silica for high-performance polymer electrolyte membrane fuel cells under low relative humidity

Journal of Membrane Science 549 (2018) 567–574

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Nanostructured composite membrane with cross-linked sulfonated poly (arylene ether ketone)/silica for high-performance polymer electrolyte membrane fuel cells under low relative humidity

T

⁎

Insung Baea, , Keun-Hwan Oha, Minhyuk Yunb, Min Kwan Kangc, Hyun Hoon Songc, Hyuk Kima a

Corporate R&D, LG Chem, 188, Munji-ro, Yuseong-gu, Daejeon, Republic of Korea IT&E Materials R&D, LG Chem, 188, Munji-ro, Yuseong-gu, Daejeon, Republic of Korea c Department of Advanced Materials, Hannam University, 1646, Yuseong-daero, Yuseong-gu, Daejeon, Republic of Korea b

A R T I C L E I N F O

A B S T R A C T

Keywords: Cross-linking Polymer electrolyte fuel cells Proton exchange membrane Silica composite Sulfonated poly(arylene ether ketone)

Proton exchange membranes (PEMs) consisting of cross-linked sulfonated poly(arylene ether ketone) with silica nanoparticles (CL-SPAEK/silica) are developed for fuel cell applications under low relative humidity (RH). CLSPAEK/silica is prepared by the reduction reaction of the synthesized SPAEK, followed by a sulfonation reaction between the benzophenone moieties of SPAEK and the hydroxyl groups of silica nanoparticles. The cross-linked structure of CL-SPAEK/silica enhanced proton conductivity that attributed to the synergistic effect of wellconnected hydrophilic proton channels and improved the dispersion of silica nanoparticles in the polymer matrix. The nanostructure of the membranes was investigated by field-emission transmission electron microscopy (FE-TEM), atomic force microscopy (AFM) and small angle X-ray scattering (SAXS). Furthermore, CLSPAEK/silica exhibited enhanced mechanical stability due to the effectively cross-linked networks between SPAEK and silica.

1. Introduction Proton exchange membrane fuel cells (PEMFCs) have great interest for environmental-friendly and efficient energy conversion devices [1–4]. Perfluorosulfonic acid (PFSA) membranes such as Nafion and Aquivion are most widely used as electrolytes because they have excellent chemical and mechanical stabilities as well as high proton conductivity. However, PFSA membranes have some specific shortcomings such as high cost, limited operating temperature and high fuel permeability [5–7]. To overcome these drawbacks of PFSA membranes, alternative polymer electrolyte membranes (PEMs) based on sulfonated aromatic hydrocarbon polymers have been intensively developed such as sulfonated poly(ether ether ketone) (SPEEK), sulfonated poly(arylene ether sulfone) (SPAES), sulfonated poly(arylene ether ketone) (SPAEK), and sulfonated polyimide (SPI) [8–14]. Among these hydrocarbon polymers, SPAEKs have great potential due to low cost, high proton conductivity, excellent thermal stability and low fuel crossover [12–14]. The proton conductivity of hydrocarbon membranes is significantly influenced by the degree of sulfonation (DS) depending on the number of sulfonic acid groups and an effective formation of hydrophilic channels for proton transport. Although high DS of hydrocarbon membranes is beneficial to fast proton conduction, increasing ⁎

the DS may lead to excessive membrane swelling which can reduce the mechanical properties of the membranes and fuel cell performance. In addition, they suffer from the drawbacks of significant deterioration of proton conductivity and poor dimensional stability by severe water loss under low humidity conditions and structural deformation according to humidity level, respectively [15,16]. Reduction of proton transport resistance under low humidity conditions and dimensional changes in the membrane are the most significant challenges of hydrocarbon PEMs for commercialization in fuel cell applications. As a promising strategy, organic-inorganic nanocomposites have been widely investigated to improve proton conductivity under low humidity conditions and dimensional stability with a change in humidity [17–19]. Various inorganic nanofillers have been reported for organic-inorganic nanocomposites, including silica [13,14,20–23], zirconia [24], graphene [25,26] and boron nitride [27]. In particular, silica nanoparticles have been introduced in sulfonated hydrocarbon polymer membranes to improve water retention increasing low humidity proton conduction and enhance the mechanical stability of the membrane with hydrogen bonding [13,14,20–23]. Despite the adoption of silica to improve water retention capacity and mechanical stability of hydrocarbon membrane, the weak interfacial interaction between the hydrocarbon polymer matrix and silica nanofillers restricts

Corresponding author. E-mail address: [email protected] (I. Bae).

https://doi.org/10.1016/j.memsci.2017.12.060 Received 12 November 2017; Received in revised form 19 December 2017; Accepted 20 December 2017 Available online 22 December 2017 0376-7388/ © 2017 Elsevier B.V. All rights reserved.

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filtered and washed with deionized water until the water was neutralized without residual solvents and dried overnight at 80 °C in a vacuum oven. The schematic details of the synthesis are described in Fig. S1.

homogeneous dispersion of silica nanofillers [20]. The agglomeration of inorganic nanoparticles particularly restricts the connectivity of hydrophilic moieties in the polymer chain, which reduces the bound water and proton conductivity under low humidity conditions [23,24]. Kim et al. reported the SPAES membrane containing sulfonated zirconia for high-temperature and low humidity operation of PEMFCs [24]. Using small angle X-ray scattering (SAXS) and atomic force microscopy (AFM), they found that the homogeneous dispersion of sulfonated zirconia in SPAES membrane improved proton conductivity of the composite membranes with effective interconnection between the sulfonic acid groups and bound water absorbed on the surface of the zirconia nanoparticles. Considering membrane morphology to improve proton conductivity under low humidity conditions, PEMs have involved a distinct phase separation between hydrophilic and hydrophobic blocks with high domain interconnectivity [8,28–30]. Cross-linking is one promising strategy to achieve the desired properties including proton conductivity and mechanical/chemical stability, thus enhancing power generation and durability [31–33]. Particularly, the cross-linked structure of organic-inorganic composite membranes has been adopted to prevent the agglomeration of inorganic nanoparticles, leading to well-connected hydrophilic proton transport channels [23,34]. Here we present the development of the nanostructured composite membrane through the facile cross-linking reaction for enhancing fuel cell performance under low humidity conditions. Benzophenone moieties of SPAEK were reduced into benzhydrol groups using sodium borohydride as a reducing agent. The reduced SPAEK (R-SPAEK) was further reacted with sulfuric acid to produce cross-linked networks between the benzhydrol groups of the polymers and the hydroxyl groups of the silica nanoparticles. The chemically bonded cross-linking between SPAEK and silica not only helped the uniform dispersion of silica nanoparticles but also promoted a well-developed nanostructure of the copolymer, resulting in the development of proton transport pathways under low humidity conditions. Moreover, the cross-linked structure increased the mechanical stability of the membrane for longterm operation.

2.3. Preparation of cross-linked composite membrane The synthesized SPAEK (5 wt% in DMSO) was dissolved in DMSO with dispersed silica (5 wt% to SPAEK) at 70 °C for 4 h. The composite solution and NaBH4 (1 wt% to SPAEK) were added to a round bottle with a reflux condenser and a drying tube for the reduction reaction of the benzophenone moieties of SPAEK polymer. The mixture was heated at 120 °C for 12 h with vigorous stirring to make a homogeneous solution. The composite solution was cast onto a clean glass plate and dried at 50 °C for 3 h followed by additional drying under a vacuum at 100 °C for 12 h. The resultant composite membranes were simultaneously cross-linked and converted to the required acid form by immersing in 1.0 M sulfuric acid solution at 80 °C for 24 h, and then rinsed with deionized water to remove the residual acid. After that, the prepared membranes were dried in a vacuum oven at 80 °C for 4 h. The reduction and cross-linking reaction details of SPAEK with silica nanoparticles are given in Fig. S2. The pristine, composite and crosslinked composite membranes were designated as SPAEK, SPAEK/silica and CL-SPAEK/silica, respectively. All membranes were carefully controlled at 15 µm thickness. 2.4. Membrane characterization The water uptake and swelling ratio of the membranes were measured by weight and volume before and after the wet treatment, respectively. The prepared membranes were dried in a vacuum oven at 80 °C for 24 h and immersed in deionized water at various temperatures of 25 °C, 50 °C, 70 °C, and 90 °C for 24 h. Then the wet membranes were quickly swabbed to remove the water on the sample surface. The water uptake (%) and swelling ratio (%) of membranes were calculated using the following equations:

Wateruptake =

2. Experimental 2.1. Materials and chemicals

∆L (or ∆T ) = 4,4′-difluorobenzophenone (DFBP), hydroquinone sulfonic acid potassium salt (HQS), potassium carbonate, sodium borohydride (NaBH4), anhydrous dimethyl sulfoxide (DMSO), 9,9-bis(hydroxyphenyl)fluorene (BHF), benzene, anhydrous ethyl alcohol and concentrated sulfuric acid (98%) were purchased from Sigma-Aldrich and used without further purification. Fumed silica (Aerosil 380, average particle size = 7 nm, BET surface area = 380 ± 30 m2 g−1) was purchased from Evonik Industries.

Wwet − Wdry Wdry

L wet − Ldry Ldry

(or

× 100

Twet − Tdry Tdry

(1)

) × 100

(2)

where Wwet, Lwet and Twet are the weight, length, and thickness of a wet membrane, respectively. Wdry, Ldry, and Tdry are similarly defined for a dry membrane. The thermal properties of the membranes were obtained using thermal gravimetric analyzers (TGA, TA Instrument Q500) and differential scanning calorimetry (DSC, TA Instrument Q100). For TGA characterization, the composite membranes were dried at 150 °C for 10 min to remove moisture and heated to 850 °C at a rate of 5 °C min−1 under N2 flow. DSC results were carried out for observing a glass transition temperature (Tg) of the membranes and investigating the physical state of water molecules in the swollen membranes. The Tg of the membranes was obtained from the second scan, after first heating up to 200 °C then cooling to 25 °C at a rate of 5 °C min−1 under N2 atmosphere. To determine the water state in the membranes, the hydrated membranes were sealed in sample pans and then cooled to −30 °C. After equilibrating for 10 min, the membranes were heated to 40 °C under N2 flux with a heating rate of 5 °C min−1. The amount of bound water (%) was obtained using the following equation using the melting phenomena of water:

2.2. Synthesis of SPAEK copolymer SPAEK copolymers were synthesized via one-pot condensation reaction [13,14]. DFBP (8.47 g, 38.82 mmol), HQS (8.22 g, 36.01 mmol), potassium carbonate (9.96 g, 72.03 mmol), DMSO (52.47 mL) and benzene (40 mL) were added into a 500 mL round bottle flask equipped with a Dean-Stock trap, mechanical stirrer, thermometer, reflux condenser and N2 inlet/outlet for hydrophilic domain of SPAEK block copolymer. After dehydration at 140 °C for 4 h, the mixture was heated up to 180 °C for 20 h. Then the flask was charged with DFBP (1.31 g, 6.00 mmol), BHF (3.58 g, 10.21 mmol), potassium carbonate (0.06 g, 0.4 mmol), DMSO (57.65 mL) and benzene (20 mL) to synthesize hydrophobic blocks of the copolymer. The reaction mixture was heated to 140 °C for 4 h and the temperature gradually heated up to 180 °C for 20 h. The highly viscous polymer solution was diluted with 40 mL of DMSO and then slowly coagulated in methanol. The precipitate was

Wb = Wt − Wf = Wt − (

Qendo × 100) Qpure

(3)

where Wt is the total amount of water calculated by Eq. (1); Wf and Wb are the amounts of free and bound water, respectively; Qendo is the 568

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Fig. 1. Photograph and schematic diagram of a) SPAEK, b) SPAEK/silica and c) CL-SPAEK/silica membranes.

groups then rinsed with deionized water followed by drying in a vacuum oven. Atomic force microscopy (AFM, Park Systems XE7) was used to examine the surface morphology of the membranes in a noncontact mode under 30% RH. Characterizing the morphology of the polymeric domains in the membranes, small angle X-ray diffraction (SAXS, Rigaku D/max-2500) was operated at 40 kV and 70 mA with CuKα as an X-ray source (λ = 1.542 Å). The AFM and SAXS measurements were carefully performed at 25 °C under 30% RH. The tensile strength and elongation of the membranes were obtained using a universal testing machine (UTM, SHIMADZU AGS-X 100 N) at 25 °C under 30% RH with a humidity control system. The membranes were cut into dumbbell-shaped strips, fixed between pneumatic grips and stretched at the speed of 5 mm min−1.

endothermic fusion enthalpy of a hydrated membrane near 0 °C during the heating process; and Qpure is 334 J g−1, the melting enthalpy of free water to the heat of pure water fusion. Attenuated total reflectance-Fourier transform infrared (ATR-FTIR) spectra of SPAEK, R-SPAEK/silica and CL-SPAEK/silica were measured with ATR-FTIR spectrophotometer (Bruker) at room temperature. The proton conductivity was evaluated with a four-electrode system using a membrane conductivity cell by the AC impedance spectroscopy method (Bio-Logics, HCP-803). The impedance measurements were performed at 70 °C under RH range of 30–100%. The proton conductivity (σ) of membranes that contacted with four platinum electrodes was calculated by using the following equation:

σ=

L R×S

(4) 2.5. Fabrication and electrochemical characterization of membrane electrode assembly (MEA)

L represents the distance between the two sensing electrodes; R is the measured resistance from a Nyquist impedance diagram; and S is the cross-sectional surface area of a membrane. The ion exchange capacity value based on the weight (IEC) of a membrane was determined by an acid-base titration. The dry membrane was immersed into 1 M NaCl solution for 24 h for a complete ion exchange from H+ to Na+ and this solution was titrated with 0.01 M NaOH solution to neutralize the exchanged H+ using phenolphthalein as an indicator. The IEC was calculated according to the following equation:

IEC =

V×M Wdry

MEA was fabricated by the transferred decal printing of prepared Pt/C layers coated on Poly(tetrafluoroethylene) (PTFE) substrate. The catalyst layers were placed on both sides of the membrane and hotpressed at 140 °C under the pressure of 102 kg f cm−2 for 5 min. The Pt loading for each electrode was controlled at 0.35 ± 0.02 mg cm−2. The prepared MEA was performed in a single cell fixture with the triple serpentine flow fields, glass fiber-reinforced gaskets sealants, and gas diffusion layer (10BB, SGL). PEMFC performance of the fabricated MEA with an active area of 25 cm2 was tested at 70 °C cell temperature using a fuel cell evaluation system (Scribner Associates Inc., USA). The flow rate of hydrogen gas to the anode and air to the cathode corresponded to the stoichiometry 1.5 and 2.0 with ambient pressure, respectively. The polarization curve of PEMFC was measured after the activation at 0.6 V for 8 h under fuel supply. Electrochemical impedance spectroscopy (EIS) was measured using an AC impedance analyzer (Bio-Logics, HCP-803) over the frequency range of 105 to 0.1 Hz and the AC amplitude (10 mV) of the applied current to obtain the Nyquist plot of a cell.

(5)

where V is the volume of NaOH solution added at the equivalent point; M is the concentration of NaOH solution; and Wdry is the mass of a dry membrane. Dark field transmission electron microscopy (TEM, FEI corp. Titan G2 80-200) was used to examine the cross-sectional morphology of the membranes by the lead stain method. The membranes were immersed in a 0.5 M lead acetate aqueous solution to exchange sulfonic acid 569

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The wet-dry cycle was executed for a mechanical durability test. The fully humidified (150% RH) and dry (0% RH) N2 gas were alternately supplied every 2 min at a flow rate of 1000 sccm, each at 80 °C. The hydrogen crossover current was measured at 80 °C with 300 sccm of H2 and N2 to the anode and cathode, respectively. The H2 crossover current density was recorded by linear sweep voltammetry from 0.6V to 0.05 V.

3. Results and discussion 3.1. Characterizations of CL-SPAEK/silica The synthesized SPAEK, SPAEK/silica and CL-SPAEK/silica are schematically depicted in Fig. 1. For the preparation of CL-SPAEK/silica, the benzhydrol groups of R-SPAEK polymers were facilitated by the reduction of benzophenone moieties using sodium borohydride followed by a cross-linking reaction between the benzhydrol groups of R-SPAEK and the hydroxyl groups of silica nanoparticles. (Fig. S2) In the cross-linking reaction, the sulfonic acid group of R-SPAEK also react with sulfonic/sulfuric acid resulting in a slight decrease of IEC in CLSPAEK/silica compared to SPAEK and SPAEK/silica. (Table S1) [32] The color of the membrane changed from pale yellow to reddish brown by the cross-linking reaction as shown in Fig. 1. Without the crosslinking reaction, silica nanoparticles easily aggregated in the polymer matrix due to weak hydrogen bonding with SPAEK. However, the strong covalent bond between the polymer and silica surface of the cross-linking structure allowed for a uniform distribution of silica as shown in the surface morphology in Fig. S3. The strong interfacial interaction was confirmed by a simple adhesion and solubility test (Figs. S4 and S5). Fig. 2a comparatively shows ATR-FTIR spectra of SPAEK/silica, RSPAEK/silica, and CL-SPAEK/silica. The strong absorption peak of benzophenone at 1651 cm−1, representing the C˭O stretch can be observed in SPAEK/silica while the peak becomes much smaller in RSPAEK/silica and CL-SPAEK/silica due to the reduction reaction. Moreover, the spectrum of R-SPAEK/silica showed an increased peak at 3458 cm−1, corresponding to the hydroxyl stretching vibration of the benzhydrol group. These hydroxyl groups of SPAEK were the reaction sites for cross-linking between the polymer chains and silica nanoparticles via sulfuric acid which resulted in the decrease of the peak intensity at 3458 cm−1 after cross-linking. Thermal properties of the membranes were investigated by the second scan of DSC as shown in Fig. 2b. SPAEK shows a glass transition temperature (Tg) of 183.4 °C and the Tg of SPAEK/silica decreased from 183.4 °C to 171.6 °C due to the enlarged free volume arising from the interface between the polymer matrix and aggregated silica fillers [35]. However, the CL-SPAEK/silica displayed an increased Tg of 209.9 °C due to restricted chain mobility of polymers by a strong interfacial interaction between the polymer main chains and silica nanoparticles. In Fig. 2c, TGA was used for further thermal characterization of the composite membranes. For an initial state of dehydrated membranes, the temperature was kept at 150 °C for 10 min under N2 flow. The decomposition reaction of SPAEK and SPAEK/silica started at 200 °C, while the degradation of CL-SPAEK/silica started around 250 °C due to the enhanced molecular interaction of the cross-linked structure. The first decomposition from 250 °C to 400 °C is caused by the cleavage of the sulfonic acid groups from the polymer main chain and the second reaction is the decomposition of the polymer backbone at above 500 °C. The weight residues of composite membranes were higher than that of SPAEK at 800 °C, representing the silica contents [23]. These thermal analytic results imply that CL-SPAEK/silica has a higher thermal stability than SPAEK and SPAEK/silica. The thermal stability of the crosslinked composite membrane was attributed to the cross-linked structure between the benzhydrol groups of R-SPAEK polymers and the hydroxyl groups of silica nanoparticles.

Fig. 2. a) ATR-FTIR spectra with cross-linking procedures of composite membranes (inset represents the magnified peaks) b) DSC c) TGA of SPAEK, SPAEK/silica and CL-SPAEK/ silica membranes.

3.2. Water uptake and proton conductivity Fig. 3a shows the water uptake of SPAEK, SPAEK/silica and CLSPAEK/silica membranes with increasing temperature. Water uptake of SPAEK and SPAEK/silica is 51.9% and 55.9% at 25 °C, respectively. The higher water uptake of SPAEK/silica is attributed to the hydrophilic character of the silica surface providing a higher capability of water molecules uptake compared to SPAEK. With increasing temperature, the water uptake of SPAEK and SPAEK/silica increased due to the thermal relaxation of the polymer chains. However, CL-SPAEK/silica shows the slightly increased value with increasing temperature, originating from the cross-linked structure in the composite membrane. The concentration of silica nanoparticles in the membrane was controlled 570

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Fig. 3. a) Water uptake b) volumetric expansion with water swelling c) DSC d) the state of water of SPAEK, SPAEK/silica and CL-SPAEK/silica membranes.

Proton transport through a membrane is described by two primary conductive mechanisms: Grotthuss mechanism and Vehicular mechanism [37]. Proton diffusion within the membrane is suggested with bound and free water through the Grotthuss mechanism and Vehicular mechanism, respectively [38]. In the Grotthuss mechanism, a proton in the ionic form of H3O+ transports to the neighboring molecules by molecular rearrangement. But in the Vehicular mechanism, the protons are carried with the water molecules forming hydrophilic ions such as H3O+, H5O2+ and H9O4+ [28]. Fig. 4 shows proton conductivity of PEMs at 70 °C as a function of RH. The controlled amount of silica nanofillers (5 wt% to polymer) in the SPAEK matrix contributed to the highest proton conductivity by higher bound water contents and effective formation that prevents the fatal aggregation of nanoparticles.

from 0 to 10 wt% to a polymer as shown in Fig. S6a. Water uptake of the composite membrane increased up to 7 wt% but decreased to 10 wt % with silica content due to severe aggregation of nanoparticles and uneven morphology of composite membranes. All of the cross-linked membranes with various silica contents had lower water uptake than the non-cross-linked membranes except for the membrane containing 10 wt% silica. After water swelling in the membranes, the volumetric expansion of the composite membranes was measured, resulting in a reduced expansion rate of SPAEK/silica and CL-SPAEK/silica. (Fig. 3b) Silica nanoparticles dispersed in the polymer matrix enhanced the dimensional stability of the membrane for water swelling, which originated from the reduction of areal expansion. Furthermore, the crosslinked structure of CL-SPAEK/silica improved the temperature-independent mechanical stability for humidity, which resulted from the area, thickness and volumetric stability as shown in Fig. S6. To further investigate the physical states of water in the membranes, DSC thermograms were investigated near 0 °C. The state of water in PEMs can be classified into free and bound water [36]. Free water behaves like bulk water without intimate bond to polymer chains and represents the sharp melting point at 0 °C by DSC analysis (Fig. 3c). The amount of free water within the membranes was calculated with the endothermic enthalpy change of the melting behavior at 0 °C according to Eq. (3). Bound water have a strong interaction with the polar groups of the polymer chain such as sulfuric acid and hydroxyl groups, thus cannot freeze under 0 °C without an endothermic heat flow. Although SPAEK and SPAEK/silica have higher water uptake than CL-SPAEK/ silica, CL-SPAEK/silica presented the highest fraction of bound water as shown in Fig. 3d and Table S1. This implication is originated from the well-dispersed silica in the polymer matrix and the cross-linked structure between the polymers and silica by sulfuric acid. The high amount of bound water in CL-SPAEK/silica is significant for improving proton transport under low humidity conditions.

Fig. 4. Proton conductivity of SPAEK, SPAEK/silica and CL-SPAEK/silica membranes.

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Fig. 5. AFM surface phase images and TEM cross-section images of a, d) SPAEK, b, e) SPAEK/silica and c, e) CL-SPAEK/silica membranes. (Scale bars represent 100 and 20 nm for AFM and TEM images, respectively.)

(Fig. S7) The proton conductivity of CL-SPAEK/silica was 3.06 mS cm−1, which was higher than SPAEK (0.32 mS cm−1) and SPAEK/silica (2.07 mS cm−1) under 30% RH. This suggests that the cross-linked structure between SPAEK and silica encouraged proton transport, which induced an increase in the interfacial networks between the polymers and silica and held more bound water. 3.3. Morphological characterization The enhanced proton conductivity under low RH could be investigated with morphological factors, such as connectivity of hydrophilic channels and phase separation between hydrophilic and hydrophobic moieties for proton transport. In the absence of free water, the nanostructure of membranes is one of the most significant factors affecting proton transport. The nano-scale hydrophilic/hydrophobic phase separation of SPAEK, SPAEK/silica and CL-SPAEK/silica was characterized by the tapping mode AFM and TEM for the surface and cross-sectional morphology of the membranes, respectively. (Fig. 5) SPAEK showed weak phase separation morphology, where the relatively dark regions represented the presence of hydrophilic domains and the light regions were assigned to the hydrophobic domain. The increased hydrophilic domains of the membranes were observed because of the hydrophilic silica nanoparticles in SPAEK/silica, and an irregular phase separation morphology and disconnected hydrophilic proton pathways were detected by the aggregation of inorganic nanofillers. Whereas CL-SPAEK/silica membrane had a well assorted and a clear nano-phase separated morphology with the cross-linked structure between the polymer chains and silica nanoparticles. Well-dispersed silica nanoparticles in the SPAEK polymer matrix with the cross-linked structure attributed to the clearly separated morphology resulting in high connectivity of hydrophilic regions. The homogeneous morphology of CL-SPAEK/silica can be contributory to the increase of proton conductivity under low humidity conditions. To quantify the correlation lengths, the molecular morphology of the phase separation in membranes was investigated by SAXS analysis under 30% RH. As shown in Fig. 6, the SAXS pattern for SPAEK exhibited the distinguished peak at the lower scattering vector, q value (q

Fig. 6. SAXS analysis of a) SPAEK b) SPAEK/silica and c) CL-SPAEK/silica membranes under 30% RH.

= 0.057 Å−1), representing the bicontinuous phase separation distance between the hydrophilic and hydrophobic moieties (d = 11.0 nm). Although the expanded inter-domain distance (d = 12.0 nm) was observed in SPAEK/silica, the scattering pattern for the phase separation was much broader than that of SPAEK due to the irregular and uneven distribution of hydrophilic domains as shown in AFM phase image of Fig. 5b. However, CL-SPAEK/silica exhibited a clear peak, which indicates more regularly phase separated morphology than SPAEK and SPAEK/silica. The Braggs distance d was increased to 13.4 nm with the uniformly dispersed silica formation in the polymer matrix by the crosslinked structure between the polymer chains and silica, which is in good agreement with the high proton conductivity under low humidity conditions. 3.4. PEMFC performance and durability Increased proton transport by higher water maintenance and well572

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Fig. 8. Strain-stress curve of SPAEK, SPAEK/silica and CL-SPAEK/silica membranes.

Mechanical stability and durability of PEM for fuel cell applications were investigated by alternating completely dry N2 (0% RH) and fully humidified N2 (150% RH) at 2 min intervals. The dimensional change of the membrane by RH cycle induces the mechanical stress during the fuel cell operation. The volumetric change of CL-SPAEK/silica was smaller than that of the other membranes and consistent with increasing temperature due to the molecular cross-linking between SPAEK polymer and silica surface. (Fig. 3b and Table S2) Stable swelling behavior, higher tensile strength and Young's modulus of CLSPAEK/silica were clearly performed in the mechanical durability test of PEMFC. In Fig. 9, SPAEK, SPAEK/silica and CL-SPAEK/silica membranes represented the hydrogen crossover current density of 1.0, 0.58 and 0.55 mA cm−2 before the wet-dry cycle, respectively. After 18 and 164 cycles, the hydrogen crossover current density of SPAEK and SPAEK/silica rapidly increased, indicating a mechanical failure of thin membranes caused by high volumetric expansions for swelling and low mechanical stability. However, CL-SPAEK/silica showed the constant current density over 600 cycles, corresponding to the results of good dimensional stability and mechanical properties without any reinforced substrates [39].

Fig. 7. a) Polarization characteristics b) Nyquist plots of impedance for SPAEK, SPAEK/ silica and CL-SPAEK/silica membranes under 32% RH.

developed phase separation morphology of CL-SPAEK/silica was investigated in a fuel cells system with electrochemical performance. Fig. 7a shows the single cell i-V curves of PEMFCs for SPAEK, SPAEK/ silica and CL-SPAEK/silica membranes under 32% RH and ambient pressure. Among these MEAs, the highest electrochemical performance of the MEA based on CL-SPAEK/silica was found to be 0.919 A cm−2 at a cell voltage of 0.6 V at 70 °C. The improved power performance of PEMFC with CL-SPAEK/silica corresponded to the elevated proton conductivity, facilitated by the developed nanostructure of the membrane with the cross-linked structure under low humidity conditions. To evaluate ohmic resistance (Rohm) of membranes, the Nyquist plots of impedance were obtained at 25 A by using the electrochemical impedance spectroscopy. As shown in Fig. 7b, CL-SPAEK/silica exhibited the lowest Rohm of 121 mΩ cm2 compared to that of SPAEK and SPAEK/ silica, 771 and 425 mΩ cm2, respectively. The cross-linked structure of the polymer-inorganic composite membrane enhanced not only proton transport but also membrane stability and single cell durability. Mechanical properties such as tensile strength, elongation at break and Young's modulus of the prepared membranes in 30% RH are shown in Fig. 8 and Table S2. Although SPAEK/silica showed higher Young's modulus than SPAEK, it had relatively low tensile strength and elongation at break due to the aggregation of silica nanoparticles in the polymer matrix. CL-SPAEK/silica, however, provided increased tensile strength and Young's modulus but decreased elongation at break among the membranes. The crosslinked structure enhanced the interfacial interaction between the organic-inorganic phases and increased the dispersion of silica nanoparticles in the polymer matrix, leading to higher mechanical properties of the membrane such as Young's modulus and tensile strength.

4. Conclusion The novel proton exchange membrane was developed using silica nanoparticles cross-linked with benzhydrol groups of R-SPAEK. The silica nanoparticles of CL-SPAEK/silica were well-dispersed in the SPAEK matrix due to the high interfacial interaction between the surface of nanoparticles and polymer chains. As a result, CL-SPAEK/silica

Fig. 9. Durability test of SPAEK, SPAEK/silica and CL-SPAEK/silica membranes by wetdry cycles of the MEAs.

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exhibited a morphological transformation such as clear hydrophilic/ hydrophobic phase separation and highly connected hydrophilic nanochannels of the synthesized block copolymer. The morphology of CLSPAEK/silica exhibited enhanced electrochemical performance under low humidity conditions compared to SPAEK and SPAEK/silica. The optimized CL-SPAEK/silica showed higher proton conductivity (3.06 mS cm−1 at 70 °C under 30% RH) and single cell performance (0.919 A cm−2 at 0.6 V under 32% RH) than those of SPAEK (0.32 mS cm−1 and 0.736 A cm−2 at 0.6 V) and SPAEK/silica (2.07 mS cm−1 and 0.818 A cm−2 at 0.6 V). Moreover, excellent durability of the composite membranes was obtained during wet-dry cycle.

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