Phase Behavior and Conductivity of Sulfonated Block Copolymers

Oct 24, 2012 - PMB) block copolymers by varying kinds of heterocyclic diazoles in ILs. A low molecular weight PSS-b-PMB copolymer (3.5−3.1 kg/mol) w...
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Phase Behavior and Conductivity of Sulfonated Block Copolymers Containing Heterocyclic Diazole-Based Ionic Liquids Onnuri Kim,† Sung Yeon Kim,† Hyungmin Ahn,‡ Chang Woo Kim,‡ Young Min Rhee,‡ and Moon Jeong Park†,‡,* †

Division of Advanced Materials Science and ‡Department of Chemistry, Pohang University of Science and Technology (POSTECH), Pohang, Korea 790-784 S Supporting Information *

ABSTRACT: We have investigated morphologies and conductivities of ionic liquids (ILs) incorporated poly(styrenesulfonate-b-methylbutylene) (PSS-bPMB) block copolymers by varying kinds of heterocyclic diazoles in ILs. A low molecular weight PSS-b-PMB copolymer (3.5−3.1 kg/mol) with sulfonation level of 17 mol % was employed as a matrix polymer, which indicates disordered morphology at entire temperature examined. The addition of different ILs results in the emergence of various ordered morphologies such as lamellar, hexagonal cylinder, and gyroid structures. Interestingly, it has been revealed that ring structures and alkyl substituents in diazoles play an important role in determining the morphologies of ILs impregnated PSS-b-PMB copolymers, attributed to the dissimilar strength of ionic interaction. Heating the ILs doped PSS-b-PMB copolymers causes intriguing thermoreversible order−order and order−disorder phase transitions, which can be rationalized by classical block copolymer thermodynamics. From conductivity measurements, it has been found that the enhanced conductivity could be achieved by increasing number of protic sites in heterocyclic diazoles. Upon exploring morphology effects on conductivities of ILs-containing PSS-b-PMB copolymers, with decoupled segmental motion of polymer chains and ion transport, similar morphology factor of 0.4 has determined if the morphologies are appeared to be lamellar and/or hexagonal cylinder structures. In contrast, the gyroid-forming sample revealed apparently high morphology factor in the range of 0.6 to 0.7, which is intimately related to better connectivity of ionic channels along cocontinuous PSS phases.

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INTRODUCTION

designs of PEMs for the access of desired ion transport properties.22,28,29 Systematic studies were carried out to underpin the factors governing ion transport and thermodynamic properties of the [CnIm][X]-incorporated PEMs, which have drawn following issues: (1) The local concentration of ions is revealed as an important parameter in enhancing the conductivity of PEMs. The ionic conductivity of IL doped PEMs becomes higher when the IL concentration in conducting domains is increased, as universally observed for diverse combinations of ILs and PEMs.9,16,29,30 (2) The length of alkyl groups in [CnIm] cations largely affects the ionic conductivity of PEMs. Pioneering studies of Colby et al. demonstrated that the enhanced ionic conductivity could be achieved by increasing the length of alkyl group (n value) of [CnIm], attributed to the lowering of Tg of PEMs.31,32 (3) The anions of ILs play an important role in determining the conductivity of PEMs, as reported by Elabd et al.18 and Park et al.9,26 This is due to the change in polarities of ILs and dissimilar diffusion coefficients of ions within PEMs, depending on the types of anions.15,26,33,34 (4) For a number of

Recently, great attention has been paid to ionic liquids (ILs) integrated polymer electrolyte membranes (PEMs) for a variety of applications such as electro-active actuators,1−3 lithium batteries,4−6 and high temperature PEM fuel cells.7−10 The tide was triggered by fascinating physicochemical properties of ILs, i.e., negligible vapor pressure, thermal stability, and large electrochemical stability window.11−14 It has been revealed that the ILs make a large alteration in the thermodynamic and electrochemical properties of the PEMs since the degree of ion dissociation,10,15 local concentration of ions,16 and glass transition temperature (Tg) of the membranes17,18 can be manipulated. Among a range of ILs, the ILs comprising 1,3-alkylimidazole, [CnIm] with n = 2−11, are most widely employed.16,19 Different polymers have been served as matrix materials for the incorporation of [CnIm][X] (X represents diverse counterions). Examples of polymers include poly(vinylidene fluoride) (PVDF),16,19 poly(ethylene oxide) (PEO),20,21 poly(methyl methacrylate) (PMMA),22,23 poly(vinylpyridine) (PVP),24,25 and poly(styrenesulfonate) (PSS).26,27 Many attempts to synthesize random, graft, or block copolymers composed of above-mentioned polymers have concerned with the optimal © 2012 American Chemical Society

Received: August 28, 2012 Revised: October 10, 2012 Published: October 24, 2012 8702

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styrene (after the reaction) to total moles of styrene (before the reaction), as determined by 1H NMR spectra with acetone-d6. The molecular structure of resulting PSS-b-PMB copolymer (3.5−3.1 kg/ mol) is shown in Figure 1a.

PEMs, it has been reported that phase-separated structures have benefits for enhancing ion transport rate.9,20,24 The results have been rationalized by the fact that well-defined phase boundaries between ion-containing domains and nonconducting phases could create less tortuous pathway for ion transport by constructing confined spaces. Such intriguing observations of morphology dependent conductivity have been a motivation for the molecular-level understanding of morphology-transport relationship of PEMs.26,35,36 While the PEMs containing 1,3-alkylimidazole have been extensively examined, only limited attention has been devoted to the PEMs composed of other heterocyclic diazoles with different ring structures.32,37−41 Accordingly, the fundamental understanding of the thermodynamics and ion transport behavior of PEMs comprising other heterocyclic diazoles is lacking substantially. This is quite surprising since it is not yet clear whether the 1,3-alkylimidazole-based ILs could be successfully integrated into real applications. For example, for high temperature PEM fuel cells, sufficient numbers of protons in the PEMs are required, however, a vast body of 1,3alkylimidazole explored to date does not possess any protic sites.18,31,32,37,40 Herein, we are motivated to explore the phase behavior and conductivity of PEMs upon incorporating a range of ILs containing different heterocyclic diazole cations. We hope to bring up following issues; (1) Can we create well-defined nanostructures for the PEMs comprising other heterocyclic diazole-based ILs and what would be the effects of nanoscale morphology on conductivity? (2) Is imidazole the best choice of cation among different heterocyclic diazoles to obtain the utmost conductivity? (3) Are Brönstead-type cations containing many protic sites beneficial in achieving the higher conductivity? In this paper, as a model PEM, we particularly focused on a low molecular weight poly(styrenesulfonate-b-methylbutylene) (PSS-b-PMB) block copolymer (3.5−3.1 kg/mol, sulfonation level =17 mol %), which exhibits a disordered morphology at entire temperature window examined. For this polymer, acquiring the equilibrium morphology in the presence of ILs would be rather easy, which helps better understanding of morphology-property relationship. In present study, we found that regardless of kinds of diazoles, the addition of ILs into the PSS-b-PMB copolymer causes the enhanced strength of phase separation (increase in Flory−Huggins interaction parameter, χ) while the relative degree of the increments was sensitive to the types of diazoles. The links between morphology and conductivity for ILs integrated PSS-b-PMBs were examined with decoupled segmental motion of polymer chains and ion transport.

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Figure 1. Molecular structures of (a) poly(styrenesulfonate-bmethylbutylene), S30MB44(17), copolymer, (b) imidazole, [Im], (c) pyrazole, [Py], (d) 1-methylimidazole, [1-MIm], (e) 2-methylimidazole, [2-MIm], (f) 2-ethyl-4-methylimidazole, [2-E-4-MIm], and (g) bis(trifluoromethane)sulfonimide, [HTFSI]. Ionic Liquids (ILs). Imidazole ([Im], ≥99.5%), pyrazole ([Py], ≥98.0%), 1-methylimidazole ([1-MIm], ≥99.0%), 2-methylimidazole ([2-MIm], ≥99.0%), 2-ethyl-4-methylimidazole ([2-E-4-MIm], ≥95.0%), and bis(trifluoromethane)sulfonimide ([HTFSI], ≥ 95%) were purchased from Sigma-Aldrich and used as received. The chemical structures of the diazoles and HTFSI are shown in Figure 1b−g. Five different nonstoichiometric ILs were synthesized by combining diazole compounds and HTFSI in a 2:1 molar ratio at 180 °C. The final composition of the synthesized ILs was 2:1, as determined by Fourier transform infrared (FT-IR) experiments. Preparation of IL Incorporated PSS-b-PMB Membranes. Inhibitor-free anhydrous tetrahydrofuran (THF, ≥ 99.9%) was used without further purification and methanol was degassed 2 times prior to use. Equimolar amounts of nonstoichiometric IL to the mole of [−SO3H] in PSS-b-PMB copolymer were added. Namely, the mole ratio of [diazole cation]:[TFSI anion]:[−SO3H] in PSS-b-PMB copolymer was 2:1:1. The 5 wt % solutions were prepared using 50/50 vol % THF and methanol mixtures, stirred overnight at room temperature. Membranes were prepared by solvent casting under an argon atmosphere for 2 days followed by vacuum drying at 70 °C for 7 days. To exclude the issue of water contamination of hygroscopic samples, sample preparations and measurements were performed under an Ar-filled glovebox with oxygen and moisture concentration below 0.1 ppm. Small Angle X-ray Scattering (SAXS). The IL-containing membranes were laminated into an airtight sample cell, which consists of an aluminum spacer, two Kapton windows, O-rings, and aluminum covers. Synchrotron SAXS measurements on these samples were performed using the 4C SAXS beamline at the Pohang light source (PLS). Sample temperature was controlled within ±0.2 °C using a sample stage provided by the PLS. The wavelength (λ) of the incident X-ray beam was 0.15 nm (Δλ/λ = 10−4), and sample-to-detector distance of 1.5 m was used yielding scattering wave vector q (q = 4π[sin(θ/2)]/λ, where θ is the scattering angle) in the range 0.1−2.0 nm−1. The scattering data were corrected for the CCD dark current and the scattering from air and Kapton windows. The resulting twodimensional scattering data were averaged azimuthally to obtain intensity versus q. Transmission Electron Microscopy (TEM). The IL-incorporated PSS-b-PMB membranes were cryo-microtomed at −120 °C to obtain thin sections with thicknesses in the 80−120 nm range using a RMC Boeckeler PT XL Ultramicrotome. The electron contrast in the

EXPERIMENTAL SECTION

Synthesis of PSS-b-PMB Copolymer. A low molecular weight poly(styrene-b-methylbutylene) (PS-b-PMB) block copolymer (3.1− 3.1 kg/mol) was synthesized by sequential anionic polymerization of styrene and isoprene, followed by selective hydrogenation of the polyisoprene.9 The molecular weight and molecular weight distribution of the PS-b-PMB copolymer were characterized by combining 1H nuclear magnetic resonance (1H NMR, Bruker AVB-300) spectroscopy and gel permeation chromatography (GPC, Waters Breeze 2 HPLC). The polydispersity index of the hydrogenated polymer was 1.03, measured on GPC with polystyrene standards in tetrahydrofuran (THF) for calibration. The styrene block of PS-b-PMB copolymer was then sulfonated using procedures described in ref 9. A sulfonation level (SL) of 17 mol % was obtained from the ratio of moles of sulfonated 8703

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Table 1. Molecular Properties of Heterocyclic Diazoles and HTFSI Used in the Present Study materials used

a

molecular weight (g/mol)

melting temperature (Tm, °C)a

van der Waals volume (Å3)b

pKa c

imidazole, [Im]

68.1

90

7.0

pyrazole, [Py]

68.1

68

2.5d

1-methylimidazole, [1-MIm]

82.1

−6

7.2c

2-methylimidazole, [2-MIm]

82.1

142

7.9c

2-ethyl-4-methyl imidazole, [2-E-4-Mim]

110.2

54

8.7c

bis(trifluoromethane) sulfonimide, [HTFSI]

281.2

60

−10e

61.6 58.6 61.4 63.6 78.9 77.4 78.2 75.6 111.9 109.2 159.2

(cation) (cation) (cation) (cation) (cation)

Measured using DSC. bCalculated using ChemAxon’s Calculator Plugins. cReference 42. dReference 43. eReference 44.

samples was enhanced by exposure to ruthenium tetroxide (RuO4) vapor for 50 min. Imaging of stained samples was performed with a JEOL JEM-2100F microscope operating at 200 kV equipped with a cold stage (−160 °C) and an Omega energy filter. Images were recorded on a Gatan 2048 × 2048 pixel CCD camera. (Gatan Inc., Pleasanton, CA). All data sets were acquired using Digital Micrograph (Gatan, Inc.) software. Conductivity Measurements. The conductivities of ILs-incorporated PSS-b-PMB membranes were measured using AC impedance spectroscopy in a glovebox. The through-plane conductivity was measured using a home-built two-electrode cell with 1.25 cm × 1.25 cm stainless steel blocking electrodes, Kapton spacers, and 1 cm × 1 cm platinum plates. The platinum plates were used as working and counter electrodes to apply a current to the membranes. The counter electrode was engraved with 0.8 cm × 0.8 cm hole and the samples (0.8 cm × 0.8 cm, 500 μm thick) were sandwiched between two platinum plates in the presence of spacers. Schematic drawing illustrating the conductivity cell is provided in Supporting Information. Data were collected using a 1260 Solatron impedance analyzer operating over a frequency range of 1−100,000 Hz. Differential Scanning Calorimetry (DSC). The glass transition temperature (Tg) of IL-containing PSS-b-PMB copolymer was measured by DSC using a Seiko Instruments (model DSC-220cu). The samples were encapsulated in aluminum pans (SSC000E33, Seiko Instruments Inc.) within glovebox prior to the measurements to exclude the issue of water contamination of hygroscopic samples. A fixed heating and cooling scan rate of 5 °C/min was employed and Tg was obtained from the change in heat capacity during the second heating. All measurements were performed with sample mass in the range of 8 to 10 mg. An empty pan was used as a reference.

[2-E-4-MIm]), and [HTFSI] (g). Table 1 summarizes the molecular properties of heterocyclic diazoles and HTFSI used in present study. Interestingly, a pair of diazoles with slightly different cyclic structures but having an identical molecular weight, i.e., [Im] vs [Py] and [1-MIm] vs [2-MIm], indicates considerably different melting temperatures (Tm). This signals dissimilar molecular interactions among diazoles depending on the ring structures. The pKa values and the van der Waals volumes of diazoles and HTFSI are also given in Table 1. The substantially low pKa value of [Py], compared to those of other diazoles is worth noting. The morphologies of IL integrated S30MB44(17) copolymers were examined by combining SAXS and TEM experiments. We first investigate the morphologies obtained with the ILs comprising unsubstituted diazoles, i.e., [Im][TFSI] and [Py][TFSI]. Figure 2 shows the SAXS profiles of [Im][TFSI] and [Py][TFSI]-incorporated S30MB44(17) copolymers, compared with that of neat S30MB44(17), measured at 25 °C. We see one broad peak centered at q = 0.71 nm−1 for the neat S30MB44(17) copolymer, signaling disordered morphology. Note in passing that the segregation strength of S30MB44 copolymer is a sensitive function of SL and the SL of 17 mol % is not yet sufficient to yield ordered morphology. (details on the SLdependent morphologies of S30MB44 copolymers are shown in Supporting Information). In contrast, the incorporation of ILs into the disordered S30MB44(17) results in ordered morphologies. With the [Py][TFSI], Bragg peaks (∇) at 1q*:2q* with q* = 2π/d100 and d100 = 9.2 nm were seen, indicating the formation of a lamellar structure (LAM). In contrast, the use of [Im][TFSI] results in a hexagonal cylinder morphology (HEX), as shown by Bragg peaks (▼) at 1 q*:√3q*:√4q*:√7q*√9q* with q* = 2π/d100 and d100 = 11.8 nm. The disorder, LAM, and HEX morphology of each sample was confirmed by cross-sectional TEM experiments on the cryo-microtomed membranes, as shown in the inset TEM images of Figure 2. PSS phases were stained by RuO4. The domain spacings of [Py][TFSI] and [Im][TFSI]-incorporated samples, obtained from TEM were 8.9 and 11.5 nm, respectively, which are in good agreement with the SAXS results. It should be noted here that the increment of domain size for [Im][TFSI]-incorporated S30MB44(17) copolymer is quite larger than the value obtained with [Py][TFSI], as shown in the inset plot of Figure 2. Such dissimilar domain spacings are intriguing since the van der Waals volume of [Im] is almost identical to that of [Py], as listed in Table 1, on account of similar 5-membered ring structures. This implies that the

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RESULTS The Morphologies of ILs-Containing PSS-b-PMB Copolymers. Figure 1a shows the molecular structure of a PSS-b-PMB (3.5-b-3.1 kg/mol) copolymer with a fixed sulfonation level (SL) of 17 mol % (five styrene units were sulfonated among 30 styrene chains), as determined from 1H NMR experiments. Hereafter the copolymer is referred to as S30MB44(17) where the subscripts indicate the degree of polymerization of each blocks. The SL value of the copolymer is given in parentheses. We aimed to investigate the morphologies and conductivities of S30MB44(17) copolymer upon incorporating ILs comprising different heterocyclic diazoles. The counterion of ILs is fixed at bis(trifluoromethane) sulfonimide ([TFSI−]), offering reasonably low Tg values of the composite membranes. In Figure 1b− g, we show the chemical structures of heterocyclic diazoles and HTFSI in their nonionic forms; imidazole (b, hereafter [Im]), pyrazole (c, [Py]), 1-methylimidazole (d, [1-MIm]), 2methylimidazole (e, [2-MIm]), 2-ethyl-4-methylimidazole (f, 8704

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Figure 2. SAXS profiles of [Im][TFSI]- and [Py][TFSI]-incorporated S30MB44(17) copolymers measured at 25 °C, as compared with that of neat S30MB44(17). The scattering profiles are vertically offset for clarity. The inverted open triangles (∇) of [Py][TFSI] integrated copolymer and the inverted filled triangles (▼) of [Im][TFSI] embedded sample indicate Bragg peaks at q*, 2q*; and at q*,√3q*,√4q*,√7q*,√9q*; respectively. The arrow (↓) for neat copolymer represents disordered morphology. Cross-sectional TEM images obtained in the absence/presence of ILs confirm the disorder, LAM, and HEX phases. SS domains were stained by RuO4 and the scale bars represent 40 nm. The domain size of each sample is plotted in the inset figure.

Figure 3. SAXS profiles of [Im][TFSI]-incorporated S30MB44(17) copolymer measured as a function of temperature. The scattering profiles are vertically offset for clarity. The arrows (↓) at 60 °C, the inverted open triangles (∇) at 90 °C, and the inverted filled triangles (▼) at 130 °C indicate Bragg peaks at q*,√3q*,√4q*,√7q*,√9q*. The inverted thick arrows (↓) at 90 °C, 130 °C, and 170 °C show Bragg peaks at q*, 2q*, indicative of the emergence of LAM phase. The dotted line is obtained during cooling scan, confirming thermoreversible characteristics of the observed phase transition via coexistence of LAM and HEX phases. The temperature-dependent domain sizes of two structures are illustrated in the inset plot.

dissimilar molecular-assemblies between diazoles and [−SO3H] groups and/or conformational changes of PSS block were induced by varying kinds of diazoles, eventually affecting the mean curvature of ionic SS domains. We also take into account the different selectivity of ILs to the PSS phase. For example, [Im][TFSI] is more selective to PSS block than [Py][TFSI] so that the chain stretching to reduce the interfacial area becomes significant for the [Im][TFSI]-incorporated sample. This bears resemblance to the well-established phase behavior of nonionic block copolymer/selective solvent mixtures,45 although the balance of interfacial area and chain stretching in the presence of ionic additives becomes more complicated. When the IL-incorporated S30MB44(17) copolymers are exposed to different temperatures, even more striking differences were observed. Figure 3 shows the SAXS profiles for [Im][TFSI] integrated S30MB44(17) as a function of temperature. At 60 °C, the sample shows qualitatively similar SAXS profile to that of 25 °C, indicative of HEX phase. At 90 °C, we find that the scattering peaks representing the HEX phase (identified in Figure 3 by ∇) persist but new peaks corresponding to a LAM phase emerge (identified by ↓ in Figure 3). Upon further heating, the relative scattering intensity of LAM over HEX morphology gradually increases and eventually, the HEX phase is no longer evident and only the LAM phase is seen at 170 °C. This indicates that the stability window of the HEX morphology becomes narrower and eventually disappears with increase in temperature and therefore, the sample undergoes order−order phase transition

(OOT) from HEX to LAM structures. The onset temperature of the OOT was found at 75 ± 5 °C. It is interesting to note that the HEX and LAM phases of [Im][TFSI]-containing S30MB44(17) coexist for quite wide temperature window from 80 to 160 °C. Such coexistence was thermally reversible, as confirmed by cooling the sample from 190 °C. The SAXS profiles obtained with cooling scan are shown by dotted lines in Figure 3. The reemergence of the HEX peaks was evident with almost identical scattering intensities to those of heating scan. This suggests that the LAM+HEX coexistence may be an equilibrium property of [Im][TFSI] embedded S30MB44(17). It should be noted here that the sample was equilibrated for 30 min at each temperature before measurements and further increase in annealing time would change the temperature window of coexistence. At the coexistence states, the domain size of the HEX phase is 7−9% larger than that of the LAM structure. For example, at 130 °C, the domain size of the HEX is 11.5 nm while that of the LAM is 10.7 nm. The temperature dependent domain sizes of the LAM and HEX structures are shown in the inset plot of Figure 3. This result is analogous to the phase behavior of nonionic block copolymer/selective solvent mixtures. The pioneering work of Lodge et al.45 demonstrated the HEX to LAM phase transition with the increase in temperature for a set of poly(styrene-b-isoprene) (PS-b-PI) block copolymers/PS− 8705

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To underpin the origin of different phase behavior of [Im][TFSI] integrated S30MB44(17) and [Py][TFSI] loaded sample, the strength and conformation of ionic interaction taking place in SS phases were calculated by ab initio calculation using density functional theory based on the B3LYP exchangecorrelation functional.46 In Table 2, we have presented 0 K

or PI−selective solvent mixtures, which was accompanied by the decrease in domain spacing. Central to these phenomena was the strong temperature dependence of the solvent selectivity toward the one of block. We thus infer that the OOT of [Im][TFSI]-containing S30MB44(17) copolymer is taken place with decreasing volume fraction of hydrophilic SS phases upon heating due to the decrease in the selectivity of [Im][TFSI] to the SS phases. For [Py][TFSI]-incorporated S30MB44(17) copolymer, it is found that the LAM morphology becomes disordered with the increase in temperature, as shown in Figure 4. The order−

Table 2. Ab Initio Interaction Energies at 0 K in a Vacuum for Different Heterocyclic Diazoles with −SO3H and HTFSI Counterions binding energy (kJ/mol)a materials

ionic pair

nonionic pair

[Im][−SO3H] [Im][HTFSI] [Py][−SO3H] [Py][HTFSI] [1-MIm][−SO3H] [1-MIm][HTFSI] [2-MIm][−SO3H] [2-MIm][HTFSI] [2-E-4-MIm][−SO3H] [2-E-4-MIm][HTFSI]

459.3 378.4 514.1 432.8 439.4 378.9 430.4 375.4 419.3 361.3

42.3 42.9 83.1 73.3 79.0 79.8 75.4 77.2 76.2 79.6

a

All calculations were performed using a DFT Exchange-Correlation Functional, B3LYP.

binding energies of ion pairs and nonion pairs between diazoles and counterions. It has been found that TFSI− anions weakly bind to diazolium cations, compared with [−SO3−] groups. The binding energies among neutral species become an order of magnitude lower than the values among ionic ones. It is worthwhile to note here that a formation of strong ionic bond is expected for [Py]-containing copolymer despite the lower nucleophilicity in [Py] than that in [Im] (see pKa values given in Table 1). Figure 5 shows molecular schemes illustrating the ionic bond formation of [Im] and [Py] in the presence of [−SO3−] and [TFSI−] counterions. The adjacent nitrogen atoms of [Im] and [Py] can serve as donors and acceptors of protons, however, the calculation suggests that both nitrogen atoms of [Py] are engaged in an ion pair formation with sulfonic acid group (and/ or TFSI−). We can infer that the long-range ionic association within SS phases would be likely reduced for the [Py]containing copolymer since it requires reorientation of the [Py] molecule, which is regarded as a rate-determining step.41,43 This would be, in part, the reason why relatively weakly segregated ionic SS phases and hydrophobic MB domains are caused for the [Py][TFSI] embedded S30MB44(17) copolymer. When the samples are exposed to elevated temperatures, weakened interactions among ionic moieties are anticipated for both [Im][TFSI] containing copolymer and [Py][TFSI] loaded sample. Particularly, a large increase in diffusion coefficient is anticipated for [Py] in SS phases on account of the lower melting temperature of [Py] (see Table 1), attributed to the emergence of disordered morphology of [Py][TFSI] embedded copolymer at high temperatures. Note that the clustering of numbers of diazole molecules via hydrogenbonding is not considered in presence study based on the assumption of evenly distributed ILs within SS phases at the equimolar IL loadings. Then, what would be the effect of alkyl-substitution of the heterocyclic diazoles on the morphological transitions of ILsincorporated S30MB44(17) copolymers? Two different [Im]

Figure 4. SAXS profiles of [Py][TFSI]-incorporated S30MB44(17) copolymer measured as a function of temperature. The scattering profiles are vertically offset for clarity. The arrows (↓) at 60 °C, the inverted open triangles (∇) at 90 °C, and the inverted filled triangles (▼) at 130 °C show Bragg peaks at q*,2q*, indicative of LAM morphology. The inverted thick arrows (↓) centered at broad scattering peak of 135 °C present disordered morphology. The dotted line is obtained during cooling scan from 150 °C, confirming thermoreversible characteristics of the observed ODT. The full-width at half-maximum (fwhm) of the primary peak as a function of temperature is plotted in the inset figure.

disorder transition (ODT) temperature can be located at 132.5 ± 2.5 °C, as viewed by large increase in the full-width at halfmaximum (fwhm) of the primary peak (see the inset figure of Figure 4). The dotted line is obtained during cooling scan from 150 °C to confirm thermoreversible characteristics of the observed ODT. Without doubt, this suggests that heating the [Py][TFSI]-containing S30MB44(17) copolymer causes weakly segregated ionic SS phase and nonionic MB domains, in turns, the decrease in effective χ parameter (χeff). The ODT behavior can be mapped on to the well-established phase diagrams of nonionic block copolymer/additive systems45 although direct comparisons between conventional nonionic block copolymers and our system may not be appropriate. 8706

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Figure 5. Molecular schemes illustrating the ionic bond formation of [Im] and [Py] in the presence of [−SO3−] and TFSI− counterions; (a) [Im][−SO3H], (b) [Im][HTFSI], (c) [Py][−SO3H], and (d) [Py][HTFSI].

derivatives are first examined where methyl group is substituted in 1 position ([1-MIm], Figure 1d) or 2 position ([2-MIm], Figure 1e). The effects of alkyl-substituent on the pKa values of heterocyclic diazoles are known to be roughly parallel to those described for parent diazoles, as listed in Table 1. The methyl substituent in [Im] causes a small but noticeable reduction in the strength of ionic interaction between the methyl imidazole and the sulfonic acid group while the degree of the weakening becomes nonsignificant for pairs of the methyl imidazole and HTFSI (see Table 2). The position of methyl-substitution does not appear to be crucial in affecting the ion pairing energies. The van der Waals volumes of [1-Mlm] and [2-MIm] are also estimated to be qualitatively the same, as shown in Table 1, analogous to the case of [Im] and [Py]. Figure 6 shows the representative SAXS data for S30MB44(17) copolymer with [1-MIm][TFSI] and [2-MIm][TFSI] loadings measured at different temperatures. Ordered morphologies were again seen with the incorporation of methyl-substituted imidazoles while the kinds of self-assembled structures are distinctly different, depending on the position of methyl group. At 60 °C, as shown in Figure 6a, the addition of [1-MIm][TFSI] into S30MB44(17) copolymer leads to the formation of a HEX structure with domain size of 12.8 nm. The HEX morphology of [1-MIm][TFSI] containing copolymer at low temperature bears resemblance to that of [Im][TFSI]incorporated sample while the domains size obtained with [Im][TFSI] was relatively small as 11.8 nm (Figure 3). It is intriguing to note that unlike the [Im][TFSI] integrated S30MB44(17), virtually no difference in the HEX morphology was detected for [1-MIm][TFSI]-incorporated copolymer in the temperature range of 60 to 180 °C, although heating the sample causes gradual decrease in domain size from 12.8 to 12.3 nm. We thus infer that χeff of [1-MIm][TFSI] embedded S30MB44(17) copolymer is enhanced by attaching methyl group at the nitrogen atom of the [Im] so that no accessible phase transition (OOT or ODT) is detected up to 180 °C. More quantitative analysis of the χeff parameters of ILs-containing PSS-b-PMB copolymers is remained as future work. In contrast, when the methyl-substitution is conducted aside the nitrogen atoms of the [Im], i.e., the 2 position of the ring structure, strikingly different phase behavior was observed. This is intriguing since the strength of ionic interaction among ion

Figure 6. SAXS profiles of (a) [1-MIm][TFSI] and (b) [2MIm][TFSI]-incorporated S30MB44(17) copolymers by varying temperature. The scattering profiles are vertically offset for clarity. (a) The inverted thick arrows (↓) indicate Bragg peaks at q*,√3q*,√4q*, indicating thermally stable HEX morphology. (b) The inverted arrows (↓, 60 °C), the inverted open triangles (∇, 90 °C), and the inverted filled triangles (▼, 135 °C) indicate LAM Bragg peaks at q*, 2q*. The inverted thick arrow (↓) centered at broad main peak of scattering intensity measured at 140 °C represents disordered morphology. The dotted line is obtained from cooling scan, confirming thermoreversible characteristics of the observed ODT. The schematic illustration of each nanostructure is shown in the inset graphic.

pairs obtained with [2-MIm] was not marked dissimilar to the value determined with [1-MIm] (see Table 2). As shown in Figure 6b, for the [2-MIm][TFSI] integrated copolymer, a LAM structure was seen with the q*:2q* Bragg peaks (↓) at 60 °C with q* = 2π/d100 and d100 = 9.5 nm. The domain spacing of 9.5 nm is significantly smaller than the value of 12.8 nm obtained with [1-MIm][TFSI], despite the similar van der Waals volumes of [1-MIm] and [2-MIm], as listed in Table 1. In addition, heating the LAM forming [2-MIm][TFSI]/ S30MB44(17) copolymer causes thermally reversible LAM-todisorder transition. The dotted line in Figure 6b obtained with cooling scan confirms the thermoreversible phase transition. The ODT temperature can be located at 137.5 ± 2.5 °C. Without doubt, this leads us to conclude that the position of alkyl-substitution makes a large alteration in the thermodynamic properties of IL-containing block copolymers. We now examined the phase behavior of IL-containing S30MB44(17) copolymer by tethering two alkyl groups to the nonprotic sites of [Im]. Figure 7 represents the temperaturedependent SAXS profiles of [2-E-4-MIm][TFSI]-incorporated S30MB44(17) copolymer. At 60 °C, the sample exhibits classical signatures of the HEX phase, as viewed by Bragg peaks (↓) at 1q*:√3q*:√4q* with q* = 2π/d100 and d100 = 11.2 nm. The increment in domain size now becomes substantial unlike the small value obtained with [2-MIm] (9.5 nm), as rationalized by significantly larger van der Waals volume of 111.9 Å3 for [2-E4-MIm] than 78.2 Å3 of [2-MIm]. Note here that, however, the domain size of 11.2 nm is still smaller than 12.8 nm of [18707

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are illustrated in Figure 8. The letters D, L, H, and G indicate disorder, lamellae, hexagonal cylinder, and gyroid morphology,

Figure 8. Phase diagrams of different heterocyclic diazole-based ILsincorporated S30MB44(17) copolymers as a function of temperature and kinds of ILs. The phase boundaries are marked with dotted lines. D, L, H, and G indicate disorder, lamellae, hexagonal cylinder, and gyroid morphology, respectively. The coexistence phases of two different ordered morphologies upon undergoing OOT are shown by shadows.

Figure 7. SAXS profiles of [2-E-4-MIm][TFSI]-incorporated S30MB44(17) copolymer measured as a function of temperature. The scattering profiles are vertically offset for clarity. The arrows (↓) at 60 °C and the thick arrows (↓) at 180 °C indicate Bragg peaks at q*,√3q*,√4q*, √6q*, √8q*, √14q*, √16q*, √20q*, √22q*, √24q*, and √26q*, respectively, which implies thermally induced order−order phase transition from HEX to gyroid. The dotted line is obtained during cooling scan, confirming thermoreversible characteristics of the observed phase transition. The schematic illustration representing each nanostructure is shown in the right-hand side of each SAXS profile.

respectively. The horizontal dotted lines in the phase diagrams show the phase boundaries for the observed ODTs and OOTs. The regions darken by shadow indicate coexistence temperature window of two phases upon undergoing OOT. Note that all observed morphologies, denoted in Figure 8, are thermally stable and reproducible with multiple heating/cooling cycles from 25 to 190 °C. While the neat S30MB44(17) copolymer shows disordered morphology for entire temperature window examined in present study, the addition of ILs into the S30MB44(17) copolymer results in ordered morphologies, indicative of the increase in χeff of the PEMs with the increased amounts of ionic moieties in SS phases. What is interesting is that a wide variety of ordered structures are obtained by varying the type of diazoles in ILs. For example, when [Py] and [2-MIm] are employed, equilibrium LAM phase was obtained at room temperature with accessible ODT temperature at around 135 ± 5 °C. In contrast, with the use of [Im] and [1-MIm], which possess different ring structures from [Py] and [2-MIm], respectively, HEX morphology was seen at room temperature. This is noteworthy since the [Im] and [Py] (also [1-MIm] and [2-MIm]) reveal similar van der Waals volumes. This leads us to conclude that such discrepancy in morphology is attributed to dissimilar degree of ionic association within SS phases depending on the ring structures of diazoles. For the [Im][TFSI]-incorporated S30MB44(17), the stability window of the HEX morphology was fairly wide and thus, unusually large coexistence window of HEX and LAM morphologies is observed. On the contrary, the phase diagram of the [1-MIm][TFSI]-incorporated sample is occupied by only the HEX phase with no accessible ODT or OOT temperatures up to 190 °C. This signals that the alkyl-substitution at the protic position of [Im] leads to the formation of a highly segregated HEX morphology. When [2-E-4-MIm] having two alkyl substituents on protic positions of [Im] is utilized, an

MIm][TFSI]-containing sample in spite of the large differences in van der Waals volumes. We found that heating the [2-E-4-MIm][TFSI]-incorporated S30MB44(17) copolymer causes intriguing OOT from HEX to gyroid. The SAXS profiles in Figure 7 reveal the development of well-defined gyroid morphology at high temperatures with Bragg peaks (↓) at √6q*, √8q*, √14q*, √16q*, √20q*, √22q*, √24q*, √26q* where q* = 2π/d211 with d211 = 10.4 nm. The Miller indices corresponding to the observed Bragg peaks are (211), (220), (321), (400), (420), (332), (422), and (431), characteristics of the Ia3d̅ space-group symmetry.47 The onset temperature of the OOT was found at 90 °C with narrow temperature window of coexistence between HEX and gyroid phases. At the existence state, the domain size of gyroid phase was found out to be approximately 2.5% smaller than that of HEX structure. At 100 °C, only gyroid phase is evident, which is thermally stable up to 180 °C. Cooling the sample from 180 to 60 °C leads to the reappearance of the HEX peaks, verifying the thermoreversible OOT between HEX and gyroid morphologies. The scattering profiles obtained from the cooling scan are shown by dotted lines. The OOT temperature can be thus located at 90 ± 5 °C. Phase Diagrams of PSS-b-PMB Copolymers Comprising Heterocyclic Diazoles-Based ILs. On the basis of the results obtained thus far, the phase diagrams of S30MB44(17) copolymers containing different heterocyclic diazoles-based ILs 8708

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analogous HEX morphology appeared at room temperature. However, heating the [2-E-4-MIm][TFSI]-containing S30MB44(17) copolymer results in intriguing OOT from HEX to gyroid at around 90 °C via narrow coexistence window. The phase diagram in Figure 8 clearly suggests us that the incorporation of heterocyclic diazole-based ILs into a PSS-bPMB block copolymer causes a large alteration in thermodynamic properties of the systems. In particular, the kinds of diazoles play a central role in determining the morphologies of the IL integrated PEMs. Although, the observed phase transitions of ILs-containing S30MB44(17) copolymers can be mapped on to the well-established phase diagrams of conventional nonionic block copolymers, the classical block copolymer theories provide no basis for understanding the origin of the diazole-sensitive morphologies seen in our system.

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DISCUSSION Effects of Heterocyclic Diazoles on the Conductivities of ILs-Incorporated Polymers. To underpin the effects of heterocyclic diazoles on the conductivities of IL-incorporated PEMs, we first investigated the conductivities of ILs-containing PSS homopolymers. This enables us to evaluate the ion transport behavior of different ILs doped PEMs with excluded morphology effects. The degree of polymerization of a PSS homopolymer was 58 where 11 styrene units were sulfonated (sulfonation level =19 mol %), referred to hereafter as S58(19). Analogous to the PSS-b-PMB block copolymer, we preserved the local concentration of ILs within the SS phases. Therefore, there has been more than 2-fold increase in actual weight fraction of ILs in the ILs-incorporated PSS homopolymers, corresponding to 71 ± 2 wt %. Under the conditions, the [2-E4-MIm][TFSI]-incorporated S58(19) stays in clear liquid form at entire temperature window examined whereas other ILsloaded S58(19) polymers are in sticky gel state at room temperature, which readily turned into liquid-like state with the increase in temperature. Temperature-dependent conductivity data of different ILsincorporated S58(19) homopolymers are shown in Figure 9. Conductivity measurements were carried out through plane of the membranes with the use of a home-built two-electrode cell. Schematic drawing of the conductivity cell is provided in Supporting Information. In the temperature range of 45 to 115 °C, [1-MIm][TFSI]-incorporated S58(19) exhibits the highest conductivity over other ILs-containing samples. In contrast, as listed in Table 3, although the Tg of [2-E-4-MIm][TFSI] integrated S58(19) is lowest as 259 K, the sample reveals considerably low conductivities at all temperatures. This signals the constrained ion motion with the use of [2-E-4-MIm][TFSI] despite the fast segmental motion of polymer chains since the bulky cation experiences large steric hindrance to make longrange ionic association within SS phases. The [Py][TFSI]containing S58(19) shows relatively higher conductivity than that of [Im][TFSI] embedded sample, intimately related to the lower melting temperature (higher diffusion coefficient) of [Py] than [Im]. The conductivity obtained with [2-MIm][TFSI] indicates steep temperature dependence, surpassing conductivities of other samples at high temperatures. It appears that for ILs-incorporated S58(19) homopolymers, the ion transport behavior is largely governed by mass diffusivity of ionic moieties, which is closely related to the intrinsic properties of ILs.

Figure 9. Temperature-dependent through-plane conductivities of ILsincorporated S58(19) homopolymers. The conductivities were obtained at an equimolar IL loading (corresponding to 71 ± 2 wt %) and the types of IL are indicated in the figure. Solid lines indicate analysis using the VTF equation.

Table 3. VTF and Arrhenius Fitting Parameters for Temperature-Dependent Ionic Conductivity S58(19) homopolymer materials used [Im][TFSI] [Py][TFSI] [1-MIm][TFSI] [2-MIm][TFSI] [2-E-4-MIm] [TFSI] a

S30MB44(17) block copolymer

Tg (K)a

B (K), VTF

Ea (kJ/mol), Arrhenius

Tg (K)a

Ea (kJ/mol), Arrhenius

280 275 271 270 259

493 453 447 650 676

31.2 27.1 25.5 34.7 35.1

347 393 331 338 325

56.4 50.5 40.0 59.8 44.8

Measured by DSC at a heating/cooling rate of 5 °C/min.

The conductivity is fitted using Vogel−Tamman−Fulcher (VTF) equation, as given below, to discover potential barriers to ion conduction in different samples.31 ⎛ −B ⎞ σ = σ∞ exp⎜ ⎟ ⎝ T − T0 ⎠

(1)

where σ∞ is the infinite temperature conductivity, B is a fitting parameter related to the potential barriers to ion conduction, and T0 is the temperature at which the polymer relaxation time becomes infinite.48 T0 was set to 50 K below Tg where the Tg was measured by DSC experiments. The best VTF fits are shown as solid lines in Figure 9 and we obtained B values in order of [1-MIm][TFSI] ≈ [Py][TFSI] < [Im][TFSI] < [2Mim][TFSI] ≈ [2-E-4-Mim][TFSI]. This tendency is in good agreement with that of the activation energy (Ea) of ion conduction, as calculated from Arrhenius analysis on the limited data point in the intermediate temperature range of each sample (data not shown here). The Ea values are 31.2, 27.1, 25.5, 34.7, and 35.1 kJ/mol for [Im][TFSI], [Py][TFSI], [1-MIm][TFSI], [2-MIm][TFSI], and [2-E-4-MIm][TFSI], respectively. 8709

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Conductivity of ILs-Incorporated S30MB44(17) Block Copolymers. We now elucidate the conductivity of the ILincorporated S30MB44(17) block copolymers to illuminate the factors governing ion transport behavior of self-assembled PEMs. We intentionally utilized 500 μm thick membranes for the conductivity measurements to rule out the concerns of anisotropic domain orientations. Representative examples of impedance profiles are shown in the Supporting Information. In contrast to the ILs-containing S58(19) homopolymers, the ILs integrated S30MB44(17) block copolymers retain their mechanical integrity at the most temperature window examined. Such difference is essentially caused by different amounts of ILs within the membranes at the equimolar concentration of ILs, i.e., the amounts of ILs within S30MB44(17) block copolymers were as small as 30 ± 2 wt %. Figure 10a represents the through-plane conductivity of S30MB44(17) copolymer upon adding different ILs. As seen in the figure, the conductivities of ILs-incorporated S30MB44(17) copolymers bear resemblance to those of S58(19) homopolymers to some extent. For example, [2-MIm][TFSI]incorporated sample reveals the utmost conductivity at high temperature regions while the achievement of high conductivity values was hampered with the use of [2-E-4-MIm][TFSI]. However, the [Py][TFSI] doped S30MB44(17) copolymer exhibits an order of magnitude low conductivity when compared with other samples, as shown in the inset figure of Figure 10a. This is contrary to what we’ve seen from S58(19) homopolymer. In fact, the conductivity measurement on the [Py][TFSI] embedded S30MB44(17) copolymer was only feasible at high temperature above 140 °C due to the high Tg value of [Py][TFSI] doped S30MB44(17) copolymer (120 °C, see Table 3). This should stem from strong interaction between [−SO3H] and [Py], as we discussed in Table 2 and Figure 5. Because of the steep temperature dependence of conductivity of [Im][TFSI]-incorporated sample, the conductivity of [Im][TFSI]-incorporated copolymer overtakes that obtained with [1-MIm][TFSI] at very high temperature. This leads us to conclude that for S30MB44(17) block copolymer, the use of [1MIm] is not all that good in achieving high conductivity, particularly, at the elevated temperatures. In contrast, the introduction of [2-MIm] appears to be worthwhile to attain improved conductivity in a wide range of temperature. Solid lines in Figure 10a indicate analysis using the Arrhenius relationship. From the slope of Arrhenius plots, the Ea are estimated at 56.4, 50.5, 40.0, 59.8, and 44.8 kJ/mol for [Im][TFSI], [Py][TFSI], [1-MIm][TFSI], [2-MIm][TFSI], and [2-E-4-MIm][TFSI], respectively. The lowest Ea is again seen with the use of [1-MIm][TFSI]. Unlike the tendency seen with S58(19) homopolymer, it is interesting to note that the [2E-4-MIm][TFSI] containing sample indicates relatively low level of Ea value. In Table 3, we summarize the Ea values of ILs embedded S 58 (19) homopolymers and IL integrated S30MB44(17) block copolymers. To take into account the different chain mobilities depending on the kinds of heterocyclic diazoles, the conductivities in Figure 10a are replotted in Figure 10b upon normalizing temperature axis by Tg of the membranes. Overall, the normalized conductivities of ILs-incorporated S30MB44(17) copolymers lie in order of [2-MIm] ≈ [Im] > [1-MIm] > [2-E4-MIm] ≫ [Py] with vanished Tg effects. It should be noteworthy that the normalization results in almost coincident conductivity profiles for [Im][TFSI] and [2-MIm][TFSI]incorporated S30MB44(17) copolymers in the temperature

Figure 10. (a) Temperature-dependent through-plane conductivities of ILs-incorporated S30MB44(17) copolymers. (b) Normalized conductivities of ILs-incorporated S30MB44(17) copolymers on the basis of Tg of the membranes. In parts a and b, the linear fits were obtained by Arrhenius analysis. The conductivity values were obtained with an equimolar IL concentration (corresponding to ca. 30 wt %) and the types of IL are indicated in the figure.

window of our interest. This leads us to conclude that the ion conduction phenomena seen with [Im][TFSI] and [2MIm][TFSI] is similar in origin on account of the preserved protic sites in the ring structure. In Figure 11, we show the proposed ion transport mechanisms in ILs impregnated PSS-b-PMB copolymers. For brevity, we only discuss a subset of diazoles listed in Table 1, i.e., [1-MIm] and [2-MIm], as representative examples. The [TFSI] anion was chosen so as to ensure low Tg of ILs embedded membranes. The direct comparison of [1-MIm][TFSI] and [2-MIm][TFSI] allows us to study the effect of the number of protic sites in diazoles on the ion transport behavior of ILs containing PSS-b-PMB copolymers, taking place within SS phases. As shown in Figure 11, the proton transfer from −SO3H sites of PSS-b-PMB to the nitrogen atom of [Im] derivatives is occurred by the conversion of hydrogen bond into ionic bond. For the case of [2-MIm][TFSI], we expect that a 8710

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the HEX and LAM at the least. To explore the morphology effects on conductivities in details, we evaluate a morphology factor ( f) of each system. The conductivity of ILs-containing S30MB44(17) copolymers was normalized by the homopolymer conductivity using eq 2 wherein the dependence of conductivity on polymer segmental motion can be factored out.36 σ f = SMB φSSσSS (2) Here σSMB is the conductivity of IL integrated S30MB44(17) copolymer and σ SS is that of IL embedded S 58 (19) homopolymer. We assume that the conductivity of block copolymer is proportional to the volume fraction of the conducting SS phases, ϕSS. For simplicity, ϕSS = 0.5 was used for LAM ([2-MIm][TFSI]) and gyroid-forming ([2-E-4MIm][TFSI]) samples while ϕSS = 0.67 was employed for HEX-forming copolymer ([1-MIm][TFSI]). These values were roughly estimated by combining the information on the change in domain sizes (SAXS) and thicknesses of SS domains and MB phases (TEM) for ILs-incorporated copolymers. Regarding the LAM/HEX coexistence ([Im][TFSI]), an average value of 0.5 and 0.67 was applied. Figure 12 represents the morphology factors obtained for ILs-incorporated S30MB44(17) copolymers. Interestingly, [Im]-

Figure 11. Schematic drawings illustrating the proposed ion transport mechanism in ILs impregnated PSS-b-PMB copolymers; (a) [2MIm][TFSI] and (b) [1-MIm][TFSI]. The proton transfer from −SO3H sites of PSS-b-PMB to the nitrogen atom of [Im] derivaties is occurred by the conversion of hydrogen bond into ionic bond. The solid line and dashed line represent ionic bond and hydrogen bond, respectively. The solid arrow and dashed arrow indicate the movement of electron and proton, respectively.

series of hopping and diffusion can carry protons from one of the −SO3H sites to neighboring [2-MIm] species and [TFSI−] counterions via consecutive formation of ionic bonds and hydrogen bonds. In contrast, the hydrogen bonding among [1-MIm] molecules is impeded by the formation of the ionic bond between −SO3H and [1-MIm]. The hydrogen bonding should involve the dissociation of the ionic bond and the rotation of [1-MIm] molecules, which is considered to be a ratedetermining step. Consequently, the ion transport in [1MIm][TFSI]-incorporated PSS-b-PMB copolymer is attributed to the diffusion of ionic moieties and/or proton relay from −SO3H sites to [1-MIm] cations via HTFSI mediators. Although the detailed analysis of the ion transfer mechanism in ILs-containing PSS-b-PMB block copolymers is a subject of future study, the fact that the we were able to achieve enhanced conductivity with the [2-MIm][TFSI] integrated S30MB44(17) copolymer by preserving protic sites in diazoles is worth noting (see Figure 10b). Morphology Effects on the Conductivities of ILsIncorporated S30MB44(17) Copolymers. The fact that [Im][TFSI]-incorporated S30MB44(17) copolymer and [2MIm][TFSI] embedded sample reveal similar normalized conductivities (see Figure 10b) despite the dissimilar morphologies of HEX+LAM coexistence and LAM, respectively, suggests that the morphology effect is not substantial for

Figure 12. Morphology factors of ILs-incorporated S30MB44(17) copolymers, as calculated by normalized conductivity using eq 2. Dashed lines mark the positions of f = 1/3 and f = 2/3.

[TFSI], [1-MIm][TFSI], and [2-MIm][TFSI] containing samples yield qualitatively similar f values at 0.4 ± 0.06 wherein the corresponding morphologies are HEX+LAM coexistence, HEX, and LAM, respectively. This again signals that the randomly oriented HEX and LAM morphologies make virtually no difference in ion transport behavior. However, the morphology factor obtained with [2-E-4-MIm][TFSI] was found out to be evidently large in the range 0.6−0.7. This is intimately related to the unique structural advantages of gyroid morphology where the ionic channels are arranged along the cocontinuous SS phases. In number of literatures,36,49 the f = 1/ 3 and 2/3 are expected for randomly organized HEX (conducting phase is a minor domain) and LAM structures, respectively, while f = 1 was predicted for gyroid structure. In 8711

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present study, the reduction in f values below the theoretical ones should be due to the existence of grain boundaries, hampering ion transport across the grains. Our results suggest that if one can create gyroid morphology with [2-MIm][TFSI] by varying chemical composition of PSS-b-PMB copolymers, even further enhanced conductivity can be accomplished. In present study, unfortunately, we were unable to examine the effects of order−order phase transitions on the conductivities since the OOT temperatures were so close to the low temperature limit to measure plateau impedance values at high frequency. Regarding the LAM-to-disorder transitions of ILs-incorporated S30MB44(17) copolymers, we were able to access high temperature, 30 °C above the TODT. As can be seen from Figure 10a, however, there is no evident transition in the conductivity profiles for both [2-MIm][TFSI]-incorporated S30MB44(17) copolymer and [Py][TFSI] embedded sample. We speculate that the complete melting of ordered LAM grains was not occurred and thus, the ion transport along the conducting SS phases is not significantly affected by the presence of the nonconducting MB phases due to the large magnitude of concentration fluctuations. Note that the highest temperature of our conductivity measurements was 180 °C (simply due to the thermal stability of the sheath of cables) where the calibrated sample temperature at such conditions was 165 °C. The ability to control the conductivity of ILs doped PEMs by adjusting kinds of ILs and the morphologies of the membranes allowed us not only to explore the link between morphology and transport in quantitative way but also to achieve efficient ion transport in PEMs. Our results should be important for diverse applications such as high temperature fuel cells and electro-active actuators.

and structural optimization of IL-incorporated PEMs are crucial for the access of desired ion transport properties.

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ASSOCIATED CONTENT

S Supporting Information *

Schematic drawing of a conductivity cell and SAXS data of S30MB44 copolymers by varying sulfonation level. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was financially supported by Basic Science Research Program (Project No. 2012-0001993) and by Midcareer Researcher Program (Project No. 2012-0005267) through the National Research Foundation of Korea (NRF), funded by the Ministry of Education, Science and Technology. We also acknowledge the Global Frontier R&D program on Center for Multiscale Energy System funded by the NRF under the Ministry of Education, Science and Technology. SAXS measurements were conducted on the beamline 4C at the Pohang Light Source (PLS) supported by the Ministry of Science and Technology of Korea.

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REFERENCES

(1) Terasawa, N.; Ono, N.; Mukai, K.; Koga, T.; Higashi, N.; Asaka, K. Carbon 2012, 50, 311−320. (2) Mukai, K.; Asaka, K.; Suigino, T.; Kiyohora, K.; Takeuchi, I.; Terasawa, N.; Futaba, D. N.; Hata, K.; Fukushima, T.; Aida, T. Adv. Mater. 2009, 21, 1582−1585. (3) Imaizumi, S.; Kokubo, H.; Watanabe, M. Macromolecules 2008, 8 (1), 265−270. (4) Seki, S.; Ohno, Y.; Kobayashi, Y.; Miyashiro, H.; Usami, A.; Mita, Y.; Tokuda, H.; Watanabe, M.; Hayamizu, K.; Tsuzuki, S. J. Electrochem. Soc. 2007, 154 (3), A173−A177. (5) Kim, G. T.; Jeong, S. S.; Xue, M. Z.; Balducci, A.; Winter, M.; Passerini, S.; Alessandrini, F.; Appetecchi, G. B. J. Power Sources 2012, 199, 239−246. (6) Choi, I.; Ahn, H.; Park, M. J. Macromolecules 2011, 44 (18), 7327−7334. (7) Steininger, H.; Schuster, M.; Kreuer, K. D.; Kaltbeitzel, A.; Bingöl, B.; Meyer, W. H.; Schauff, S.; Brunklaus, G.; Maier, J.; Spiess, H. W. Phys. Chem. Chem. Phys. 2007, 9, 1764−1773. (8) Tang, Q.; Wu, J.; Tang, Z.; Li, Y.; Lin, J. J. Mater. Chem. 2012, 22, 15836−15844. (9) Kim, S. Y.; Kim, S.; Park, M. J. Nat. Commun. 2010, 1, 88. (10) Mistry, M. K.; Subianto, S.; Choudhury, N. R.; Dutta, N. K. Langmuir 2009, 25, 9240−9251. (11) Ohno, H. Electrochemical Aspects of Ionic Liquids: John Wiley & Sons, Inc.: Hoboken, NJ, 2005. (12) Armand, M.; Endres, F.; MacFarlane, D. R.; Ohno, H.; Scrosati, B. Nat. Mater. 2009, 8, 621−629. (13) Sowmiah, S.; Srinivasadesikan, V.; Tseng, M. C.; Chu, Y. H. Molecules 2009, 14, 3780−3813. (14) Green, M. D.; Long, T. E. Polym. Rev. 2009, 49, 291−314. (15) Hou, J.; Zhang, Z.; Madsen, L. A. J. Phys. Chem. B 2011, 115, 4576−4582. (16) Lalia, B. S.; Yamada, K.; Hundal, M. S.; Park, J. J.; Park, G. G.; Lee, W. Y.; Kim, C. S.; Sekhon, S. S. Appl. Phys. A: Mater. Sci. Process 2009, 96, 661−670.

CONCLUSIONS We hope to conclude our paper by making answers for the questions that we have brought up in Introduction. (1) We successfully created various well-defined nanostructures for sulfonated block copolymers comprising different heterocyclic diazole-based ILs. It has been found that the kinds of diazoles play an important role in determining the morphologies of ILsincorporated copolymers. For example, the addition of [Im][TFSI] into disordered PSS-b-PMB copolymer results in the development of HEX morphology whereas in the presence of [Py][TFSI] and [2-MIm][TFSI], the copolymers exhibit LAM morphology. The enhanced segregation strength between PSS and PMB microdomains was observed upon substituting methyl group at the protic site of [Im], i.e., [1-MIm]. Morphology effects on conductivities seemed to certainly exist in the IL impregnated PSS-b-PMB copolymers although further experiments on the extended set of samples are remained as future study. (2) Compared with [Py], [Im] seems to be a better proton source, particularly for the block copolymer matrix. This is due to the restricted rotational motion of [Py] molecules by making strong ionic bond between [Py] and [−SO3H], which limits long-range ion transport. (3) From the conductivity results obtained with [1MIm][TFSI] and [2-MIm][TFSI], [2-MIm] is appeared to be a preferable cation to achieve enhanced conductivity values owing to the preserved protic sites. However, even with two protic sites, the Brönstead-type diazole having long alkylsubstituents undoubtedly limits efficient ion transport among ionic moieties, as seen with [2-E-4-MIm]. Our results demonstrate that both molecular design of heterocyclic diazoles 8712

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Article

(17) Green, M. D.; Salas-de la Cruz, D.; Ye, Y.; Layman, J. M.; Elabd, Y. A.; Winey, K. I.; Long, T. E. Macromol. Chem. Phys. 2011, 212, 2522−2528. (18) Ye, Y.; Elabd, Y. A. Polymer 2011, 52 (5), 1309−1317. (19) Fernicola, A.; Panero, S.; Scrosati, B. J. Power Sources 2008, 178, 591−595. (20) Cho, E. K.; Park, J. S.; Sekhon, S. S.; Park, G. G.; Yang, T. H.; Lee, W. Y.; Kim, C. S.; Park, S. B. J. Electrochem. Soc. 2009, 156, B197− B202. (21) Lee, H. N.; Newel, N.; Bai, Z.; Lodge, T. P. Macromolecules 2012, 45 (8), 3627−3633. (22) Gwee, L.; Choi, J. H.; Winey, K. I.; Elabd, Y. A. Polymer 2010, 51 (23), 5516−5524. (23) Susan, M. A. B. H.; Kaneko, T.; Noda, A.; Watanabe, M. J. Am. Chem. Soc. 2005, 127, 4976−4983. (24) Hoarfrost, M. L.; Tyagi, M. S.; Segalman, R. A.; Reimer, J. A. Macromolecules 2012, 45, 3112−3120. (25) Virgili, J. M.; Hexemer, A.; Pople, J. A.; Balsara, N. P.; Segalman, R. A. Macromolecules 2009, 42, 4604−4613. (26) Kim, S. Y.; Yoon, E.; Joo, T.; Park, M. J. Macromolecules 2011, 44, 5289−5298. (27) Chen, L.; Hallinan, D. T.; Elabd, Y. A.; Hillmyer, M. A. Macromolecules 2009, 42 (16), 6075−6085. (28) Elabd, Y, A.; Hickner, M. A. Macromolecules 2011, 44 (1), 1−11. (29) Missan, H. P. S.; Lalia, B. S.; Karan, K.; Maxwell, A. Mater. Sci. Eng., B 2010, 175, 143−149. (30) Du, Q.; Fu, X.; Liu, S.; Niu, L. Polym. Int. 2012, 61, 222−227. (31) Choi, U. H.; Lee, M.; Wang, S.; Liu, W.; Winey, K. I.; Gibson, H. W.; Colby, R. H. Macromolecules 2012, 45, 3974−3985. (32) Lee, M.; Choi, U. H.; Cruz, D. S.; Mittal, A.; Winey, K. I.; Colby, R. H.; Gibson, H, W. Adv. Funct. Mater. 2011, 21, 708−717. (33) Liu, Y.; Liu, S.; Lin, J.; Wang, D.; Jain, V.; Montazami, R.; Heflin, J. R.; Li, J.; Madsen, L.; Zhang, Q. M. Appl. Phys. Lett. 2010, 96, 223503. (34) Kreuer, K. D.; Paddison, S. J.; Spohr, E.; Schuster, M. Chem. Rev. 2004, 104, 4637−4678. (35) Weber, R. L.; Ye, Y.; Schmitt, A. L.; Banik, S. M.; Elabd, Y. A.; Mahanthappa, M. K. Macromolecules 2011, 44 (14), 5727−5735. (36) Wanakule, N.; Panday, A.; Mullin, S. A.; Gann, E.; Hexemer, A.; Balsara, N. P. Macromolecules 2009, 42, 5642−5651. (37) Salas-de la Cruz, D.; Green, M. D.; Ye, Y.; Elabd, Y. A.; Long, T. E.; Winey, K. I. J. Polym. Sci., Part B: Polym. Phys. 2012, 50 (16), 338− 346. (38) Mangiatordi, G. F.; Hermet, J.; Adamo, C. J. Phys. Chem. A 2011, 115 (12), 2627−2634. (39) Endo, T.; Kato, T.; Nishikawa, K. J. Phys. Chem. B 2010, 114 (28), 9201−9208. (40) Schuster, M. F. H.; Meyer, W. H.; Schuster, M.; Kreuer, K. D. Chem. Mater. 2004, 16, 329−337. (41) Kreuer, K. D.; Fuchs, A.; Ise, M.; Spaeth, M.; Maier, J. Electrochem. Acta 1998, 43 (10−11), 1281−1288. (42) Lenarcik, B.; Ojczenasz., P. J. Heterocyclic Chem 2002, 39, 287− 290. (43) S., Ü .; Bozkurt, A.; Hosseini, S. S. Prog. Polym. Sci. 2012, 37, 1265−1291. (44) Miran, M. S.; Kinoshita, H.; Yasuda, T.; Susan, M. A. B. H.; Watanabe, M. Phys. Chem. Chem. Phys. 2012, 14, 5178−5186. (45) Hanley, K. J.; Lodge, T. P.; Huang, C.-I. Macromolecules 2000, 33 (16), 5918−5931. (46) Stephens, P. J.; Devlin, F. J.; Chabalowski, C. F.; Frisch, M. J. J. Phys. Chem. 1994, 98, 11623−11627. (47) Alexandridis, P.; Olsson, U.; Lindman, B. Langmuir 1998, 14 (10), 2627−2638. (48) Vogel, H. Phys. Z. 1921, 22, 645. (49) Sax, J.; Ottino, J. M. Polym. Eng. Sci. 1983, 23 (3), 165−176.

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