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. 2022 Dec 22;8(1):1121–1130. doi: 10.1021/acsomega.2c06514

Effects of a Nanophase-Separated Structure on Mechanical Properties and Proton Conductivity of Acid-Infiltrated Block Polymer Electrolyte Membranes under Non-Humidification

Takato Kajita †, Haruka Tanaka †, Yumiko Ohtsuka †, Tsuyoshi Orido †, Atsushi Takano †, Hiroyuki Iwamoto ‡, Albert Mufundirwa ‡, Hideto Imai §, Atsushi Noro †,∥,*
PMCID: PMC9835166  PMID: 36643438

Abstract

graphic file with name ao2c06514_0011.jpg

Acid-infiltrated block polymer electrolyte membranes adopting a spherical or lamellar nanophase-separated structure were prepared by infiltrating sulfuric acid (H2SO4) into polystyrene-b-poly(4-vinylpyridine)-b-polystyrene (S–P–S) triblock copolymers to investigate the effects of its nanophase-separated structure on mechanical properties and proton conductivities under non-humidification. Lamellae-forming S–P–S/H2SO4 membranes with a continuous hard phase generally exhibited higher tensile strength than sphere-forming S–P–S/H2SO4 membranes with a discontinuous hard phase even if the same amount of Sa was infiltrated into each neat S–P–S film. Meanwhile, the conductivities of lamellae-forming S–P–S/H2SO4 membranes under non-humidification were comparable or superior to those of sphere-forming S–P–S/H2SO4 membranes, even though they were infiltrated by the same weight fraction of H2SO4. This result is attributed to the conductivities of S–P–S/H2SO4 membranes being greatly influenced by the acid/base stoichiometry associated with acid–base complex formation rather than the nanophase-separated structure adopted in the membranes. Namely, there are more free H2SO4 moieties that can release free protons contributing to the conductivity in lamellae-forming S–P–S/H2SO4 membranes than sphere-forming S–P–S/H2SO4, even when the same amount of H2SO4 was infiltrated into the S–P–S.

1. Introduction

Fuel cells generate electrical energy through an electrochemical reaction between hydrogen and oxygen gases, producing only water as a product; therefore, these cells are promising as clean power generation systems.1,2 Polymer electrolyte fuel cells (PEFCs) using proton-conductive polymer electrolyte membranes (PEMs) consisting of the important components in fuel cells are operated at relatively low temperatures, approximately 70–90 °C,3−12 as compared with other fuel cells based on other electrolytes such as solid oxides and phosphoric acid, which are typically operated at 200 °C or higher; thus, PEFCs can be used in a limited space such as fuel cell vehicles6 and household fuel cell cogeneration systems.7 The most well-known proton-conductive PEM is a perfluorosulfonic acid polymer such as Nafion.8,9 Nafion exhibits high proton conductivities of 0.05 S cm–1 at 80 °C and 70 %RH, and 0.1 S cm–1 at 80 °C and 90 %RH,10 where the fluorinated phase contributing to the mechanical strength of a membrane is separated from the proton-conductive sulfonic acid group/water mixed phase at a microscopic scale.11,12 However, Nafion exhibits almost no conductivity under non-humidification; therefore, a PEFC using Nafion requires a system that controls not only temperature but also humidity.

To address the abovementioned problem, PEMs exhibiting proton conductivities even under non-humidification have been widely studied.13−31 A representative example of proton-conductive PEMs that can be operated under non-humidification is a phosphoric acid-doped polybenzimidazole (PBI) membrane, where PBI is a type of super-engineered plastic with aromatic rings in the backbone. At an early stage in development of phosphoric acid-doped PBI membranes, the conductivity was ∼0.04 S cm–1 at a relatively high temperature of 190 °C under non-humidification for the membrane with 61 wt % phosphoric acid of a low-molecular-weight electrolyte.13 Further development allowed the preparation of PBI with 93 wt % phosphoric acid exhibiting a high conductivity (0.16 S cm–1) at 120 °C under dry conditions via a sol–gel process by heating the monomer of PBI with polyphosphoric acid.16 As other PEMs, an ionic liquid (IL)-doped PBI membrane17,19 and an IL-doped sulfonated polyimide membrane20 have also been reported; however, they exhibited lower conductivities than the phosphoric acid-doped PBI membranes because the conductivities of neat phosphoric acid32 under non-humidification are typically higher than those of neat ILs.33

Vinyl polymer-based PEMs have also been developed as further alternatives of proton-conductive PEMs under non-humidification.34−43 For example, Narayanan and coworkers prepared a PEM by infiltrating sulfuric acid (H2SO4) or phosphoric acid into poly(4-vinylpyridine) with a basic group,35 although the conductivity of the PEM with 48 wt % H2SO4 was a mere ∼2 × 10–4 S cm–1 at 140 °C under non-humidification. To attain higher conductivities, the amount of infiltrated H2SO4 contributing to high proton conduction44 should be larger, but the mixture of poly(4-vinylpyridine) and a larger amount of H2SO4 easily becomes a fluid because the glass transition temperature (Tg) of such a mixture tends to be lower than the operation temperature. Therefore, to keep the mixture of poly(4-vinylpyridine) and H2SO4 solid, we have recently prepared a proton-conductive PEM by chemically cross-linking poly(4-vinylpyridine), which can be swollen with a large amount of H2SO4, exhibiting conductivities over 0.1 S cm–1 at around 100 °C under non-humidification.42 A highly proton-conductive PEM under non-humidification has also been prepared by infiltrating H2SO4 into a nanophase-separated polystyrene-b-poly(4-vinylpyridine)-b-polystyrene (S–P–S) triblock copolymer that can be prepared by living addition polymerization45,46 of vinyl monomers. The effect of the acidity of the specific acid (H2SO4 or phosphoric acid) on the conductivity of the S–P–S/acid membranes was evaluated as well, revealing that H2SO4 was more effective in preparing highly conductive acid-infiltrated PEMs than phosphoric acid.43

Depending on the composition of block polymers, nanophase-separated structures of block polymers vary from a spherical structure with isolated discontinuous spheres in a continuous matrix to a lamellar structure with stacked continuous layers.47,48 If nanophase-separated structures formed in PEMs of acid-infiltrated block polymers are different, the PEMs should also exhibit different conductivities and mechanical properties. To date, the effects of the nanophase-separated structure, probably one of the key factors determining the properties of PEMs of acid-infiltrated block polymers, has not been sufficiently investigated. Therefore, in this study, block polymer-based PEMs infiltrated with H2SO4 have been prepared by using two S–P–S triblock copolymers adopting different nanophase-separated structures with different phase continuities (Figure 1) to investigate the effects of a phase-separated structure on the mechanical properties and conductivities of the PEMs under non-humidification. The effects of H2SO4 weight fraction in the PEMs are also investigated. The study uses H2SO4 as an infiltrated acid since block polymer-based PEMs infiltrated with H2SO4 are expected to exhibit higher conductivities under non-humidification than block polymer-based PEMs infiltrated with phosphoric acid when the amount of infiltrated acid is the same.43 It should also be noted that S–P–S triblock copolymers were used instead of S-P diblock copolymers because S–P–S triblock copolymer-based membranes exhibit better mechanical properties than S–P diblock copolymer-based membranes due to formation of polymer network by bridging P center blocks with the hard S domains in the S–P–S triblock copolymer-based membranes.49

Figure 1.

Figure 1

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Schematic illustration of the preparation of acid-infiltrated block polymer electrolyte membranes adopting different nanophase-separated structures: (a) Spherical structure. (b) Lamellar structure.

2. Experimental Section

2.1. Synthesis and Characterization of Triblock Copolymers

Two S–P–S triblock copolymers with almost the same molecular weight (Mn,total), but with different ϕS values, were synthesized by reversible addition-fragmentation chain transfer (RAFT) polymerization,50,51 as previously reported.43 First, styrene purified by passage through an activated alumina column was polymerized by using a bifunctional RAFT agent and 2,2′-azobis(isobutyronitrile) (AIBN) at 130 °C. After polymerization, the product of polystyrene (abbreviated as S) was purified by reprecipitation with methanol. Then, S–P–S was synthesized by polymerizing 4-vinylpyridine by using the S with the RAFT agent residue as a macro-RAFT agent at 80 °C. Two S–P–S with smaller and larger S volume fractions were coded as S–P–S(s) and S–P–S(l), respectively. Molecular characteristics of the two types of S–P–S were determined as previously reported,43 and Table 1 summarizes molecular characteristics of the two synthesized variants of S–P–S. To estimate the molecular weight distribution of the two S–P–S, gel permeation chromatography was also conducted by using an HPLC system (HPLC pump: Shimadzu LC-20AD; column oven: Shimadzu CTO-20A; RI detector: Shimadzu RID-10; eluent solvent: N,N-dimethylformamide; flow rate: 1.0 mL min–1) equipped with three TSK gel G4000HHR columns (Tosoh Corp.) at 40 °C (see GPC chromatograms in Figure S1). The degree of polymerization, number-average molecular weight, and composition of the two S–P–S were determined by 1H NMR spectroscopy with an Ascend 500 MHz (Bruker Corp.). The solvent used for spectroscopy was deuterated chloroform. Figure 2 shows the 1H NMR spectra of S–P–S(s) and S–P–S(l), respectively. The integral of the signals originating from three phenyl protons (positions a and b) in the spectrum of S–P–S(l) was obviously larger than in the spectrum of S–P–S(s), indicating that the S fraction of S–P–S(l) was larger than that of S–P–S(s). See also Figure S2 regarding how to determine the molecular weights of S–P–S(s), S–P–S(l), and the precursor S.

Table 1. Molecular Characteristics of Neat S–P–S.

sample Mn,Sa Mn,totalb Mw/Mnc φSd
S–P–S(s) 14,000 180,000 1.6 0.08
S–P–S(l) 33,000 169,000 1.7 0.20
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a

Number-average molecular weight of a precursor S determined by 1H NMR.

b

Total number-average molecular weight of an S–P–S calculated by using a molecular weight of the precursor S and the molar fraction of the S–P–S estimated from 1H NMR.

c

Molecular weight distribution determined by GPC. The molecular weight was calibrated by using polystyrene standards.

d

Volume fraction of S blocks calculated by using the molar fraction of S blocks in S–P–S estimated from 1H NMR and the room-temperature bulk densities of component polymers, i.e., 1.05 g cm–3 for S and 1.17 g cm–3 for P.52

Figure 2.

Figure 2

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1H NMR spectra of S–P–S(s) (top) and S–P–S(l) (bottom).

2.2. Preparation of S–P–S/H2SO4 Membranes

S–P–S/H2SO4 membranes were prepared by infiltrating H2SO4 into neat S–P–S as previously reported.43 First, neat S–P–S films were prepared by a solution casting method with pyridine as a solvent, followed by vacuum drying at 50 °C. The neat S–P–S film was immersed into a solution of H2SO4 in methanol. Note that no phase transition presumably occurs after infiltration of H2SO4 because methanol dissolves P and H2SO4 but does not dissolve S. After slowly evaporating the methanol from the solution at 50 °C for 12 h, methanol was added again to homogeneously infiltrate the acid, followed by solution casting at 50 °C for 36 h and vacuum drying at 50 °C for 24 h. The weight content of H2SO4 in the S–P–S/H2SO4 PEMs ranged from 50 to 80 wt %. S–P–S/H2SO4 membranes with less than 50 wt % H2SO4 were not prepared because conductivity values of the membranes with less than 50 wt % H2SO4 are assumed to be very low under non-humidification according to previous reports.42,43 For S–P–S(l)/H2SO4 with 80 wt % H2SO4, all the H2SO4 used for membrane preparation cannot be infiltrated in S–P–S(l). This is attributed to both the smaller fraction of H2SO4-retaining P of S–P–S(l) than that of S–P–S(s) and the larger interfacial area between the S phase and the P/H2SO4 mixed phase in lamellae-forming S–P–S(l) than that of sphere-forming S–P–S(s), where S and H2SO4 contact each other. Probably due to such reasons, the easier leaching of H2SO4 from S–P–S(l)/H2SO4 than S–P–S(s)/H2SO4 was observed. The homogeneous PEMs prepared are coded as S–P–S(X)/H2SO4(wH2SO4), where X is s or l, and wH2SO4 represents the weight percent of H2SO4 in the PEM. The molar ratio of H2SO4 to a pyridyl group in S–P–S, termed the acid doping level (ADL), can be calculated from eq 1:

2.2. 1

where nH2SO4 and MH2SO4 are the molar amount and molecular weight (98 g mol–1) of H2SO4, respectively, while nP, wP, and MP,monomer are the molar amount, weight amount, and molecular weight (105 g mol–1), respectively, of the 4-vinylpyridine monomer unit in the PEM. The values wS-P-S and w’P are the weight fraction of S–P–S in the PEMs and weight fraction of P in neat S–P–S, respectively.

2.3. Measurements

Transmission electron microscopy (TEM) observation was carried out for the two types of neat S–P–S. The sample specimen of neat S–P–S was embedded into epoxy resin, followed by preparation of ultrathin microtome sections with a thickness of ∼80 nm in a wet condition. The sections were stained with iodine (I2) vapor at 50 °C for 50 min. The instrument used for TEM observation was JEM-2100 Plus (JEOL Ltd.), and the acceleration voltage was 200 kV. Note that TEM observation for PEMs infiltrated with H2SO4 was not performed because highly acidic H2SO4 can damage the TEM instrument if H2SO4 is leached out from the PEM during observations.

SAXS measurements were performed at room temperature under an argon atmosphere to acquire quantitative nanostructural information of neat S–P–S and S–P–S/H2SO4 membranes. To prevent the membranes from exposure to an air atmosphere, hard samples were enclosed in a glass capillary with a diameter of 1.5 mm whereas soft samples were sandwiched between Kapton films with a thickness of ∼7.5 μm. The measurements were carried out at room temperature using BL-40B2 of the SPring-8 facility, Hyogo, Japan, at a wavelength of 0.15 nm53,54 and a camera length of 6.25 m with the Pilatus 2M detector.

Tensile tests for the strip-shaped PEMs with a dimension of 20 mm × 4 mm × 0.5 mm were carried out with an Autograph AGS-X (Shimadzu) equipped with a 50 N load cell and 50 N pneumatic flat grips. The tests were conducted at an initial between-jigs distance of about 10 mm with an initial strain rate of approximately 0.10 s–1 (an elongation rate of 1.0 mm s–1) at room temperature.55,56

The conductivities of the PEMs were determined as previously reported42,43 by alternating current impedance spectroscopy with a potentio/galvanostat VSP-300 (BioLogic Science Instruments) in the frequency range of 1 × 100 to 7 × 106 Hz at a signal amplitude of 50 mV by using a two-probe method. A test cell was prepared by fixing the distance (l) between the platinum electrodes at 7 mm and using the test specimen with sectional area (A) of ∼2 mm2 (see also the schematic of the impedance spectroscopy setup in Figure S4). The temperature and humidity were controlled in the benchtop-type environmental chamber SH-242 (ESPEC Corp.) within a temperature range of 20–95 °C below the Tg of polystyrene (∼100 °C) at no humidification, where the humidity in the chamber at all temperatures was determined to be close to 0 %RH by the thermohygrometer Testo 645 (Testo SE & Co. KgaA). To estimate the conductivity (σDC) by using eq 2, a bulk resistance (R) of the PEMs was evaluated by reading the extrapolated value of the plot on the horizontal axis (nonzero Z′ intercept in the Nyquist plot39,43,57), where Z′ is the real part of the complex impedance Z = Z′ – iZ″ (see also the Nyquist plots of S–P–S/H2SO4 membranes in Figure S5).

2.3. 2

3. Results and Discussion

3.1. Nanophase-Separated Structure of Neat S–P–S and PEMs

Typical TEM images of neat S–P–S(s) and neat S–P–S(l) are displayed in Figure 3a,b, respectively, where the S phase appears brighter while the P phase looks darker because of I2 vapor staining.58,59 A brighter spherical S phase with discontinuity in the darker P phase as a continuous matrix was observed in the TEM image of neat S–P–S(s). Based on the TEM image, domain spacing (D) of neat S–P–S(s) was estimated to be ∼25 nm. On the other hand, neat S–P–S(l) showed a lamellar structure with D ∼ 45 nm composed of a thin S-layered phase and a thick P-layered phase. The lamellar structure was formed in neat S–P–S(l) in spite of a relatively small φS, probably due to the relatively large molecular weight distribution of S–P–S(l) (Mw/Mn = 1.7).60−62

Figure 3.

Figure 3

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TEM images of (a) neat S–P–S(s) and (b) neat S–P–S(l).

SAXS measurements were also carried out to investigate quantitative nanostructural information of the two neat S–P–S. SAXS profiles of neat S–P–S(s) and neat S–P–S(l) are shown in Figure 4a,b, respectively. The profile of S–P–S(s) at the bottom of Figure 4a (profile with black lines) showed a peak at a scattering vector q (= 4π sin θ/λ) of 0.23 nm–1, where λ and 2θ are the wavelength of X-rays and the scattering angle, respectively. Taking account also of both the TEM results of S–P–S(s) and a general self-assembly manner of block polymers, the S spherical domains in neat S–P–S(s) are assumed to be packed in a bcc lattice,63,64 and D of neat S–P–S(s) is estimated to be ∼24 nm by using the first peak q position (q1) of the SAXS profile and the eq D = (3/4)1/2 × 2π/q1,59 which roughly agreed with the TEM result (see also q1 and D values in Table S1). Note that the second peak was not clearly observed in the profile of neat S–P–S(s), indicating spherical domains were arranged in a poorly ordered manner. On the other hand, the SAXS profile of neat S–P–S(l) at the bottom of Figure 4b clearly showed integer-order peaks relative to the first peak (q1 = 0.13 nm–1), indicating formation of the lamellar structure with D (= 2π/q159) ∼49 nm in S–P–S(l). This result was also consistent with D estimated from the TEM image of neat S–P–S(l).

Figure 4.

Figure 4

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SAXS profiles of (a) S–P–S(s)/H2SO4 and (b) S–P–S(l)/H2SO4.

SAXS measurements for S–P–S/H2SO4 membranes were also conducted to acquire information about the nanostructures of S–P–S infiltrated with H2SO4. Figure 4a shows SAXS profiles of S–P–S(s)/H2SO4 membranes as well as the profile of neat S–P–S(s). In a profile of the S–P–S(s)/H2SO4 membrane with 50 wt % H2SO4, scattering intensities at √2q1 and √3q1 were relatively strong, compared with the weaker intensities in a profile of neat S–P–S(s). This outcome may be attributed to improved ordering of the spherical domain packing resulting from stronger segregation between hydrophobic S and hydrophilic P/H2SO4 mixed phases than between S and P phases.53 Furthermore, as the amount of H2SO4 in S–P–S(s)/H2SO4 membranes increased, the first peak also shifted to lower q, while the profile pattern was mostly retained (see also q1 values in Table S1). These results indicated that domain spacing (D) became larger without morphology transition by selectively infiltrating H2SO4 into the P matrix phase (see also D values in Table S1). Similarly, all the profiles of S–P–S(l)/H2SO4 membranes in Figure 4b showed the integer-order peak as neat S–P–S(l), indicating that the lamellar structure remained in the S–P–S(l)/H2SO4 membranes even in the presence of infiltrating H2SO4. In addition to S–P–S(s)/H2SO4, the first peak of S–P–S(l)/H2SO4 membranes also shifted to the lower q side with the addition of more Sa to the neat S–P–S(l). Table 2 summarizes D of S–P–S/H2SO4 membranes estimated from the SAXS profiles, and D is plotted against wH2SO4 in Figure 5. For both S–P–S(s)/H2SO4 and S–P–S(l)/H2SO4, D seems to depend more strongly on wH2SO4 in the range from 50 to 80 wt % compared with the range below 50 wt %. This dependence is probably due to acid–base complexation65 between H2SO4 and the pyridyl group in the P/H2SO4 mixed phase. At the ADL below unity or wH2SO4 below approximately 50 wt %, P block chains largely shrink by strong ionic interactions attributed to acid–base complexation. On the other hand, when the ADL was larger than unity, the P chains were swollen because an excess amount of H2SO4 behaves like a plasticizer, leading to a rapid D increase. In addition, the D of S–P–S(l)/H2SO4 was more strongly dependent on wH2SO4 compared with D of S–P–S(s)/H2SO4. This result is probably due to the difference in the nanostructure dimensionality and the swelling behavior. Note that the effect of one-dimensional swelling of polymer chains in S–P–S(l)/H2SO4 resulting in an increase in D was larger than that of three-dimensional swelling observed in S–P–S(s)/H2SO4.

Table 2. Properties of Neat S–P–S and S–P–S/H2SO4 Membranes.

sample wH2SO4a(wt %) ADLb Dc (nm) EYd (MPa) σmaxe (MPa) εbf (%) Tgg (°C) σDC,95 °Ch(S cm–1)
neat S–P–S(s) 0 0 24 n.d.i n.d.i n.d.i 138 n.d.i
S–P–S(s)/H2SO4(50) 50 1.2 30 12 1.7 58 52 1.5 × 10–4
S–P–S(s)/H2SO4(60) 60 1.7 35 0.60 0.43 580 –61 0.012
S–P–S(s)/H2SO4(70) 70 2.7 38 0.17 0.18 510 –75 0.062
S–P–S(s)/H2SO4(80) 80 4.6 45 0.039 0.032 210 –80 0.14
neat S–P–S(l) 0 0 49 n.d.i n.d.i n.d.i 141 n.d.i
S–P–S(l)/H2SO4(50) 50 1.3 57 43 2.7 47 31 0.0011
S–P–S(l)/H2SO4(60) 60 2.0 84 33 1.4 120 –60 0.039
S–P–S(l)/H2SO4(70) 70 3.1 91 4.2 0.75 99 –75 0.061
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a

Weight fraction of H2SO4 in a sample.

b

Acid doping level calculated from eq 1.

c

Domain-spacing estimated from SAXS (see also Table S1).

d

Young’s modulus estimated from the slope from the strain range from 1 to 3%.

e

Tensile strength.

f

Elongation at break.

g

Glass transition temperature of the P or P/H2SO4 mixed phase determined by DSC (Figure S3). Note that Tg of the S phase was not clearly observed by DSC due to the small volume fraction of S in S–P–S/H2SO4. However, Tg derived from S should exist because the S phase was observed by TEM and SAXS.

h

Conductivity at 95 °C under non-humidification.

i

Not determined.

Figure 5.

Figure 5

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Domain spacing of S–P–S/H2SO4 as a function of wH2SO4.

3.2. Mechanical Properties of S–P–S/H2SO4 Membranes

To investigate the effect of the phase-separated structure of the block polymer-based PEMs on the mechanical properties, tensile tests for S–P–S/H2SO4 membranes were performed. Figure 6a,b shows tensile stress–strain curves of a series of S–P–S(s)/H2SO4 and S–P–S(l)/H2SO4 membranes, respectively, when measured at room temperature. Table 2 also summarizes Young’s modulus (EY), tensile strength (σmax), and elongation at break (εb) of the S–P–S/H2SO4 membranes. The EY, σmax, and εb values of the S–P–S(s)/H2SO4(50) membrane were 12 MPa, 1.7 MPa, and 58%, respectively, indicating a brittle plastic-like behavior.66 In contrast, the S–P–S(s)/H2SO4(60) membrane exhibited a lower EY (0.60 MPa) and σmax (0.43 MPa) but a larger εb (580%), which indicated that the membrane shows an elastomeric behavior.55 As the H2SO4 content increased further, EY, σmax, and even εb decreased (Table 2). At wH2SO4 = 50 wt %, that is, ADL is close to 1, Tg of the P/H2SO4 mixed phase was much higher than room temperature due to formation of a hard acid–base complex of H2SO4 and the pyridyl group. Since Tg of the S phase in S–P–S(s)/H2SO4(50) was approximately 100 °C, also higher than room temperature, the S–P–S(s)/H2SO4 (50) membrane behaved like a brittle plastic at room temperature. On the other hand, when wH2SO4 exceeded 60 wt %, Tg of the P/H2SO4 mixed phase became lower than room temperature by the presence of an excessive amount of H2SO4 serving as a plasticizer in the mixed phase, causing the membranes to behave like an elastomer.

Figure 6.

Figure 6

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Tensile stress–strain curves of a series of (a) S–P–S(s)/H2SO4 and (b) S–P–S(l)/H2SO4 membranes.

The S–P–S(l)/H2SO4(50) membrane also exhibited a brittle plastic-like behavior similar to the S–P–S(s)/H2SO4(50) membrane. Similarly, as wH2SO4 increased, S–P–S(l)/H2SO4 membranes also behaved like an elastomer; namely, EY and σmax of S–P–S(l)/H2SO4 membranes decreased and its εb increased as well as the S–P–S(s)/H2SO4 membranes (Table 2). However, it should also be noted that the elastomer-like S–P–S(l)/H2SO4 exhibited much better mechanical strength compared with the elastomer-like S–P–S(s)/H2SO4 membranes with the same wH2SO4. For example, EY and σmax of S–P–S(l)/H2SO4(70) were 25 times and 4.2 times, respectively, higher than the same properties of the S–P–S(s)/H2SO4(70) membrane. Since the morphology of the hard S phase in S–P–S(l)/H2SO4 membranes is a planar layer with a higher degree of continuity while that of the hard S phase in S–P–S(s)/H2SO4 membranes is a discontinuous sphere, the continuity of the hard S-layered phase in S–P–S(l)/H2SO4 membranes probably contributes greatly to the higher mechanical strength. Therefore, if a hard three-dimensionally continuous S phase such as gyroid phase is formed in the membrane, such a membrane can exhibit higher mechanical strength than that of lamella-forming S–P–S(l)/H2SO4 membranes with a two-dimensionally continuous S phase.

3.3. Conductivity of S–P–S/H2SO4 Membranes under Non-Humidification

Figure 7 shows the plot of σDC for S–P–S/H2SO4 membranes against the reciprocal of the absolute temperature (T) (see also the Nyquist plots for S–P–S/H2SO4 membranes in Figure S5). The σDC of the S–P–S(s)/H2SO4 membranes under non-humidification increased with increasing temperature for all samples. The σDC of S–P–S(s)/H2SO4 membranes was also fitted by the Vogel–Fulcher–Tammann (VFT) equation (eq S1), which is useful for fitting σDC influenced by Tg(39) (see also Table S2 and Figure S6 for the fitting parameters). The σDC of the S–P–S(s)/H2SO4(50) membrane exhibited a low σDC of 1.5 × 10–4 S cm–1 under non-humidification even at the highest temperature of 95 °C adopted in this study, while the S–P–S(s)/H2SO4(60) membrane showed a moderately high σDC of 1.2 × 10–2 S cm–1, two orders of magnitude larger at the same temperature under dry conditions. As the H2SO4 content increased further, membranes such as S–P–S(s)/H2SO4(70) and S–P–S(s)/H2SO4(80) exhibited much higher values of σDC of 6.2 × 10–2 and 1.4 × 10–1 S cm–1, respectively. These results are consistent with our previous study.42,43 Similar to S–P–S(s)/H2SO4 membranes, the σDC values of S–P–S(l)/H2SO4 membranes with the lamellar structure under non-humidification also increased with increasing wH2SO4, but notably, the absolute value of σDC of S–P–S(l)/H2SO4 was generally higher than that of S–P–S(s)/H2SO4 under non-humidification when wH2SO4 was the same. Namely, the S–P–S(l)/H2SO4(50) membrane exhibited the σDC of 1.6 × 10–3 S cm–1 under non-humidification, a conductivity that is one order of magnitude higher than that of the S–P–S(s)/H2SO4(50) membrane at 95 °C (1.5 × 10–4 S cm–1). The σDC of S–P–S(l)/H2SO4(60) was 4.6 × 10–2 S cm–1, a conductivity approximately three times higher than that of S–P–S(s)/H2SO4(60) at the same temperature (1.2 × 10–2 S cm–1). The σDC of S–P–S(l)/H2SO4(70) was 6.1 × 10–2 S cm–1, which is comparable to that of S–P–S(s)/H2SO4(70).

Figure 7.

Figure 7

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Temperature dependence of conductivity of a series of (a) S–P–S(s)/H2SO4 and (b) S–P–S(l)/H2SO4 membranes under non-humidification. Solid lines are fits by the VFT equation.

To compare the σDC of the S–P–S(l)/H2SO4 membranes with that of S–P–S(s)/H2SO4 when wH2SO4 is the same, the σDC at 95 °C under non-humidification was plotted against wH2SO4 (Figure 8a). The σDC values are summarized in Table 2. The conductivity of both S–P–S(s)/H2SO4 and S–P–S(l)/H2SO4 increased as wH2SO4 increased, while the conductivity of S–P–S(l)/H2SO4 was higher than or comparable to that of S–P–S(s)/H2SO4 with the same wH2SO4. Taking acid–base complexation between H2SO4 and the pyridyl group of S–P–S in the membranes into account, the molar ratio of acid to base, i.e., ADL, should directly affect the σDC under non-humidification because the number of protons released from free H2SO4 is strongly dependent on acid–base complexation, which consumes the free protons.

Figure 8.

Figure 8

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Conductivity at 95 °C under non-humidification for a series of S–P–S/H2SO4 membranes against (a) wH2SO4 and (b) ADL.

The relationship between the σDC under non-humidification and ADL of S–P–S/H2SO4 is exhibited in the plot in Figure 8b. Surprisingly, data points of σDC for both S–P–S(s)/H2SO4 and S–P–S(l)/H2SO4 membranes under non-humidification followed almost the same curve. In other words, the conductivity dependence for ADL is almost the same for S–P–S(s)/H2SO4 and S–P–S(l)/H2SO4, and the σDC of the S–P–S/H2SO4 membranes was greatly influenced by the acid/base stoichiometry associated with acid–base complex formation rather than the nanophase-separated structure adopted in the membranes. Although the experimental data was not measured for the region of ADL <1 in this study, the σDC at ADL <1 is expected to be very low, below 10–4 S cm–1 according to our previous studies.42,43 In the region of ADL >1, the membranes exhibit the σDC of 10–4 S cm–1 or higher, and the larger the ADL, the higher the σDC. For example, the σDC was 0.039 S cm–1 at ADL = 2.0, but when ADL >2, the degree of increase in conductivity decreased.

The very low σDC in the region of ADL <1 may reflect the few free protons directly contributing to conductivity, because almost all free H2SO4 in the membranes contributed to the formation of rigid ionic acid–base complexes consisting of a hydrogen sulfate anion and the protic pyridinium cation (Figure 9a). When the ADL exceeds unity, the excess H2SO4 not used for acid–base complex formation is mixed with rigid acid–base complexes as a plasticizer, and free protons promoting proton transport are generated by ionization of free H2SO4, attaining proton conduciveness of the membranes (Figure 9b). As the ADL became larger, the ionic interactions that arise in the acid–base complexes are further weakened and the acid–base complexes become softer due to the infiltration of more excess H2SO4 as a plasticizer (Figure 9c). Moreover, the absolute number of free protons generated by ionization of excess H2SO4 also increases, inducing the higher conductivity of the membranes. Note that the degree of the increase in the fraction of free protons declines in the region of ADL > 2, resulting in the gradual increase in σDC. In short, since the σDC dependence on ADL is almost the same for S–P–S(l)/H2SO4 and S–P–S(s)/H2SO4, S–P–S(l)/H2SO4, which has a smaller fraction of P, exhibits comparable or higher σDC than that of S-P-S(s)/H2SO4, even when wH2SO4 is the same for both membranes.

Figure 9.

Figure 9

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Schematic illustration of H2SO4 and pyridyl groups in S–P–S/H2SO4 at the molecular level for (a) ADL <1, (b) 1< ADL <2, and (c) ADL >2.

4. Conclusions

We investigated the effects of the nanophase-separated structure on the mechanical properties and proton conductivity of acid-infiltrated block polymer electrolyte membranes adopting a spherical or lamellar nanophase-separated structure by infiltrating sulfuric acid into S–P–S triblock copolymers. SAXS measurements revealed that no morphology transition occurred even after the infiltration of H2SO4 into neat S–P–S films. S–P–S(l)/H2SO4 membranes with a continuous hard S phase generally exhibited higher tensile strength than S–P–S(s)/H2SO4 membranes forming a spherical structure, even if the same amount of H2SO4 was infiltrated into each type of neat S–P–S film. Meanwhile, the conductivities of S–P–S(l)/H2SO4 membranes under non-humidification were higher or comparable to those of S–P–S(s)/H2SO4 membranes with the same wH2SO4. This outcome reflects the strong dependence of the σDC of the S–P–S/H2SO4 membranes on ADL, i.e., the stoichiometric ratio of acid to base, rather than on the nanophase-separated structure adopted in the membranes. In other words, there are more free molecules of H2SO4 that are not consumed for acid–base complexation, which can then release more free protons in S–P–S(l)/H2SO4 compared to S–P–S(s)/H2SO4 with the same weight fraction of H2SO4. This feature originates from the molecular characteristics of the two types of S–P–S, i.e., the difference in the fraction of the basic P block in S–P–S. Although the infiltrated acid is easily dissolved out in water from the S–P–S/H2SO4 membranes if the membranes are put into water, the S–P–S/H2SO4 membranes exhibit good conductivities even under non-humidification unlike the Nafion membrane; therefore, the S–P–S/H2SO4 membranes still have a high potential for application into PEFCs that can generate electricity even under non-humidification, if dissolving out of acid into water can be suppressed. The findings obtained in this study will help to design high-performance PEMs for development of next-generation fuel cells.

Acknowledgments

The authors thank Dr. Masaki Ando, Dr. Naoki Nakamura, and Mr. Seiji Sano at Toyota Motor Corporation for their kind advice on preparation and evaluation of proton-conductive polymer electrolyte membranes. This work was supported through a project (JPN P20003) subsidized by the New Energy and Industrial Technology Development Organization (NEDO) and KAKENHI grant numbers (21 K05197 (A.N.)) from JSPS, Japan. This research was technically supported by NEDO FC-Platform. The synchrotron experiment was performed at BL40B2 at SPring-8 under approval of the Japan Synchrotron Radiation Research Institute (JASRI) with proposal no. 2021A2054.

Supporting Information Available

The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/acsomega.2c06514.

  • GPC chromatograms of S–P–S(s) and S–P–S(l); 1H NMR spectra of S–P–S(s), S–P–S(l), and precursor S; DSC thermograms for a series of S–P–S(s)/H2SO4 and S–P–S(l)/H2SO4 membranes; the impedance spectroscopy setup; the q1 and D of the SAXS profile; Nyquist plots of S–P–S/H2SO4 membranes; and VFT fitting (Figures S1–S6 and Tables S1–S2) (PDF)

Author Present Address

# Fuel Cell Cutting-Edge Research Center Technology Research Association, 2-3-26 Aomi, Koto-ku, Tokyo 135-0064, Japan

Author Contributions

The manuscript was written through contributions of all authors. A.N. conceived the original idea and designed the study with T.K. T.K. carried out the tensile tests and DSC measurements. H.T. and Y.O. performed polymer synthesis, sample preparation, and conductivity measurements. T.O. and A.T. conducted TEM observation. H.Iw., A.M., and H.Im. carried out SAXS measurements. A.N. and T.K. analyzed the data and co-wrote the paper with input from all authors. All authors have given approval to the final version of the manuscript.

The authors declare no competing financial interest.

Supplementary Material

ao2c06514_si_001.pdf (1.2MB, pdf)

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