenging as often different electrolytes and conditions were used in the literature, a passive swelling behavior as low as found for P3HHT was reported only for poly [6-(thiophen-3-yl)hexane-1-sulfonate tetramethylammonium]-co-poly(3-hexyl thiophene) (PTHS - TMA + -co-P3HT) comprising a large fraction of P3HT (77 mol%; i.e., a relatively apolar material that does not allow significant OECT operation). PTHS -TMA + -co-P3HT copolymers of low P3HT content featured similar behavior as p(g2T-TT) [21]; while poly(7-glycol-naphthalene-1,4,5,8-tetracarboxylic-diimide-bisthiophene) [22] (i.e., p(gNDI-T2)) seems to passively swell more than PEDOT:PSS (values extracted from literature are reported in Table S1 in the SI). We proceeded to investigate the electrolyte uptake behavior of both P3HHT and p(g2T-TT), using again QCM-D but this time with a set-up equipped with an electrochemical module (e-QCM-D) to apply, in a pulsed fashion, a bias of 0.6 V, which leads to ion/electrolyte injection into the bulk of the film structures (Figure 2). A few observations can be immediately made. We can infer from the small change in the oscillation frequency upon application of a voltage bias that P3HHT undergoes a small mass uptake upon electrolyte exposure/application of a bias pulse (Figure 2a,b). This is very different for p(g2T-TT), where we observe a notable difference in frequency, f, in the presence/absence of an electric field, indicating a significant mass uptake upon application of the 0.6 V bias [Note that the frequency of the e-QCM-D oscillations changes when the mass on the sensor changes [23]; further details are given in the e-QCM-D section of the SI, in addition to frequency and dissipation data for multiple harmonics are given in Figure S11 andS12).]

Figure 2 Comparison of e-QCM-D data for P3HHT and p(g2T-TT) upon application of a bias pulse of ∆V = 0.6 V showing minor changes in a) frequency and b) dissipation variation for P3HHT with respect to p(g2T-TT). c) By fitting the frequency and dissipation data over multiple harmonics of the QCM, the mass uptake was calculated, from which a volumetric change can be deduced (see Supplementary Information for details). Notably different responses for the two polymers are found. Coupling e-

QCM-D results with electrical data recorded simultaneously, contributions to mass uptake from ions and water injected during application of a bias can be obtained, given here in dashed-dotted lines and dashed lines, respectively.

Additional information can be obtained from the dissipation of the signala measure of energy loss during sensor oscillation that reflects the softness of the sample [23]. We record relatively small changes in dissipation for P3HHT compared to p(g2T-TT), suggesting that swollen films of P3HHT are mechanically stiffer than p(g2T-TT). This implies that P3HTT not only takes up less electrolyte (as deduced from the smaller change in frequency), but also appears less plasticized by the electrolyte solution.

Encouraged by our 'active-swelling' results, we went on and quantified the mass uptake in presence of an electrolyte solution and upon application of a +0.6 V bias, using multiple harmonics of the QCM resonance frequency and dissipation, and taking into account the energy losses during vibrations, assuming the same density for both polymers. Remarkably, while we infer from the high mass uptake of p(g2T-TT) an expansion in thickness by (+75 to +80) % in this scenario, we calculate the thickness of P3HHT films to only increase by less than +10 % (around +9 %; Figure 2c). Moreover, upon removal of the electrical bias, P3HHT almost fully restores its initial thickness as attested by the minor change in QCM-D frequency (∆f ≈ -10 Hz) between its initial state and after the bias is removed, which translates to a +1.8 % thickness increase. In contrast, p(g2T-TT) remains substantially swollen, around (+30 to +35) % of its initial thickness (∆f ≈ -110 Hz). This difference seems to predominantly result from water retention after the electric field is switched "OFF". Indeed, separating the contribution of anions and water by combining e-QCM-D data with simultaneously recorded electrical data (see SI Figure S13 for details), we identify the remaining thickness change of P3HHT after bias removal to be due to the fact that 20 % of overall water molecules are trapped in the films, but essentially no ions. In p(g2T-TT), 45 % of the water molecules remain in the system as well as some ions (Figure 2c, right panel), explaining why, after the electrical field is switched "OFF", this material keeps its mechanical softness, typical for plasticized material, while P3HHT mostly recovers its materials stiffness, as already alluded to above (Figure 2b).

Intriguingly, the different swelling behavior between P3HHT and p(g2-TT) cannot be explained entirely by the significantly higher ion injection into p(g2T-TT) at 0.6 V. Comparison of the total mass uptake indicates that, for P3HHT, each injected Cl -is accompanied by 7 water molecules, while for p(g2T-TT), it is 15 H 2 O molecules for each Cl -. At a lower bias of 0.2 V, where we approximately achieve the same accumulated charge in both polymer films (see Figure S14 in SI), the ratio of H 2 O molecules for each Cl -increases for p(g2T-TT) to 20. This experimental evidence demonstrates that p(g2T-TT) absorbs more water than P3HHT regardless of the bias that is applied, and it further highlights that p(g2T-TT) consistently retains more solvent molecules after the bias is removed. This observation may explain the electrochemical stability of P3HHT structures when subjected to repeated cycling over increasing voltage ranges up to 0.75 V (see Figure S15).

We hypothesize that this behavior is linked to the lower density of polar sites per repeat unit of P3HHT that can interact with water molecules, compared to those of p(g2T-TT). While water retention in polymers depends on many factors, water-polymer interactions can be expected to be one of the main driving forces. The framework of a proposed hierarchy of interactions (cation-anion > water-ion > water-polar = polar-polar = water-water > water-hydrophobic) [START_REF] Rowland | Water in Polymers[END_REF], thereby, helps to understand the concert of competing forces acting in these complex systems and the propensity of a material to absorb and retain water molecules.

Important for its usage in devices such as OECTs, the ion uptake during application of a positive bias results in chemical doping of P3HHT. We used electrochemical absorption spectroscopy [START_REF] Zhai | [END_REF]26], which is a commonly used method in the organic mixed conductor field [9,[14][15][16][17] for the purpose of tracking ion uptake/doping (details on this methodology are given in the Supplementary Information). Data obtained using an 0.1 mol/L KCl solution as electrolyte and a similar measurement geometry as employed for the e-QMC-D experiments, reveals bleaching of the ground-state absorption centered around 550 nm upon application of a 0.75 V-bias while, simultaneously, an increase of the polaron band (centered around 800 nm), characteristic for doped thiophene-based polymers, is observed (see Figure 3a). In agreement with our e-QMC-D data, upon removal of the bias, the same ground state absorption is restored (Figure 3b), indicating nearly complete de-doping, without the need to apply a negative bias. This favorable behavior, in addition to the fact that the OH-functionalization seems not to affect the hole mobility of materials as measured in a thin-film transistor (TFT) using P(3HT-3HHT) copolymers (see Supplementary Information Figure S16), has a direct impact on OECT device performance (Figure 3). For instance, the ease with which the material can be de-doped, as evidenced with our electrochemical spectroscopy measurements, leads to consistently low "OFF"-currents (see transfer characteristics displayed in Figure 3c). As a consequence, a high "ON"/"OFF" ratio of 10 4 , i.e., a pronounced difference between the "ON" and "OFF" current, is realized, which is on par/on the higher end of reported "ON"/"OFF" ratios reported in literature on p-type devices [9,[START_REF] Giovannitti | Proc. Natl. Acad. Sci[END_REF]21]. Very clean output characteristics were also recorded (Figure 3d). [Note: In case ion flow in-and out-of the polymers is very slow when compared to the gate voltage sweeping times, and/or ions remain in the active OECT layer even when the gate voltage is removed, a higher "OFF"-current would be observed, similar to TFTs when the active layer is slightly doped [27] and free charges remain present when the device is switched "OFF". Gate leakage can contribute to a high "OFF"-current as well, but we generally observed low gate currents, I g , in the 1 nA-regime. This effect was further limited by using in some cases insulated, patterned devices (see OECT section of the SI for details on device fabrication)] A large set of devices of different geometries (different channel lengths, L; channel widths, W; and thickness d; see Figure S17 in the SI for details), including two sets of insulated, patterned, small-channel devices with two different thicknesses, d (see Figure S18), and simple, unpatterned devices, allowed us to extract the figure of merit often used to compare OECT device performance [28], that is µ•C*, where µ is the OECT mobility and C* the volumetric capacitance, by plotting the transconductances, g m = , deduced from the transfer characteristics, against (V th is the threshold voltage, V g is the gate voltage). We obtain µ•C*(OECT) to be (35 ± 2) F•cm -1 •V -1 •s -1 (Figure 3e) which is only slightly lower to what is reported for PEDOT:PSS, (47 ± 2) F•cm -1 •V -1 •s -1 , but at the benefit of drastically lower volumetric expansion/retraction upon swelling.

The above discussion shows that P3HHT can be easily integrated in devices such as OECTs, leading to clean device characteristics that scale with device geometry (channel length, channel width; patterned vs. unpatterned) and come at par with devices fabricated, e.g., with the widely available PEDOT:PSS.

Patterned devices also allowed us to perform accurate bandwidths measurements (Figure 3f) to obtain the cut-off frequency of such OECTs. This enables estimation of their time response. We extract a cut-off frequency of 75 Hz. This is somewhat slower than what is measured for, e.g., p(g2T-TT). Yet, the lower swellability and, thus, the expected minimized structural changes to the film by water accumulation, could render P3HHT and its derivatives attractive for sensing applications, e.g., for tracking of biological signals [29,30]. Indeed, preliminary OECT stability measurements suggest that P3HHT-based devices are only minimally affected by immersion in water over multiple days (see Figure S19), contrary to commonly processed PEDOT:PSS [5].

In conclusion, despite the plethora of possible chemical structures that can be designed and synthesized, so far, focus has been on the ubiquitous PEDOT:PSS and its derivatives, as well as common semiconducting polymers used in the organic thin-film transistor field, substituted with polar side chains such as oligoethylene-glycol moieties. We showed here that hydroxyl-alkyl functionalization provides an interesting, alternative side-chain optionespecially if low swelling (i.e., small volumetric changes) upon doping is required. We attribute the low swelling behavior of P3HHT compared to glycolated organic mixed conductors to the fact that hydroxyl alkyl side-chains offers less sites available for water-polymer interactions, especially when compared to fully glycolated side-chain polymers such as p(g2T-TT). The OH-functionalization also has other benefits. It hardly affects the typical molecular packing; when compared to P3HT, only an increase in the lamellar stacking distance due to the slightly longer side-chain length is found (see structural characterization section in SI and GIWAXS data reported in Figure S20).

As a consequence, electronic charge transport as assessed in TFTs is not negatively impacted by the OHfunctionalization. Use of hydroxyl-alkyl side chains also leads to a higher melting temperature, T m , (290 °C; see Figure S21 for DSC thermogram) compared to, e.g., regio-regular P3HT of similar molar mass displays, T m  240 °C to 260 °C [31,32,33], promising a somewhat higher temperature stability. The benefits come at an expense. In order to minimize the hysteresis between forward and backwards scans when measuring OECT transfer characteristics, relatively slow sweeping rates of 0.25 V•s -1 (compared to p(g2T-TT) [START_REF] Giovannitti | Proc. Natl. Acad. Sci[END_REF]) are needed. In addition, the figure of merit of µ•C* (OECT) deduced for P3HHT, while rather comparable to PEDOT:PSS, is lower than for materials such as p(g2T-TT) for which a µ•C* (OECT) of (240±90) F•cm -1 •V -1 •s -1 is reported [28]. Yet, the large error of the µ•C* (OECT) values reported for p(g2T-TT) indicates potential issues with accurately determining this figure of merit for this material, possibly resulting from, among other things, its variations in glass transition temperature (due to, e.g., different plasticizing degrees, which will affect mass transport) and its generally strong dependence of properties on the environment. Clearly the use of hydroxyl-alkyl functionalization will expand future materials design options, especially when such side chains are combined with backbones with a better electronic charge-transport capability, such as the one used in p(g2T-TT) [34]. Considering that the synthetic pathway is relatively straight-forward, requiring only a protection/de-protection step and enabling synthesis of highly regioregular materials because the oxygen atom is well distanced from the aromatic moiety of the monomer and, thus, does not affect its reactivity during polymerization, our work promises an alternative chemical design tool towards more reliable mixed conducting structures. This is achieved via materials that undergo lower volumetric changes upon doping/de-doping cycles (as suggested by the high electrochemical stability of P3HHT structures when subjected to repeated cycling over increasing voltage ranges up to 0.75 V that we observe; see Figure S15), that display a smaller environmental sensitivity and feature higher thermal transition temperatures, while hole transport seems not to be affected, as thin-film transistor measurements show (Figure S16), likely because the OHfunctionalization does not affect molecular packing nor the edge-on orientation often found for thiophenebased polymers [35], as already alluded to above. We thus conclude that hydroxyl-alkyl functionalization leads to model systems that can assist to systematically establish relevant structure/property interrelations with respect to the fundamental requirements that organic materials need to fulfill to sustain ion flow/electrical conduction without adverse side effects such as excessive swelling.

  volumetric changes can have undesirable consequences, e.g., negatively affecting hole/electron conduction and severely shortening device lifetime. Here, we present a new material poly[3-(6hydroxy)hexylthiophene] (P3HHT) able to transport ions and electrons/holes, as tested in electrochemical absorption spectroscopy and organic electro-chemical transistors, that exhibits low swelling, attributed to the hydroxylated alkyl side-chain functionalization. P3HHT displays a thickness change upon passive swelling of only +2.5 %, compared to +90 % observed for the ubiquitous PEDOT:PSS, and (+10 to +15) % for polymers such as poly(2-(3,3'-bis(2-(2-(2methoxyethoxy)ethoxy)ethoxy)-[2,2'-bithiophen]-5-yl)thieno[3,2-b]thiophene) (p(g2T-TT)). Applying a bias pulse during swelling, this discrepancy becomes even more pronounced, with the thickness of P3HHT films changing by <10 % while the one of p(g2T-TT) structures increases by (+75 to +80) %.

Figure 1 a

 1 Figure 1 a) Chemical structures of P3HHT, left, and p(g2T-TT), right. b) 7 th Harmonic of QCM-D data obtained on P3HHT films with respect to a bare gold sensor reference. Passive swelling leads to negligible differences between the dry and wet state of a 90 nm P3HHT film (bold lines) while QCM-D data for a 50 nm film of p(g2T-TT) (thin lines) shows a more noticeable increase in frequency variation. c) Sauerbrey-derived thickness from the 7 th harmonic of the QCM-D for P3HHT in the wet and dry state (solid and dotted lines, respectively; further details are given in the SI).

Figure 3 a

 3 Figure 3 a) Time evolution of P3HHT absorption recorded upon application of 0.75 V bias. b) Absorption intensity at 560 nm normalized at t = 0 s versus time. c) Transfer loop characteristics of an unpatterned device (W = 100 µm, L = 10 µm) at varying sweep rates. d) Forward output characteristics of the same device using a V d -scan rate of 0.5 V/s, while the gate voltage varied between (+0.1 to -0.7) V with voltage steps of 0.05 V. e) Scaling of transconductance, g m , at V g = 0.6 V with channel dimensions using data obtained from simple, unpatterned devices (black squares) and patterned, insulated devices (red, empty circles; see details of channel dimensions and devices in SI). f) Bandwidth measurement of a patterned OECT (W = 100 µm, L = 10 µm) at V d = -0.6 V showing the half-power point to be at 75 Hz.

Acknowledgements

The authors thank Jonathan Rivnay and Sahika Inal for highly fruitful discussion and preliminary OECT measurements. Tommaso Nicolini, Georges Hadziioannou and Natalie Stingelin are grateful for financial support provided by the IONBIKE RISE project, which has received funding from the European Union's Horizon 2020 research and innovation programme under the Marie Skłodowska-Curie grant agreement No 823989. Tommaso Nicolini, Georges Hadziioannou and Natalie Stingelin also acknowledge funding from the MARBLE project (IdEX). Olivier Dautel acknowledges funding from the MAPLE project (Institut Carnot Chimie Balard Cirimat). Lee J. Richter and Tommaso Nicolini thank NSLS-II for the access to the CMS BM11 beamline of the Brookhaven National Laboratory NY, USA which is a U.S. DOE Office of Science Facilities, at Brookhaven National Laboratory under Contract No. DE-SC0012704.

Declaration of conflict of interest

The authors declare no conflict of interest.