Mechanically Robust Anion Exchange Membranes via Long

Mar 9, 2017 - †Department of Materials Science and Engineering and ‡Department of Chemistry, The Pennsylvania State University, University Park, P...
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Mechanically Robust Anion Exchange Membranes via Long Hydrophilic Cross-Linkers Liang Zhu,† Tawanda J. Zimudzi,‡ Ying Wang,§ Xuedi Yu,† Jing Pan,† Juanjuan Han,† Douglas I. Kushner,† Lin Zhuang,§ and Michael A. Hickner*,† †

Department of Materials Science and Engineering and ‡Department of Chemistry, The Pennsylvania State University, University Park, Pennsylvania 16802, United States § College of Chemistry and Molecular Sciences, Hubei Key Lab of Electrochemical Power Sources, Wuhan University, Wuhan 430072, China ABSTRACT: Mechanically tough, cross-linked anion exchange membranes (AEM) based on poly(2,6-dimethylphenylene oxide) (PPO) were achieved by introducing a hydrophilic and flexible Jeffamine (O,O′-bis(2-aminopropyl)polypropylene glycol-block-poly(ethylene glycol)-block-polypropylene glycol 500) cross-linker into the cationic macromolecular network. The Jeffamine cross-linked AEMs demonstrated outstanding strength and flexibility and were mechanically tougher than AEM samples based on benzyltrimethylammonium (BTMA) functionalized PPO alone. The hydrated BTMA40 membrane showed 52% elongation at break, while the Jeffamine-based J10PPO sample had a 167% elongation at break. In addition, the hydroxide (OH−) conductivity of the J10PPO sample was 52 mS/cm at 80 °C with a swelling ratio of 61%, while BTMA60 suffered severe swelling above 60 °C. The alkaline stabilities of the AEMs with different degrees of Jeffamine cross-linking were evaluated in 1 M NaOH at 80 °C for 500 h. During the 500 h degradation test, J10PPO exhibited the greatest cation stability. The OH− conductivity of this membrane decreased by 30% over 500 h. In contrast, the OH− conductivity of BTMA40 decreased to 9.6 mS/cm at 20 °C, which is 60% lower than the value measured for the sample before the stability test. Based on the highperformance Jeffamine cross-linked AEM, a Pt-catalyzed fuel cell with a peak power density of 314 mW/cm2 was demonstrated at 60 °C under 100% related humidity.

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INTRODUCTION Large scale, environmentally friendly energy harvesting, storage, and conversion technology will be required in the 21st century to secure a sustainable future for humankind.1 Inexpensive fuel cells are one of the promising emerging technologies for large scale energy conversion and storage. In a fuel cell, electricity is generated from chemical fuels by means of an electrochemical reaction that is promoted by an overall negative Gibb’s free energy of reaction between the fuel and the oxidant. Unlike batteries and internal combustion engines, fuel cells directly convert the chemical energy of gaseous fuels into electrical energy at relatively low temperatures and with no intermediate steps between fuel and electricity. The capacity of the fuel cell can be adjusted by the size of the chemical fuel storage vessel and the size of the chemical fuel storage and the device output can be ramped across a wide power range. Proton exchange membrane fuel cells (PEMFCs) possess many advantages such as compact device architectures, fast startup, and high discharge power densities, rendering them promising power conversion devices for vehicles and portable electronic devices.2−4 The most frequently used proton exchange membrane (PEM) is Nafion, a perfluorinated sulfonic acid (PFSA) resin with a perfluorinated backbone and pendant perfluorosulfonic acid side groups, produced by DuPont.5,6 Nafion and its derivatives © XXXX American Chemical Society

exhibit high proton conductivity (above 0.1 S/cm at operating conditions), excellent strength (25 MPa tensile strength) and flexibility (180% elongation at break), and outstanding chemical stability, ensuring high operational performance and long-term durability of PEMFCs.7−9 However, the widespread commercialization of PEMFC devices faces roadblocks. First, the potential environmental impact of synthesizing and processing PFSAs is of great concern.10 Second, PEMFCs require expensive precious metals, such as platinum (Pt), which are stable and efficient electrocatalysts in the acidic environment but are rare and difficult to economically deploy in large scale for the automotive industry.10 In contrast, AEM fuel cells (AEMFCs) which utilize AEMs as a solid polymer ion-conducting membrane are a promising alternative to PEMFCs due to their potential to employ nonprecious electrocatalysts, thus greatly lowering the cost of fuel cell technology.10−12 Compared to PEMFCs, AEMFCs have potentially lower overpotential electrode reactions13,14 and more options for cathode catalysts, such as nonprecious transition metals and their oxides, resulting in performance Received: June 28, 2016 Revised: August 8, 2016

A

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improvements and reduction in device costs.15−18 On the other hand, AEMs generally exhibit inferior properties compared to PEMs including lower OH− mobility and a less developed phase-segregated morphology and disorganized ionic domain structure due to their nonfluorinated backbones.7 In order to enhance ionic conductivity and maintain an acceptable swelling ratio in the hydrated materials, researchers have employed various strategies for cross-linking ionic polymers19−30 and controlling the phase-segregated morphology of cationic polymers to achieve high-performance AEMs.10,31−35 Although rigid aromatic polymer backbones ensure good tensile strength (higher than 20 MPa), the poor flexibility (lower than 25% elongation at break) of aromatic-based membranes is unsatisfactory for fuel cell applications.35−42 In low relative humidity environments, poor flexibility renders AEMs brittle, causing problems during MEA fabrication in the dry or wet state.7,15 Also, long-term fuel cell performance under dynamic conditions is closely related to the flexibility and toughness of AEMs since the fuel cell experiences humidity and temperature fluctuations which cause mechanical stresses in the membrane.7,15 One promising solution to resolve the brittleness issue present in many types of aromatic ion-conducting membranes is to introduce flexible structures into the rigid macromolecular architecture. For example, Yang et al.43 reported that grafting poly(ethylene glycol) (PEG, Mn = 350, 750, and 1000 g/mol) onto a poly(styrene-ethylene-co-butylene-styrene) (SEBS)

main chain enhanced the elongation at break of the material and enlarged the size of the ion-conducting channels, resulting in improvement of flexibility and ion conductivity of the resulting AEMs. However, the tensile strength of the PEGSEBS AEM materials was lower than 10 MPa, which is unsatisfactory for fuel cell applications.7,15 In our previous report,7,44 we synthesized tough and chemically stable semiinterpenetrating network anion exchange membranes (SIPN AEMs) by introducing a hydrophilic, thiol−ene cross-linked, flexible poly(ethylene glycol)-co-poly(allyl glycidyl ether) (xPEG−PAGE) component into the rigid quaternized poly(2,6-dimethylphenylene oxide) (BTMA PPO) matrix. This network architecture demonstrated sufficient mechanical strength and much better flexibility than conventional BTMA PPO samples. To circumvent brittleness in ion-containing membranes, we were inspired by the successful example of elastic gels and networks in which long, flexible cross-links are often employed to improve the mechanical strength and flexibility of the materials.45,46 Herein, we employ a unique approach to toughen cationic PPO-based AEMs by creating a cross-linked network with commercially available Jeffamine (O,O′-bis(2aminopropyl)polypropylene glycol-block-poly(ethylene glycol)block-polypropylene glycol 500) cross-linkers to impart flexibility to the resulting materials. Introducing the hydrophilic cross-linked network into the AEM greatly enhanced the toughness and chemical stability of the materials. In addition, B

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Characterization. 1H nuclear magnetic resonance (NMR) spectra were recorded at 300 MHz on a Bruker AV 300 spectrometer (Billerica, MA), and chemical shifts were listed in parts per million (ppm) downfield from tetramethylsilane (TMS). Tensile measurements were obtained for samples equilibrated in liquid water and ambient air at room temperature. The hydrated and dry membranes were cut into a dumbbell shape (12 mm × 3 mm in the test area), and tensile measurements were performed using an Instron-5866 (Norwood, MA) instrument at a crosshead speed of 5 mm/min at room temperature (25 °C). Ionic conductivity (σ) was measured using impedance spectroscopy on a Solartron 1260 A impedance/gain-phase analyzer (Farnborough, Hampshire, UK) with a two-point, in-plane geometry at frequencies ranging from 100 kHz to 100 Hz.48 The membrane resistance was obtained from the real value of the impedance where the imaginary response was zero. The ionic conductivity, σ in mS/cm, of each sample was calculated from σ = L/RA, where L is the distance between reference electrodes, R is the resistance of the sample, and A is the cross-sectional area of the sample. Bicarbonate conductivities were measured by exchanging the bromide form membranes in 1 M sodium bicarbonate for 24 h followed by rinsing to remove excess salt. Chloride conductivities were measured by exchanging the bromide form membranes in 1 M sodium chloride at room temperature for 24 h followed by extensive rinsing to remove excess salt. Hydroxide conductivities were measured by exchanging the bromide form membranes in 1 M sodium hydroxide for 24 h followed by rinsing to remove excess salt with degassed and deionized water. The membranes were subsequently placed into conductivity cells and immersed in deionized water that was degassed and blanketed with flowing argon gas. Water uptake was measured after drying the membrane in the desired counterion form at 60 °C under vacuum for 24 h. The dried membrane was immersed in water and periodically weighed on an analytical balance until a constant mass was obtained, giving the massbased water uptake. Water uptake was calculated by WU = (mhyd − m0)/m0, where mhyd is the hydrated sample mass and m0 is the dry sample mass. The hydration number (λ), or the number of water molecules per ionic group, was calculated from

the Jeffamine cross-linked membranes maintained a good balance between ion conductivity and mechanical properties and could be fabricated in large scale. In the previous work,7,44 we broke new ground in the design and optimization of networked AEMs. While that original work informed the current report, the PEG−PAGE network had to be synthesized as a custom polymer, which can hold back large-scale deployment of these materials. In the present work, we employed commercially available Jeffamine as an elastic crosslinker. This strategy enables kilogram-scale synthesis of AEMs and is an important step in the deployment of these materials as there is currently a dearth of low resistance, commercially available AEMs that are suitable for electrochemical applications. In addition, for SIPN AEMs, the introduction of thiol groups may have a negative impact on the fuel cell performance. In the SIPN work, fuel cell performance was not reported.7,44

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EXPERIMENTAL SECTION

Materials. Poly(2,6-dimethyl-1,4-phenylene oxide) was purchased from Sigma-Aldrich and dried under vacuum at room temperature overnight. N,N-Dimethylhexylamine, O,O′-bis(2-aminopropyl)polypropylene glycol-block-poly(ethylene glycol)-block-polypropylene glycol 500 (Jeffamine ED-900, CAS Number 65605-36-9), Nbromosuccinimide, and 2,2′-azobis(2-methylpropionitrile) were obtained from Sigma-Aldrich and used as received. The brominated PPO polymers with a degree of functionalization (DF) of 60 mol % of the repeat units on the backbone with one benzyl bromide group were synthesized according to the reported literature.47 Preparation of BTMA40 Membrane. To a round-bottom flask was added poly(2,6-dimethyl-1,4-phenylene oxide) (24 g, 200 mmol) and chlorobenzene (200 mL) with stirring until the polymer was completely dissolved. To the flask was added N-bromosuccinimide (22.5 g, 126.6 mmol) and 2,2′-azobis(isobutyronitrile) (1.20 g, 7.14 mmol). The flask was fitted with a reflux condenser and heated to 125 °C for 3 h. The product of the mixture was cooled to 25 °C and slowly poured into a flask of stirring methanol (∼1 L) at room temperature to precipitate the product residue. The precipitate was filtered and washed with methanol (2 × 500 mL). The product was collected and dried under vacuum (0.1 Torr) for 12 h to afford a brown powder that was benzyl brominated PPO (Br-PPO). To a round-bottom flask was added the aforementioned brown powder Br-PPO precursor (2.0 g), N-methylpyrrolidone (NMP) (20 mL), and aqueous trimethylamine (∼45 wt %, 3.2 mL). The mixture was stirred at 25 °C for 24 h. The obtained solution was poured into 100 mL of toluene or hexane to precipitate the polymer. The product was filtered and further washed with hexane and toluene several times. A dark-yellow powder was obtained and dried at 50 °C under vacuum overnight. The BTMA40 polymer (2 g) in bromide form was dissolved in NMP (20 mL) to yield a 10 wt % solution. The solution was then cast onto a leveled Teflon mold and dried at 82 °C under ambient pressure for 24 h followed by vacuum drying for another 24 h at 80 °C to give an ∼60 μm thick, transparent, tough polymer film. Fabrication of Membranes. Brominated PPO with a DF of 60 (2.5 g, 14.88 mmol) was dissolved in 30 mL of NMP. Then, N,Ndimethylhexylamine (0.97 g, 7.47 mmol) was added slowly. The mixture was stirred at room temperature for 48 h. Jeffamine (0.54 g, 0.89 mmol) was added to the resulting mixture. Subsequently, the solution was then cast onto a leveled Teflon mold and and dried at 82 °C under ambient pressure for 24 h followed by vacuum drying for 24 h at 80 °C to give an ∼100 μm thick, transparent, tough film. The resulting membrane was J10PPO (10% Jeffamine cross-linked, the J stands for Jeffamine cross-linker; as shown in Scheme 1, the 60 refers to the degree of aromatic bromination where 60 mol % of the PPO repeat units are brominated; the x indicates that the x mol % of the brominated PPO repeat units react with N,N-dimethylhexylamine, and the 60 − x is equal to 10) in the bromide form.

⎛ mhyd − m0 ⎞⎛ 1000 ⎞ ⎟⎟ λ=⎜ ⎟⎜⎜ m0 ⎝ ⎠⎝ M H2O· IEC ⎠

(1)

where mhyd is the hydrated membrane mass, m0 is the mass of the dry membrane, MH2O is the mole mass of water (18 g mol−1), and IEC is the ion exchange capacity with units of milliequivalents of ions per gram of dry polymer. The swelling degree (SW) was characterized by the linear expansion ratio of the material, which was determined using the difference between the wet and dry dimensions of a membrane sample (3 cm in length and 1 cm in width). The calculation was based on the equation SW (%) =

X wet − Xdry Xdry

× 100% (2)

where Xwet and Xdry are the lengths of the wet and dry membranes, respectively. To calculate the titrated gravimetric IEC values, membranes in the OH− form were immersed in 50 mL of 0.01 M HCl standard solution for 24 h. Then, the solutions were titrated with a standardized NaOH (0.01 M) solution to pH = 7. Subsequently, the samples were washed and immersed in deionized water for 24 h to remove the residual HCl and then dried under vacuum at 50 °C overnight and weighed to calculate the dry masses in the Cl− form. The IEC of the membranes was calculated via eq 3:

IEC =

n i(H+) − nf(H+) mdry(Cl)

(3)

where mdry(Cl) is the mass of dry membranes, ni(H+) is the initial amount of H+ in the HCl solution, and nf(H+) is the final amount of H+ in the HCl solution. C

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Table 1. Ion Exchange Capacity, Water Uptake, Ionic Conductivity, and Mechanical Properties of Jeffamine Cross-Linked AEMs sample

Jeffamine (mol %)

IECa (mmol/g)

IECb (mmol/g)

J5PPO J10PPO J15PPO J20PPO BTMA20 BTMA30 BTMA40 BTMA60

5 10 15 20 0 0 0 0

2.51 2.22 1.94 1.67 1.61 2.13 2.66 3.50

2.35 2.08 1.83 1.58 1.47 1.92 2.50 3.21

WU (wt %) (OH−)c 189 228 165 121 21 59 116 450

± ± ± ± ± ± ± ±

15 20 11 9 2 5 8 19

σ (mS/cm) (OH−)c

λ (OH−)

gel fractiond (%)

± ± ± ± ± ± ± ±

44 56 43 33 8 17 27 78

79 81 84 85 0 0 0 0

29 24 23 23 6 14 24 31

2 1 1 1 1 2 2 2

Calculated from the polymer composition and the degree of bromination. bTitrated values (standard error: ±10%). cMeasured at room temperature in liquid water. dStandard error: ±3−5%.

a

After cross-linking, the membranes remained insoluble in all common solvents, and gel fraction was used to obtain information on the completeness of the cross-linking reaction. The gel fraction of the cross-linked membranes was calculated from the ratio of the weight of the polymer after extraction with NMP at 80 °C for 1 day and the initial weight of the sample before extraction. FTIR spectra were obtained using a Bruker Vertex 70 FTIR spectrometer (Billerica, MA) equipped with a liquid nitrogen cooled MCT detector. Transmission spectra were collected using 7 mm KBr pellets prepared by mixing 2.5 mg of polymer with 60 mg of anhydrous KBr (Sigma-Aldrich). The spectra were signal averaged over 100 scans at 4 cm−1 resolution with a 0.5 mm aperture size and a nitrogen purge at ambient temperature. All spectra were processed using Bruker OPUS 6.5 software. Cast membrane samples were cross-sectioned at −120 °C using a LeicaUltracut UC6 ultramicrotome with an EMFC6 cryo attachment. The sections were supported on carbon/Formvar-coated grids. Transmission electron images were collected on a JEOL JEM 1200 EXII microscope equipped with a tungsten emitter operating at 80 kV. Images were captured on unstained samples recorded on a CCD camera using TCL software. Membrane−Electrode Assembly Fabrication and Fuel Cell Testing. Pt/C (60%, Johnson Matthey Co.) was mixed with BTMA40 ionomer solution and sprayed onto both sides of J10PPO (IEC = 2.08 mmol/g) or BTMA30 (IEC = 1.92 mmol/g) membrane (100 ± 5 μm in thickness), respectively, to prepare the catalyst-coated membranes (CCMs) with the same ionic material in the membrane and in the electrode. The weight percentage of ionomer in both the anode and the cathode was calculated to be 20 wt %. The Pt loading in both anode and cathode was 0.4 mg/cm2, and the area of the electrodes was 4 cm2. The resulting CCMs were pressed between two pieces of carbon cloth (CeTech W1S1005) to form the membrane electrode assembly (MEA) which was then tested in a fuel cell device. After the H2 and O2 gases were provided to the anode and cathode, the humidity and temperature of the fuel cell system were gradually increased to the desired set point values. After the voltage data for the galvanostatic IV curve was stable over time, we continued to increase the current to the next set point. The equilibration time between data points for the individual IV curve points was approximately 5 min. The H2−O2 fuel cell performance of the MEAs was measured under a galvanic mode using fully humidified H2 and O2 gases flowing at a rate of 250 mL/min and at temperature of 60 °C with 0.1 MPa backpressure (850e Multi Range, Scribner Associates Inc.).

Menshutkin reaction of the olefinic tertiary amine (N,Ndimethylhexylamine shown above) with the bromobenzyl group of Br-PPO (Scheme 1). Subsequently, the cross-linked samples were fabricated by the reaction of the remaining bromobenzyl groups of the Br-PPO with the primary amine groups of the Jeffamine cross-linker during membrane casting. The initial DB of Br-PPO was maintained at 60 to obtain samples with calculated IECs ranging from 1.67 to 2.51 mmol/ g (Table 1). Figure 1 shows a typical 1H NMR spectrum for the resultant BTMA material; peaks were observed at chemical shifts of 4.3

Figure 1. 1H NMR of BTMA30 in DMSO-d6/D2O (10:1 w:w).

and 3.1 ppm corresponding to the H atoms in the benzyl group and the quaternary ammonium group, respectively. In this work, by setting the DB of Br-PPO at 20, 30, 40, and 60%, BTMA samples with calculated IECs ranging from 1.61 to 3.50 mmol/g (Table 1) were obtained. FTIR was used to confirm the reaction of the bromobenzyl group on Br-PPO with the −NH2 of Jeffamine to form the cross-linked network. Compared to the PPO, there were distinct v(C−Br) peaks at 634 and 987 cm−1 for the brominated PPO, which appeared after bromination, as shown in Figure 2.49,50 Both of the peaks assigned to v(C− Br) disappeared after cross-linking. In the FTIR spectrum of complete cross-linked membranes by the Jeffamine cross-linker, the new peaks at 1106 cm−1 associated with v(C−O) verified the presence of the Jeffamine cross-linker in the AEMs.51

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RESULTS AND DISCUSSION The synthetic route for the Jeffamine cross-linked AEMs is shown in Scheme 1. Br-PPO possessing a degree of bromination (DB) of 60 mol % at the benzyl position was synthesized using N-bromosuccinimide and 2,2′-azobis(2methylpropionitrile) in a refluxing chlorobenzene solution for 3 h according to our previous reports.35 Then, a series of crosslinkable comb-shaped ionomers was achieved by the D

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exhibited high tensile strength (24 MPa) but possessed low flexibility (52% elongation at break). Compared to the BTMA40 membrane, the J10PPO demonstrated a significantly higher elongation at break of 167%. The area under the stress− strain curves for these samples was integrated in order to determine the material toughness. The toughness of the BTMA40 and J10PPO samples was 1.05 × 104 and 1.5 × 104 kJ/m3, respectively, representing a 150% increase in toughness for the J10PPO sample. It is clear that introducing the hydrophilic, flexible cross-linker into the cationic network enhanced the mechanical performance of the J10PPO material. The mechanical performance of the Jeffamine cross-linked AEMs is better than the previous reported values for crosslinked AEMs.19−30 For example, the cross-linked PE-AEMs synthesized by Robertson et al.26 exhibited tensile strength at break of less than 28 MPa in the dry state. Lin et al.28 reported cross-linked [PVMIm][OH]-DVB with an IEC = 0.75 mmol/g exhibiting elongation at break of 103% at room temperature. The cross-linked, poly(arylene ether sulfone)-based AEMs reported by Zhuo et al.21 exhibited elongation at break of less than 9% in the dry state. He et al.36 reinforced benzyltrimethylammonium polysulfone-based AEMs with a cross-linked poly(styrene-co-divinylbenzene) network. The elongation at break of the cross-linked AEM materials with this hydrophobic network was less than 10%. Pan et al.7 synthesized SIPN AEMs by introducing a hydrophilic poly(ethylene glycol)-co-poly(allyl glycidyl ether) (xPEG−PAGE) component into the rigid BTMA PPO matrix. The resulting cross-linked AEMs exhibited an elongation at break as high as 100% in the wet state. The comparison of tensile strength between the hydrated Jeffamine cross-linked and conventional BTMA40 samples is shown in Figure 4a. The tensile strength decreased by increasing the content of the Jeffamine cross-linker due to the flexible and hydrophilic properties of the cross-linker. In addition, the increase in the Jeffamine cross-linker content reduced the concentration of rigid PPO backbones, which decreased the strength of the AEMs. Compared with J5PPO, the J10PPO showed a higher tensile strength due to the tradeoff between decreased content of the rigid backbone and IEC. The comparison of elongation at break between the hydrated Jeffamine cross-linked and conventional BTMA40 AEMs is shown in Figure 4b. When the degree of cross-linking increased from 0 to 10%, the elongation at break increased by boosting the Jeffamine content. For example, J10PPO exhibited a 177%

Figure 2. FTIR spectra of PPO, brominated PPO, Jeffamine and 100% Jeffamine cross-linked membranes.

Considering these results, a strong covalent cross-linking network was formed by bromine and the secondary amine reaction which was further confirmed by the gel fraction (>80%) in NMP, Table 1. Mechanical Properties. The comparison of mechanical properties between J10PPO and BTMA40 hydrated AEMs is shown in Figure 3. As a conventional rigid AEM, BTMA40

Figure 3. Stress−strain curves for J10PPO and BTMA40 hydrated AEMs.

Figure 4. Mechanical properties for Jeffamine cross-linked and BTMA40 hydrated AEMs: (a) tensile stress at break; (b) maximum elongation at break. E

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Figure 5. Hydroxide conductivity of Jeffamine cross-linked and BTMA AEMs in liquid water at 20 °C as a function of (a) IEC and (b) hydration number (λ).

elongation at break, while BTMA 40 showed only 52% elongation at break. However, the elongation at break decreased as the degree of cross-linking increased from 10 to 20% due to the trade-off between the content of Jeffamine and degree of cross-linking. Increasing the content of Jeffamine cross-linker in the sample rendered materials with similarly high gel fractions (Table 1). High ionic conductivity is one of the most important properties for the application of AEMs in devices. To study the ion transport behavior of the Jeffamine cross-linked AEMs, OH− conductivities were measured for AEMs with different IEC values. As expected, for the same type of AEM, the ionic conductivity decreased with decreasing IEC, as shown in Figure 5a. The J5PPO membrane displayed an OH− conductivity of 29 mS/cm at room temperature with an IEC of 2.35 mmol/g compared to 24 mS/cm for BTMA40 (IEC = 2.50 mmol/g) under the same experimental conditions. For a better comparison among the different IEC membranes, the hydroxide conductivity was plotted as a function of λ in Figure 5b. For the BTMA membranes, the approximate trend was that the hydroxide conductivity increased with λ due to the increasing number of water molecules per ionic site which facilitated the hydroxide transport in the membranes. Jeffamine cross-linked membranes with similar IEC values showed relatively higher hydroxide conductivity compared to the BTMAx membranes at much higher hydration number (λ). For example, J5PPO had a hydroxide conductivity of 29 mS/cm at λ = 44, while the BTMA40 sample showed a hydroxide conductivity of 24 mS/ cm at λ = 27. This result is attributed to the introduction of the hydrophilic cross-linker in the Jeffamine-based sample that helps to boost the conductivity of AEMs without over hydration and large swelling degrees. One of the barriers to develop high-performance AEM materials is the trade-off between ion conductivity and swelling ratio. As shown in Figure 6, the BTMA30 and BTMA40 samples exhibited low swelling ratios of 7.5% and 36% at 80 °C, respectively. However, the low swelling ratio was usually associated with low ion conductivity. The BTMA30 sample with an IEC of 2.10 mmol/g showed an OH− conductivity of 36 mS/cm at 80 °C. The OH− conductivity of the BTMA membranes was enhanced by increasing the IEC of the samples. As depicted in Figure 6, BTMA60 with an IEC of 3.22 mmol/g displayed an OH− conductivity of 43 mS/cm at 60 °C, which was much higher than that of BTMA30 sample (29 mS/cm at

Figure 6. Hydroxide conductivity of Jeffamine cross-linked and BTMA AEMs in liquid water as a function of swelling ratio and temperature.

60 °C). However, enhancing ion conductivity through increasing the IEC resulted in significant reduction in the mechanical properties of the material. As shown Table 2, Table 2. Mechanical Properties of OH− Type Jeffamine Cross-Linked AEMs and BTMA Samples in the Dry and Wet Statesa tensile strength at break (MPa) sample J5PPO J10PPO J15PPO J20PPO BTMA40 BTMA60 a

dry 35 41 38 37 47 21

± ± ± ± ± ±

wet 3 5 3 3 4 2

10 15 8 7 24 16

± ± ± ± ± ±

elongation at break (%) dry

1 2 1 1 3 2

21 26 22 18 17 15

± ± ± ± ± ±

wet 2 2 1 2 3 2

91 177 109 75 52 14

± ± ± ± ± ±

7 9 10 6 7 1

Properties were measured at room temperature.

BTMA60 exhibited very low tensile strength and elongation at break, which is not desirable for long-term fuel cell applications. In contrast, the Jeffamine cross-linked membranes maintained a good balance between ion conductivity and mechanical properties. Introducing the hydrophilic Jeffamine cross-linker as part of the network enhanced the toughness of the AEM materials and maintained a high hydration number to ensure sufficient ion conductivity at elevated temperature. As depicted in Figure 6, J10PPO demonstrated a hydroxide conductivity of F

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To evaluate the long-term alkaline stability of the AEMs, Jeffamine cross-linked AEMs and BTMA40 samples were immersed in argon-saturated 1 M NaOH solution at 80 °C for 500 h, with replacement of the 1 M NaOH every 3 days during the testing period. Figure 9 illustrates the changes in OH−

Figure 9. Chemical stability of the quaternary ammonium cation in Jeffamine cross-linked and BTMA40 samples in 1 M NaOH solution at 80 °C, OH− conductivity as a function of degradation time at 20 °C.

Figure 7. Hydroxide conductivity of Jeffamine cross-linked and BTMA membranes in an Arrhenius-type temperature plot showing activation energies in kJ/mol.

conductivity during chemical stability testing of the samples. During the 500 h test, J10PPO exhibited the greatest cation stability. The OH− conductivity of this membrane decreased by 30% over the evaluation period. In contrast, the OH − conductivity of BTMA40 decreased to 9.6 mS/cm at 20 °C, which 60% lower than the value measured for the sample before the stability test. The increased stability of the Jeffamine cross-linked AEM compared to the BTMA control sample was observed due to the cross-linking network and an increase in the microphase separation of hydrophilic and hydrophobic domains in this material, which is in good agreement with other reported results.53 While not completely mitigating the degradation of the ammonium-based polymers, the addition of hydrophilic moieties, like PEG, increases the degradation resistance of AEMs as shown in our previous work.7 Fuel Cell Performance. Selected AEMs, J10PPO (IEC = 2.08 mmol/g) and BTMA30 (IEC = 1.92 mmol/g), were evaluated in fuel cells to ascertain the suitability of the materials in operational devices. As is displayed in Figure 10, with similar membrane thickness (100 ± 5 μm) and Pt catalyst loading (0.4 mg Pt/cm2 on the anode and cathode), the assembled H2−O2 single cells with the same material in the electrode and the

temperature dependence promoted by the thermal activation of water motion.52 The apparent activation energy of ion conduction estimated from the slopes of the ln(σ) vs 1000/T curves was ∼10 kJ/mol and was similar among the measured membranes. The apparent activation energy of the tested AEM samples was somewhat lower than those of reported AEMs (10−23 kJ/mol), indicating that Jeffamine cross-linked membranes have similar aqueous OH− conduction attributes to other membrane samples.52 The addition of poly(ethylene glycol) (PEG) and other hydrophilic, nonionic components into the AEMs seems to have a positive effect of lowering the activation energy for conduction, which may be desirable in certain applications.7 Figure 8 shows transmission electron micrographs for the BTMA40 and J10PPO samples where the dark areas represent

Figure 8. TEM images of dry membranes in the bromide form: (a) BTMA40 and (b) J10PPO.

the ionic domains, while the bright areas represent the nonionic domains. As shown in Figure 8, for the J10PPO membrane, ion aggregation in the TEM image was obvious, and hydrophilic/ hydrophobic phase separation was observed tying the increased conductivity of these samples to their well-developed morphologies. In contrast, no obvious ion aggregation was observed in the TEM image for BTMA40, which is characteristic of BTMA-based AEMs and results in lowered conductivity.

Figure 10. Fuel cell performance of the AEM samples under H2/O2 conditions. G

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Macromolecules

Research Institute and the Penn State Institutes of Energy & the Environment.

membrane show that the type of AEM influenced the cell performance. Specifically, at a cell operating temperature of 60 °C and gas flow rates of 250 mL/min with a back-pressure of 0.1 MPa, the peak power densities for the J10PPO and BTMA30 AEM devices were 318 and 96 mW/cm2 under current densities of 610 and 200 mA/cm2, respectively. The introduction of a hydrophilic cross-linked network yielded up to a 300% increase in the peak power density from 96 mW/cm2 for BTMA30 to 318 mW/cm2 for J10PPO as a result of a more efficient ion transport architecture and enhanced toughness of the AEM materials.36

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CONCLUSIONS

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AUTHOR INFORMATION

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REFERENCES

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In summary, a unique approach was employed to toughen AEMs by cross-linking the functionalized PPO cationic network using commercially available Jeffamine as a cross-linking agent. Compared to the BTMA40 membrane, the Jeffamine crosslinked membrane demonstrated significantly higher elongation at break and toughness. The hydrated BTMA40 membrane showed a 52% elongation at break, while the 10% Jeffamine cross-linked membrane had a 167% elongation at break. Clearly, the introduction of a hydrophilic cross-linked network greatly enhanced the toughness of the AEMs. Compared to BTMA40, the Jeffamine-based J10PPO had a 150% increase in toughness. In addition, the Jeffamine cross-linked membranes maintained a good balance between ion conductivity and mechanical properties. Introducing a hydrophilic Jeffamine cross-linker into the cationic network enhanced the toughness of the AEM materials and maintained a high hydration number to ensure sufficient ion conductivity at elevated temperature. Compared to BTMA-based AEMs, Jeffamine cross-linked membranes exhibited higher hydroxide conductivity and better long-term alkaline stability compared to PPO-based AEMs, which is critical for alkaline fuel cell applications. Employing the high performance J10PPO membrane as a separator, an AEMFC using Pt catalysts was successfully demonstrated, which showed a peak power density of 318 mW/cm2 at a current density of 610 mA/cm2 at a device operation temperature of 60 °C. Overall, the Jeffamine cross-linked AEM materials exhibited better mechanical properties, especially in regards to elongation at break, compared to our previously reported materials.35 More importantly, the introduction of a hydrophilic cross-linked network yielded up to a 300% increase in the peak power density from around 100 mW/cm2 for C6D40 to 318 mW/cm2 for J10PPO as a result of a more efficient ion transport architecture, increased hydrophilicity, and enhanced toughness of the AEM materials.31

Corresponding Author

*E-mail [email protected] (M.A.H.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was funded by the NSF DMREF program under award number CHE-1534326. M.A.H. also acknowledges the Corning Foundation for support through the Pennsylvania State University, College of Earth and Mineral Sciences Corning Faculty Fellowship. Infrastructure support was provided by The Pennsylvania State University Materials H

DOI: 10.1021/acs.macromol.6b01381 Macromolecules XXXX, XXX, XXX−XXX

Article

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DOI: 10.1021/acs.macromol.6b01381 Macromolecules XXXX, XXX, XXX−XXX