Enhancement of hydroxide conduction by incorporation of metal

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Enhancement of hydroxide conduction by incorporation of metal-organic frameworks into a semi-interpenetrating network Nosaibe Anahidzade, Mohammad Dinari, Amir Abdolmaleki, Koorosh Firouz Tadavani, and Mohammad Zhiani Energy Fuels, Just Accepted Manuscript • DOI: 10.1021/acs.energyfuels.9b00650 • Publication Date (Web): 07 Jun 2019 Downloaded from http://pubs.acs.org on June 8, 2019

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Enhancement of hydroxide conduction by incorporation of metalorganic frameworks into a semi-interpenetrating network

Nosaibe Anahidzadea, Mohammad Dinaria, Amir Abdolmalekia,b,1, Koorosh Firouz Tadavania, Mohammad Zhiania

aDepartment

bDepartment

of Chemistry, Isfahan University of Technology, Isfahan 84156-83111, Iran

of Chemistry, College of Sciences, Shiraz University, Shiraz 71467-13565, Iran

1

Corresponding author: Tel.: +98 3133913249; Fax: +98 3133912350 E-mail address: [email protected], [email protected] (A. Abdolmaleki)

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ABSTRACT In order to overcome the trade-off barrier between conductivity and stability in anion-exchange membranes and also to investigate the effect of the metal-organic framework to resolve this limitation, two membranes (semi-IPN, and semi-IPN MOF) were prepared through using the solution casting method. The membranes were characterized by FT-IR, 1H-NMR, BET, scanning electron microscopy, and thermogravimetric analysis. Moreover, the effect of metal-organic frameworks was accurately investigated on the membrane features such as ion exchange capacity, water uptake, swelling ratio, hydroxide ion conductivity, thermal and mechanical properties, methanol crossover, single cell performance, and alkaline stability. The results indicated that the enhancement of ion exchange capacity (1.92𝑚𝑒𝑞.𝑔 ―1 vs 2.40𝑚𝑒𝑞.𝑔 ―1) caused by increased quaternary functional groups resulted in higher water uptake (48% vs 73%). In contrast, the crosslinked networks along with the robust metal-organic frameworks prevented the membranes from the excessive swelling ratio (a swelling ratio of 7%). Finally, the robust, porous, and hydrophilic Cr-MIL-101-NH2 frameworks via the construction of well-connected hydrophilic nanochannels significantly enhanced and facilitated hydroxide ion conductivities (0.07 S cm-1vs 0.01 Scm-1 at 30 °C). This strategy is a promising method to resolve the trade-off issue between hydroxide ion conductivity and swelling in anion exchange membranes.

Keywords: Post-modified metal-organic framework, Semi-interpenetrating networks, Ion nanochannels, Hydroxide conductivity, Anion-exchange membranes


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1.

Introduction

The alkaline anion exchange membrane fuel cells (AAEMFCs) are a kind of low-cost green energy-conversion devices, which are recognized as highly efficient power conversion technologies for stationary and mobile applications 1-2. The insufficient hydroxide ion conductivity of AEMs is one of the key barriers to the development of AEMs in contrast to proton-exchange membranes (such as Nafion). Ideal AEMs should maintain the following advantages: good mechanical stability, low degrees of swelling, high hydroxide conductivity, and sufficient chemical and thermal stabilities 3. High ionic conductivity generally needs to design AEMs with high ion exchange capacity (IEC). Increasing the membrane ion concentrations by the higher degree of quaternary functional groups may be useful in overcoming it, frequently causing excessive water uptake (WU) and swelling ratio (SR) 4. Thus, the trade-off effect between conductivity and SR still exists as a scientific challenge 5-8. The covalent cross-linking strategies have been frequently employed to sustain an adequate swelling of membranes at high IECs 9-10. However, some of the membranes with crosslinked linkages are slightly brittle in dry states and also hydroxide conductivity declines sharply because of the lack of well-connected ion conducting nanochannels and reduced WU 11-12. Currently, there has been an interest in semi-interpenetrating (semi-IPN) network anion exchange membranes containing both crosslinking networks and linear (flexible) polymer chains that interpenetrate one another. They have exhibited improved mechanical strength and flexibility, but their conductivity and WU need further improvements 13-14. On the other hand, hybrid ion conductive membranes, which merge the advantages of both polymeric and inorganic materials, for instance, SiO2 15-16, TiO2 17-18, ZrO2 19, graphene oxide 20-21, layered double hydroxide 22-24 and metal-organic


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frameworks (MOFs) 25-29, are effective to improve the ion conductivity and physicochemical stability of membranes. Chen et al. synthesized a new class of AEMs based on triple-cationic side-chain poly(2,6-dimethyl-1,4-phenylene oxide) (TC-QAPPO) membranes by spaying quaternary-ammonium-modified layered double hydroxide (QA-LDH) on the surface of the TCQAPPO membrane. The QA-LDH/TC-QAPPO composite membranes showed higher dimensional stability and hydroxide ion conductivity in comparison with TC-QAPPO membrane 30.

Li et al. synthesized a series of quaternized poly(arylene ether sulfone)/nanozirconia

composites. The modification with nano-ZrO2 improved the WU, dimension stability, and ion conductivity of composite membranes 19. Liang et al. prepared MOF-PVP hybrid membranes including protonated tertiary amines as proton carriers and high proton conductivity was observed at 25 °C and 53% RH 31. Despite the many applications of the cationic MOFs, the majority of MOFs are neutral. Mao et al. reported a post-synthetic modification method (anion stripping) based on the differential affinity between distinct metal ions with framework anionic species. In this way, Al+3 is used to strip F- anions away from Cr+3 sites in the MIL-100-Cr-F, and subsequently anion exchange with OH- leading to the network with mobile OH- anions. This material exhibit much improved ionic conductivity compared to the original unmodified MOFs 27.

Wu et al. reported the preparation of a novel anion exchange membrane (AEM) from porous

bromomethylated poly(2,6-dimethyl-1,4-phenylene oxide) (BPPO) entrapped cationic MOFs with polyvinyl alcohol (PVA) coating on the two sides. The entrapped cationic MOFs can work as the OH- conductive channels 26. Sadakiyo et al. reported the first example of the MOF-based hydroxide ion conductors. Alkylammonium hydroxides were introduced as ionic carriers in ZIF8 pores by host-gust interactions. ZIF-8 containing OH− ions showed four times higher ionic conductivity than blank ZIF-8 32. However, this system can only work in dry conditions because 4 ACS Paragon Plus Environment

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of the leaching in aqueous media. Thus, Gao et al. synthesized a polyvinyl(benzyl trimethylammonium hydroxide) (PVBTAH) caged in ZIF-8 that exhibited premier ion-exchange kinetics compared to traditional ion-exchange resins 33. One of the ZIF-8 problems was its microporous cages that made hard the impregnations of monomers into the MOF pours. Therefore, in their subsequent work, Gao et al. used Cr-MIL-101 because of its mesoporous cages and ease of impregnation of monomers. This group polymerized the sodium poly(4-styrene sulfonate) as a cationic polyelectrolyte into Cr-MIL-101 34-35. Li et al. reported the poly (ionic liquid) (PIL) as OH- ion carriers within the pores of Cr-MIL-101. The MOF membrane, prepared by a hot-press method, exhibited an efficient ion transport 25. Despite the benefits of functionalized MOFs-filled hybrid membranes, there is a dearth of research in this area. Herein, we selected Cr-MIL-101, with the empirical formula [Cr3 (O)(BDC)3 (F, OH) (H2O)2] (where BDC = benzene-1,4-dicarboxylate), because of its unique structure. It has two types of inner mesoporous cages with diameters of 29 and 34 Å and the pore size of the windows of up to 16 Å in diameter 36-37. In order to improve the WU and hydroxide ion conductivity of Cr-MIL-101, we introduced hydrophilic amino functionality (-NH2) via the post-synthetic modification (PSM) method in the terephthalate linkers in MOF. Then, a semiIPN MOF membrane was prepared by in situ polymerization assembly of N-vinylimidazole, Nvinyl-2-pyrrolidone and divinylbenzene (DVB) inside the cavities of Cr-MIL-101-NH2 in the presence of polyamide as the linear polymer. Well-connected nanochannels with a high density of ionic groups, high WU, and the formation of a continuous pathway for the facile hydroxide transfer owing to the existence of MOFs pores were leading to a sharp increase in ion conductivity of the semi-IPN MOF membrane.


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2.

Experimental

2.1. Materials Materials including 4-Fluoronitrobenzene (≥98%), cesium fluoride (CsF), calcium chloride (CaCl2, ≥99.9%), Pd/C 10%, N-methyl-2-pyrrolidone (NMP, ≥98%), dimethyl sulfoxide (DMSO), N,N-dimethyl formamide (DMF, ≥99.5%), pyridine (Py), triphenyl phosphite (TPP), panisidine (≥99%), N-vinyl-2-pyrrolidone (VP, ≥97%), N-vinylimidazole (VI, ≥99%) divinyl benzene (DVB, ≥80%), azobisisobutyronitrile (AIBN, 98%), iodomethane (CH3I, ≥99%), potassium hydroxide (KOH, ≥80%), chromium nitrate nonahydrate [Cr (NO3)3·9H2O, 99%], 1,4benzene dicarboxylic acid (H2BDC, 99%), concentrated sulfuric acid (98%), nitric acid (65 wt%), concentrated hydrochloric acid (37%), and SnCl2.2H2O, were purchased from Merck and Sigma-Aldrich Chemical Co. All the solvents were dried by conventional methods. Deionized (DI) water was used all over the experiments. 2.2. Preparation of Cr-MIL-101-NH2 Cr-MIL-101-NH2 was prepared from Cr-MIL-101 by the following procedure, as explained in details in the previous article 38. Terephthalic acid (0.164 g, 1.0 mmol), Cr(NO3)3·9H2O (0.4 g, 1.0 mmol), and HNO3 (1.0 mmol) were dissolved in water (5 mL) and then transferred to a Teflon-lined stainless steel autoclave. This acidic solution was heated at 220 °C for 8h. Afterward, the nitration of Cr-MIL-101 was performed by a post-synthetic modification method using a mixture of concentrated nitric acid and sulfuric acid. Eventually, Cr-MIL-101-NH2 was synthesized via SnCl2 reduction of Cr-MIL-101-NO2.


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2.3. 4, 4'-dinitro-4"-methoxytriphenylamine (dinitro intermediate (1)) Cesium fluoride (CsF) (1.52 g, 10 mmol) was added in 9 mL of dried dimethyl sulfoxide (DMSO) and stirred at room temperature 39. Afterward, p-anisidine (0.502 g, 4 mmol) and 4fluoronitrobenzene (1.168 g, 8.2 mmol) were added to the mixture, respectively, and heated with vigorous stirring at 120 °C for 24 h under nitrogen atmosphere. After cooling, the mixture was precipitated into 30 mL of methanol, collected by filtration and washed several times with a mixture of water: methanol (1:1 V/V) to obtain 1.307 g (89% in yield) of orange-red powder. 1H NMR (400 MHz, DMSO-d6) δ (ppm): 3.87 (s, 3H, OCH3), 7.14 (d, J= 8.90 Hz, 2H, Hc), 7.25 (d, J= 9.20 Hz, 4H, Hb), 7.31 (d, J= 8.90 Hz, 2H, Hd), and 8.24 (d, J= 9.20 Hz, 4H, Ha). 2.4. 4, 4'-Diamino -4"-methoxytriphenylamine (diamine monomer (2)) The dinitro compound (7.34 g, 20 mmol) and 10% Pd/C (0.13 g) were stirred with 100 mL of ethanol under a nitrogen atmosphere and heated to reflux 39. Thereafter, hydrazine monohydrate (7 mL) was added dropwise to the resulting mixture and the mixture was allowed to stir at reflux temperature for further 9 h. Eventually, the Pd/C was filtered and the resulting solution was cooled to precipitate light-green crystals. Subsequently, the product was collected by filtration and dried in a vacuum oven at 80 °C to give 4.88 g (80% in yield) of diamine. 1H NMR (400 MHz, DMSO-d6) δ (ppm): 3.72 (s, 3H, OCH3), 4.87 (s, 4H, NH2), 6.55 (d, J= 8.60 Hz, 4H, Ha), 6.72 (d, J= 8.60 Hz, 4H, Hb), and 6.74–6.78 (m, 4H, Hc + Hd). 2.5.

Synthesis of polyamide (3)

A mixture of 4,4'-diamino-4"-methoxytriphenylamine (0.382 g, 1.25 mmol), 4,4'-oxybis (benzoic acid) (0.322 g, 1.25 mmol), calcium chloride (0.15 g), TPP (1.25 mL), Py (0.6 mL), and 7 ACS Paragon Plus Environment

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NMP (1.25 mL) was stirred at 105 °C for 3 h 39. The polymer solution was slowly poured into 300 mL of methanol. Subsequently, the fiber-like precipitate was collected by filtration, washed completely with hot water and methanol, and dried in a vacuum oven at 100 °C to afford 0.538 g (99% in yield) of polyamide. 1H NMR (400 MHz, DMSO-d6) δ (ppm): 3.83 (s, 3H, OCH3), 6.878.17 (m, 20H, aromatics), and 10.33 (s, 2H, -NH-CO-). 2.6. Synthesis of semi-IPN polymer The synthesized polyamide (0.5 g) was dissolved in DMF (10 mL). Then, N-vinylimidazole: Nvinyl-2-pyrrolidone (1: 1, 0.5 g), divinylbenzene (DVB, 1 wt% of the total monomer), and AIBN (3 wt% of the total amount of monomers and DVB) were added to the solution and heated to 90 ºC under N2 atmosphere for 3 days. Afterward, CH3I (3 mL) was added to this suspension and stirred 2 days at room temperature and 3 days at 50 ºC. 2.7. Synthesis of semi-IPN polymer with encapsulating (VI-co-VP) within Cr-MIL-101-NH2 For encapsulating of N-vinylimidazole (VI) and N-vinyl-2-pyrrolidone (VP) monomers into CrMIL-101-NH2, activated Cr-MIL-101-NH2 (0.25 g) was suspended in dichloromethane (2 mL). N-vinylimidazole: N-vinyl-2-pyrrolidone (1:1, 0.5 g), DVB (1 wt% of the total monomers), and AIBN (3 wt% of the total amount of monomers and DVB) were added to this suspension, and to allow the penetration of monomers into Cr-MIL-101-NH2, the mixture was stirred at room temperature for 3 days. Due to the uniform distribution of the monomers as well as the further solubility of AIBN, dichloromethane was added to this suspension. Subsequently, dichloromethane was removed under vacuum at room temperature. Thereafter, the synthesized polyamide (0.5 g) was dissolved in DMF (10 mL), added to the N-vinylimidazole and N-vinyl-2-


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pyrrolidone filled Cr-MIL-101-NH2, and heated to 90 ºC under N2 atmosphere for 3 days. Afterward, CH3I (3 mL) was added to this suspension and stirred for 2 days at room temperature and 3 days at 50 ºC. 2.8. Membranes preparation For the film preparation, in each step, the resulting viscous solution was poured into a 9-cm glass petri dish and placed in an oven overnight to slowly removing the solvent at 90 ºC 40. Then, for further drying, the petri dish was stored in a vacuum oven at 120 ºC for an additional 8 h. Subsequently, the film was stripped off from the petri dish, washed with DI water to remove the unreacted monomers, and soaked in a 2 M KOH solution for the ion exchange process. The obtained membranes thickness was about 60-80 µm after drying in a vacuum oven for 24 h and was used for ion exchange conductivity tests. 2.9. Characterization The Fourier transform infrared (FTIR) spectra of the samples were recorded for analyzing the bending and vibrations of chemical bonding from 400-4000 cm-1 using a JASCO FT-IR 680 spectrometer (Japan). 1H NMR spectroscopy (Bruker Ultrashield 400, Germany) was performed in DMSO-d6 as the solvent. Thermogravimetric analysis (TGA, STA 503 win TA, BahrThermoanalyse GmbH, Hüllhorst, Germany) was performed to determine the thermal stability of the samples. The TGA measurements were conducted by heating the samples to 800 °C using the heating rate of 10 °C per minute under a nitrogen atmosphere. The prepared membranes morphology and microstructure were observed via a scanning electron microscope (SEM) with an accelerating voltage of 10 kV after sputter coating by gold. For cross-sectional SEM, the


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samples were obtained by fracturing in liquid nitrogen. Brunauer−Emmett−Teller (BET) surface areas and porosity were measured by (BELSORP mini II, Japan) analyzer at 77 K. For each run, ~0.05 g sample was evacuated in a vacuum oven at 120 °C for 4h. The WU of the membranes was determined by immersing the samples in a 1000 mL beaker filling with deionized water at 25 °C for a period of 24 h. Then, the samples were wiped with a filter paper before weighing. The dry mass was obtained after drying the membranes at 80 °C for a period of 24 h in a vacuum oven. The values of the WU can be calculated as follows: (𝑊ℎ𝑦𝑑 ― 𝑊𝑑𝑟𝑦)

Water uptake (%) =

𝑊𝑑𝑟𝑦

(1)

× 100%

where Wdry and Whyd represent the masses of dry and hydrated membranes, respectively. Similar to the WU, the dimensional changes were characterized by immersing the samples in a 1000 mL beaker filling with deionized water at 25 °C for a period of 24 h. The length (lhyd) and thickness (thyd) of the hydrated membranes were measured swiftly after wiping with a tissue paper. Then, the dry length (ldry) and thickness (tdry) of the membranes were measured after the membranes dehydration at 80 °C for 24 h in a vacuum oven. The dimensional changes can be calculated as follows: (𝑙ℎ𝑦𝑑 ― 𝑙𝑑𝑟𝑦)

Dimensional change (%) Δl =

𝑙𝑑𝑟𝑦

× 100%, Δt =

(𝑡ℎ𝑦𝑑 ― 𝑡𝑑𝑟𝑦) 𝑡𝑑𝑟𝑦

× 100%

(2)

The mechanical properties of the membranes were measured using a Zwick universal tester 1446-60, 1994 (Germany) at 25 °C. The membrane samples were about 20×30 mm2 and tested at the tensile rate of 1 mm min-1. The tests were performed three times, and the results were reported as an average. 10 ACS Paragon Plus Environment

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The IEC ( 𝑚𝑒𝑞.𝑔 ―1) of the membranes were measured by using a typical back-titration method. The membranes in OH− form were soaked in 25 mL of 0.1 M HCl for 48 h, after which the HCl solution was back titrated by 0.1 M NaOH standard solution. The IEC values were calculated according to: ∆𝑉𝑁𝑎𝑂𝐻 × 𝐶𝑁𝑎𝑂𝐻

IEC =

𝑊𝑑𝑟𝑦

(𝑚𝑒𝑞.𝑔 ―1)

(3)

where ∆V is the volume of the NaOH solution, C is the concentration of NaOH and Wdry is the mass of dry membrane samples. The crosslinking density in membranes was measured by the gel fraction method. Approximately 0.05 g of the dried membranes were placed in a filter paper and kept in DMF at 70 °C for 24 h. The undissolved part of the samples was dried and weighed. The gel fraction was calculated based on the following equation 41. 𝑊

Gel fraction (%) = 𝑊0 × 100%

(4)

Where W0 is the initial weight of the samples before extraction and W is the weight of the samples after extraction. The ion conductivities (𝜎) of the quaternary membranes were investigated by the electrochemical impedance spectroscopy method from a frequency range of 10 to 105 HZ 38. The hydroxide conductivity of the membranes was measured from 30 to 80 °C in an insulated test tube and full hydrated conditions by the following equation: 𝑙

𝜎𝑂𝐻 ― (𝑆.𝑐𝑚 ―1) = 𝑅𝐴

(5)


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where l is the distance between the two electrodes (cm), R is the ohmic resistance of the samples (Ω), and A is the cross-sectional area of the membrane samples (cm 2). The alkaline stability or long-term durability of the membranes was investigated in the KOH solution (2 mol L−1) at 80 °C for 12 days, with the replacement of the 2 M KOH solution every day during the testing period. The membrane electrode assembly (MEA) was prepared by mixing Pt/C (Pt content 30%) with 5% Nafion solution. The mixture was sonicated for homogeneity and then coated onto a carbon paper (CeTech W1S1005). The Pt loading for both the electrodes was 0.5 mg cm-2. The MEA with an active area of 5cm2 was fabricated by hot-pressing the anode and cathode electrodes on both sides of the semi-IPN MOF membrane at 110 °C for 2 min. Fully humidified H2 and O2 on the anode and cathode, respectively, were fed with a flow rate of 250 mL min-1. The methanol crossover was measured employing a drive mode cell. The one side of MEA was fed by 2M MeOH and the other side by air flow and followed by applying potential in the range of 0-1.2 V. The current density explores the methanol crossover through the membrane.

3.

Result and discussion

3.1. Synthesis of Cr-MIL-101-NH2 Originally, Cr-MIL-101 was synthesized through hydrothermal conditions from chromium nitrate and terephthalic acid, as previously described in the literature 42. The covalent postsynthetic modification of Cr-MIL-101 with the nitrating agent (concentrated HNO3/concentrated 12 ACS Paragon Plus Environment

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H2SO4) afforded Cr-MIL-101-NO2. Finally, Cr-MIL-101-NH2 was achieved by the subsequent reduction of Cr-MIL-101-NO2 via SnCl2.2H2O/concentrated HCl (Scheme. 1) 43. The successful covalent modification of Cr MIL-101 by amine functional groups could be followed by FT-IR spectroscopy (Fig. 1). The characteristic absorption band at around 1255 cm-1 was attributed to the C-N stretching vibration of aromatic amines (Fig. 1b). Observation of a peak splitting around 1400 cm-1 is clear evidence of intermolecular hydrogen bonding between grafted amino and carboxyl groups in the MOF. Here, the lower energy peak (1386 cm-1) is related to the C-O stretching vibrations affected by hydrogen bonding and the higher energy peak (1428 cm-1) is assumed independent of these interactions 44. 3.2. Synthesis of monomers and polyamide Aromatic diamine, (4,4'-diamino-4"-methoxytriphenylamine), was easily achieved by catalytic reduction of dinitro intermediate (1) from the nucleophilic fluoro-displacement reaction of 4fluoronitrobenzene by p-anisidine in the presence of CsF (Scheme. 2) 39. The structures of dinitro intermediate (1) and diamine monomer (2) were identified by the FT-IR and 1H NMR spectroscopic techniques. The FT-IR spectra of dinitro intermediate (1) represented two nitro characteristic bands at about 1578 cm-1 and 1313 cm-1 (–NO2 asymmetric and symmetric stretching vibration) (Fig. 2 a). After reduction, the disappearance of the characteristic vibrations of the nitro group was observed, and the asymmetric and symmetric N–H stretching vibration pair appeared at 3451 cm-1 and 3412 cm-1, respectively, which confirmed the presence of amino groups (Fig. 2 b). The 1H NMR spectrum also confirmed the successful transformation of nitro (1) to amino (2) groups by the high field shift of the aromatic protons and appearance of the signals related to the amino protons at around 5 ppm (Fig.S1 b). Eventually, polyamide (3) was 13 ACS Paragon Plus Environment

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polymerized from the diamine (2) and 4,4'-oxybis (benzoic acid) via solution polycondensation by using TPP in the NMP-pyridine solution, as depicted in Scheme. 2 45-46. The inherent viscosity of the synthesized polyamide (3) was measured in the NMP solution at 30 °C and found to be 0.54 g dL-1. From the FT-IR and 1H NMR spectroscopy, it can be reliably inferred that polyamide was achieved successfully (Fig. 2c and S1 c). The characteristic FT-IR absorption bands of the amide group were around 3350 (N–H stretching) and 1638 cm-1 (amide carbonyl). Meanwhile, the resonance peaks performing at 10.33 ppm in the 1H NMR spectrum also approved the formation of amide linkages (Fig. S1 c). For all the membrane samples, broadening the band between 3400 and 3500 cm-1 can be attributed to the stretching vibrations of OH- groups (Figs. 2d, 2h, and 2i). 3.3. Preparation of semi-interpenetrating polymer network (semi-IPN) membranes As shown in Scheme. 3, for the preparation of semi-IPN membrane, the synthesized polyamide was dissolved in DMF and then the monomers (N-vinylimidazole and N-vinyl-2-pyrrolidone), a crosslinking agent (DVB), and initiator (AIBN) were added to the resulting solution. The in situ radical polymerizations was allowed to proceed at 90 ºC for 3 days. After polymerization, for quaternization, MeI was added to the solution and stirred for another 3 days at 60 ºC. For the preparation of semi-IPN membrane with MOF, the monomers (N-vinylimidazole and N-vinyl-2pyrrolidone), DVB, and AIBN were added to CH2Cl2 as a solvent and impregnated into Cr-MIL101-NH2. After 3 days of stirring at room temperature, CH2Cl2 was removed by vacuum evaporation and the dissolved polyamide in DMF was added to the mixture. Finally, the in situ radical polymerizations was allowed to proceed at 90 ºC for 3 days. Then, MeI was added to the solution and stirred for another 3 days at 60 ºC to quaternize the polymer. Eventually, the 14 ACS Paragon Plus Environment

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synthesis of the semi-IPN and semi-IPN MOF membranes was investigated by the FT-IR technique (Fig. 2 (h, i)). By comparing the FT-IR spectra of the monomers and semi-IPN polymer networks, the most significant changes were the disappearance of the peaks at around 3043 cm-1 and 3054 cm-1 (=CH stretching vibration) and appearance of the absorption band at around 2937 cm-1 (-CH2 stretching vibration), which were attributed to the monomers and semiIPN polymer networks, respectively. The characteristic band at 1670 cm-1 was ascribed to the C=O stretching vibrations of the semi-IPN membranes. The peak at around 1388 cm-1 was observed in the semi-IPN MOF membrane and Cr-MIL-101-NH2. The FT-IR analysis quantitatively confirmed the successful incorporation of MOFs and polymerization of monomers within the polyamide matrix. 3.4. The SEM images of the semi-IPN and semi-IPN MOF membranes One of the most important parameters that influence the conductivity of membranes is a membrane's morphology 47. Consequently, the morphology of Cr-MIL-101-NH2, as well as the surface and cross-sectional morphologies of the semi-IPN and semi-IPN MOF membranes were studied (Fig. 3). The octahedral morphology of Cr-MIL-101-NH2 could be observed from the SEM image (Fig. 3a). The results showed that the membrane surfaces were clearly smooth, dense and relatively homogeneous. Moreover, the changes in the membrane morphology and development of the cavities and porosity due to the addition of MOFs were observed from the cross-sectional images. The performance of the semi-IPN MOF membrane can be significantly affected by this unique morphology.


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3.5. Thermogravimetric and BET analysis For the investigation of the thermal degradation behavior of the samples, the thermogravimetric analysis (TGA) was performed at a heating rate of 10 °C min−1 under an N2 atmosphere. As shown in Fig. 4, Cr-MIL-101-NH2 had a two-step degradation profile. The first initial weight loss between 100 and 150 ºC was attributed to the evaporation of the water and residual trapped solvents. The second weight loss above 320 ºC was related to the degradation of the Cr-MIL101-NH2 frameworks, confirming the excellent thermal stability of MOFs. For aromatic polyamide, the slight weight loss up to 150 °C may come from the evaporation of the residual solvent, and the weight loss above 300 °C was assignable to the degradation of the polymer main chain. Furthermore, for the semi-IPN membrane, the first weight loss started around 230 °C could be due to the decomposition of the copolymer's backbone; moreover, the subsequent weight loss of the membrane may be caused by the degradation of the polyamide backbone. Finally, the degradation behavior of the semi-IPN MOF membrane may be evidence for supporting the copolymer interpenetration into the cavity of MOFs 48. The first weight loss between 100 and 150 ºC for Cr-MIL-101-NH2 without trapped copolymer was attributed to the evaporation of the water and residual trapped solvents from the cavities. Nevertheless, the same weight loss stage could not be found for the semi-IPN MOF membrane degradation profile. This observation implied that the copolymer was caging within the framework, where the copolymer replaced the solvent and filled the cavities. On the other hand, the first weight loss in the semiIPN MOF membrane occurred around 280 °C that was higher than the same degradation stage in the semi-IPN membrane (230 °C). This observation could also be considered as another evidence for interpenetrating the copolymer into the cavity of MOFs, where the copolymer was protected


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from the early thermal degradation process. Eventually, the second weight loss started at 490 °C could be related to the degradation of the residual Cr-MIL-101-NH2 frameworks and polyamide skeleton. Generally, all the membranes displayed good thermal stability below 250 °C. The BET surface areas and porosity were measured to afford further evidence of caging the copolymer into pores of the MOFs. The decreasing of the pore volume from 1.019 to 0.002 cm3 g-1, implying the pores of the Cr-MIL-101-NH2 were filled mostly with the copolymer (Table. 1). 3.6. Ion exchange capacity (IEC), water uptake (WU), swelling ratio (SR), and gel fraction (GF) of the membranes The IEC, WU, and SR of the samples were investigated and the results are summarized in Table. 2. The more requirement of the hydroxyl ionic groups to maintain electrical charge neutrality was as a result of the more N+ sites. Consequently, the IEC values increased with the OHincreasing. However, excessive water absorption was induced by the large IEC, which may cause deterioration of the mechanical and dimensional stability of the membranes. As listed in Table. 2, the semi-IPN membrane showed higher WU and IEC value compared with the polyamide membrane due to the increase of the quaternary functional groups. However, decreasing in the SR was attributed to the restriction of the crosslinking structure of the semi-IPN membrane. On the other hand, in comparison with the semi-IPN membrane, the WU and IEC values of the semiIPN MOF membrane increased due to the higher N+ sites in the Cr-MIL-101-NH2 structure. However, no significant differences were observed in the SR of these membranes (semi-IPN and semi-IPN MOF) because the extra absorbed waters entered into the porosity of MOFs. The robust pore walls of Cr-MIL-101-NH2 with mechanical and structural durability protected the membrane from the excessive swelling effects 49. The gel fraction (GF) is one of the widely used 17 ACS Paragon Plus Environment

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methods as an indirect measurement of the cross-linking density of the polymer networks. In order to confirm the cross-linking reaction, the amounts of GFs of polyamide and semi-IPN membranes were determined. As shown in Table. 2, polyamide was completely dissolved in DMF at 70 °C but semi-IPN polymers partially dissolved which confirmed the successful crosslinking of these polymers 40, 50. 3.7. Mechanical properties The mechanical properties of the AEMs were evaluated at room temperature under dry and hydrated conditions and the results are presented in Table. 3. The polyamide membrane had the tensile strength and elongation at the break values of 25.50 MPa and 18.50% in the dry and 17.16 MPa and 25.68% in the hydrated conditions, respectively. However, the semi-IPN and semi-IPN MOF membranes exhibited an increase in the tensile strength and a decrease in the elongation at break in both the dry and hydrated conditions (Table. 3). The mechanical properties of the semiIPN membrane were found to be affected by the cross-linking segments. On the other side, a slight increase in the tensile strength and decrease in elongation at break of the semi-IPN MOF membrane compared with the semi-IPN membrane was attributed to the hydrogen bonding formation between amide linkages in the polyamide and –NMe2 groups in the post-modified MOFs structures. This slight increase in the rigidity of the semi-IPN MOF membrane might be due to the movement restriction of the polyamide chains caused by MOFs. The tensile strength and elongation at the break values of the semi-IPN membranes were in the range of the values of membranes showing good fuel cell performance.


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3.8. Hydroxide conductivity The fundamental purpose of this study was to synthesize AEMs with high hydroxide ion conductivity and good dimensional stabilities. It has long been known that the ion conductivity of AEMs is commonly governed by three factors: the ion intrinsic mobility, ion-exchanging group density, and nanochannels morphology 51-53. Generally, higher concentrations of quaternary functional groups, which yield higher IECs, result in excessive WU 54. On the other hand, one of the necessities of anions transportation and nanochannels continuity is the appropriate water content in AEMs 55, but the dimensional stability was reduced by excessive WU 56. Consequently, the trade-off between ion conductivity and the SR is one of the drawbacks to developing high-performance AEMs 41. In this study, we attempted to generate highly conductive AEMs by introducing well-connected ion-conducting nanochannels. With regard to this purpose, we first attempted to synthesize a semi-IPN membrane and incorporate the ionic groups into the polymer backbone with the aim of improving the hydrophilicity, IEC, and WU of the membrane for increasing the ionic conductivity. However, no significant increase in conductivity was observed (Fig. 5a), which could be because of crosslinking. Generally, crosslinking limits the development of ion nanochannels by reducing the WU 57. However, in the semi-IPN MOF membrane, by associating the polymeric chains within the hydrophilic nanopores of Cr-MIL-101-NH2, makes it possible to array MOFs and causes the interconnected nanochannels (Scheme. 4). Consequently, the resulting interconnected ionic channels operated as special highways for water and OH- transportation. Fig. 5a shows the OH- conductivities of the polyamide (as control), semi-IPN, and semi-IPN MOF membranes as a function of temperature. Generally, due to the higher mobility of OH- in higher temperatures, the conductivities of the


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membranes increased by increasing the temperature. Ascribing to its highest WU, IEC, and also well-connected anion nanochannels a superior conductivity of 0.12 S cm-1 was observed at 80 °C for the semi-IPN MOF AEM. Moreover, in comparison with the semi-IPN membrane, the semiIPN MOF AEM recorded a large jump in hydroxide conductivity (0.12 vs 0.02) due to the lack of interconnected nanochannels in the semi-IPN membrane, as mentioned previously. Hence, it was speculated that improving in ion conductivity in the semi-IPN MOF membrane may be related to (i) an increase in the hydration level and IEC of the membrane; (ii) the enhanced basic domain due to hydrophilic (–NR3 +) functional groups for OH- migration; (iii) the porous structure of MOFs that led to the well-interconnected nanochannels; and (iv) the existence of copolymer, like a robust chain, connecting the porous framework of MOFs to create an ion transport highway (Scheme. 4). Thus, the connection of the porous and robust Cr-MIL-101-NH2 frameworks into the semi-IPN MOF membrane successfully solved the obstacle between conductivity and dimensional stability of the membrane. As shown in Fig. 5b, the conductivity of the AEMs exhibited an approximate Arrhenius-type temperature dependence. The ion transport activation energy Ea (kJ mol-1) of the AEMs was calculated using the following equation: Ea= -b×R

(6)

where b is the slope of the regressed linear ln(σ) vs. 1000/T plots and R is the gas constant (8.314 J (mol K)-1). The Arrhenius plots of the membranes (Fig. 5b) showed that the activation energy (Ea) of the semi-IPN MOF membrane was lower than that of the semi-IPN membrane, indicating low resistance to ion transport in the semi-IPN MOF membrane. However, the amounts of Ea for the semi-IPN (13.63 kJ mol-1) and semi-IPN MOF (8.18 kJ mol-1) membranes were similar and somewhat lower than the results of AEMs usually reported in the literature 20 ACS Paragon Plus Environment

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(10−23 kJ mol-1). This implied that these membranes had a water-assisted OH− conduction mechanism 58. Essentially, hydroxide ions were transported along the membrane through the formation and cleavage of hydrogen bonds of a chain of water molecules (Scheme. 4) 41, 58-61. In order to consider the properties of the semi-IPN MOF membrane (reasonably) with currently reported advanced AEMs, the WU, SR, IEC, and conductivity were compared and the results are presented in Table. 4. The results clearly showed that this new anion exchange membrane could make a good candidate for possible application in AEMFCs. 3.9. Alkaline stability The chemical stability of alkaline anion exchange membranes is an essential indicator to determine their lifetime for working under operating conditions. Accordingly, an accelerated alkaline stability test of the semi-IPN and semi-IPN MOF membranes was carried out under harsh conditions (2 M KOH aqueous solution at 80 °C for 12 d). No significant decrease was observed in OH- conductivity and IEC during the test time (Fig. 5c and Table. 2). Fig. 6, also illustrated that there were no obvious changes in the surface morphology of the semi-IPN MOF membrane after the stability test. The results indicated that the prepared AEMs in addition to the high ionic conductivity and excellent thermal stability also exhibited adequate alkaline stability for long-term applications in alkaline fuel cell devices. 3.10. Methanol crossover Methanol tends to behave as a surfactant and enhanced of the free volume of polymer membranes by weakening the interaction of the polymer backbone. Methanol transport is an immediate consequence, which is commonly called “methanol crossover” 62. Ideal membranes


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for direct methanol fuel cells (DMFCs) must have both high ionic conductivity and low methanol permeability. In order to investigate the application potential of semi-IPN MOF membrane in DMFC, the methanol crossover of semi-IPN MOF membrane was measured at 70 °C (Fig. 5d). The methanol crossover of the semi-IPN MOF membrane showed 16.30 mA cm−2 current density, which was remarkably lower than Nafion 117 membrane 63. 3.11. Fuel cell test Considering the appropriate ionic conductivity of the semi-IPN MOF membrane was chosen to study with polarization and power density curves in a single cell test (Fig. 7). The single-cell performance was operated at 70 °C, and the thickness of the membrane was around 40-50 µm. The open circuit voltage (OCV) of the single cell is 0.92 V with a peak power density of 42.48 mW cm-2. Generally, the single-cell performance could be affected not only by the membrane itself but also associates with many factors such as catalyst loading, MEA fabrication procedure, ionomer preparation, and operating conditions 64. The maximum power density of this membrane is higher 65 or comparable with some other reported AEMs 66-68.

4.

Conclusions

In summary, a semi-IPN MOF membrane was synthesized via in situ copolymerizations of monomers in the pores of Cr-MIL-101-NH2 in the presence of polyamide as a linear polymer and compared with a semi-IPN membrane without MOF. The TGA profile showed an improvement in the thermal stability of the semi-IPN MOF membrane compared to the semi-IPN membrane due to the trapping of the copolymer into Cr-MIL-101-NH2 cavities. Furthermore, the 22 ACS Paragon Plus Environment

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morphological changes resulting from the contribution of MOFs was observed from the crosssectional images. Afterward, the results showed that the morphological changes and hydroxide ion transport nanochannels induced by Cr-MIL-101-NH2, in addition to the higher WU and IEC, could further facilitate the hydroxide ion transfer in the semi-IPN MOF membrane. The semiIPN MOF membrane displayed high hydroxide ion conductivity up to 0.07 S cm-1 at 30 °C (663% increase compared with the semi-IPN membrane without MOF). This large jumping in ion conductivity was due to the specific highway's construction with an array of MOFs for water and OH- transportation that accorded typically with a hopping (Grotthus) mechanism. Despite the increase in the WU and IEC, excellent dimensional stability (an SR of 7%) could be due to both the crosslinking and robust pore walls of MOFs. The semi-IPN MOF membrane also demonstrated good mechanical, alkaline stability, methanol crossover, and single cell performance. In the end, the semi-IPN MOF membrane introduced in this work provides a feasible approach to resolve the trade-off issues between conductivity and dimensional stability.

Acknowledgments The authors are grateful to the Research Affairs Division of the Isfahan University of Technology (IUT), Isfahan, and the Renewable Energy Organization of Iran (REOI) for financial support of this project.


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51. Li, Z.; Jiang, Z.; Tian, H.; Wang, S.; Zhang, B.; Cao, Y.; He, G.; Li, Z.; Wu, H., Preparing alkaline anion exchange membrane with enhanced hydroxide conductivity via blending imidazoliumfunctionalized and sulfonated poly (ether ether ketone). J. Power Sources. 2015, 288, 384-392. 52. Wu, X.; Chen, W.; Yan, X.; He, G.; Wang, J.; Zhang, Y.; Zhu, X., Enhancement of hydroxide conductivity by the di-quaternization strategy for poly (ether ether ketone) based anion exchange membranes. J. Mater. Chem. A. 2014, 2 (31), 12222-12231. 53. Liang, N.; Tu, Y.; Xu, J.; Chen, D.; Zhang, H., Hybrid anion exchange membranes with self-assembled ionic channels. Adv Polym Tech. 2018, 37 (6), 1732-1736. 54. Gu, S.; Cai, R.; Yan, Y., Self-crosslinking for dimensionally stable and solvent-resistant quaternary phosphonium based hydroxide exchange membranes. Chem. Commun. 2011, 47 (10), 2856-2858. 55. Vandiver, M. A.; Horan, J. L.; Yang, Y.; Tansey, E. T.; Seifert, S.; Liberatore, M. W.; Herring, A. M., Synthesis and characterization of perfluoro quaternary ammonium anion exchange membranes. J. Polym. Sci., Part B: Polym. Phys. 2013, 51 (24), 1761-1769. 56. Zhou, J.; Unlu, M.; Vega, J. A.; Kohl, P. A., Anionic polysulfone ionomers and membranes containing fluorenyl groups for anionic fuel cells. J. Power Sources. 2009, 190 (2), 285-292. 57. He, Y.; Wu, L.; Pan, J.; Zhu, Y.; Ge, X.; Yang, Z.; Ran, J.; Xu, T., A mechanically robust anion exchange membrane with high hydroxide conductivity. J. Membr. Sci. 2016, 504, 47-54. 58. Tanaka, M.; Fukasawa, K.; Nishino, E.; Yamaguchi, S.; Yamada, K.; Tanaka, H.; Bae, B.; Miyatake, K.; Watanabe, M., Anion conductive block poly (arylene ether) s: synthesis, properties, and application in alkaline fuel cells. JACS. 2011, 133 (27), 10646-10654. 59. Lin, B.; Qiao, G.; Chu, F.; Wang, J.; Feng, T.; Yuan, N.; Zhang, S.; Zhang, X.; Ding, J., Preparation and characterization of imidazolium-based membranes for anion exchange membrane fuel cell applications. Int. J. Hydrogen Energy. 2017, 42 (10), 6988-6996. 60. Lu, W.; Zhang, G.; Li, J.; Hao, J.; Wei, F.; Li, W.; Zhang, J.; Shao, Z.-G.; Yi, B., Polybenzimidazolecrosslinked poly (vinylbenzyl chloride) with quaternary 1, 4-diazabicyclo (2.2. 2) octane groups as highperformance anion exchange membrane for fuel cells. J. Power Sources. 2015, 296, 204-214. 61. Merle, G.; Wessling, M.; Nijmeijer, K., Anion exchange membranes for alkaline fuel cells: A review. J. Membr. Sci. 2011, 377 (1-2), 1-35. 62. Kreuer, K.-D., Ion conducting membranes for fuel cells and other electrochemical devices. Chem. Mater. 2013, 26 (1), 361-380. 63. Jiang, R.; Zhang, Y.; Swier, S.; Wei, X.; Erkey, C.; Kunz, H. R.; Fenton, J. M., Preparation via supercritical fluid route of Pd-impregnated nafion membranes which exhibit reduced methanol crossover for DMFC. Electrochem. Solid State Lett. 2005, 8 (11), A611-A615. 64. Dekel, D. R., Review of cell performance in anion exchange membrane fuel cells. J. Power Sources. 2018, 375, 158-169.


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Scheme. 1. The synthesis path for Cr-MIL-101-NH2

Scheme. 2. Synthesis of dinitro intermediate (1), diamino monomer (2), and polyamide (3) 30 ACS Paragon Plus Environment

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Scheme. 3. The synthetic route of the a) semi-IPN, and b) semi-IPN MOF membranes

Scheme. 4. The schematic illustration of the water-assisted OH− conduction mechanism in the well-connected nanochannels across the semi-IPN MOF membrane 31 ACS Paragon Plus Environment

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Fig. 1. The FT-IR spectra of (a) Cr-MIL-101, and (b) Cr-MIL-101-NH2

Fig. 2. The FT-IR spectra of the (a) dinitro compound (1), (b) diamine monomer (2), (c) polyamide (3), (d) Q-polyamide membrane, (e) N-vinylimidazole monomer, (f) N-vinyl-2-


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pyrrolidone monomer, (g) Cr-MIL-101-NH2, (h) semi-IPN membrane, and (i) semi-IPN MOF membrane (the as-prepared anion exchange membranes were in OH- form (d, h and i))

Fig. 3. The SEM images of the (a) Cr-MIL-101-NH2, (b) semi-IPN, and (c) semi-IPN MOF membranes; the cross-sections of (d) semi-IPN, and (e) semi-IPN MOF membranes


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Fig. 4. The thermogravimetric (TGA) curves of Cr-MIL-101-NH2, and AEMs


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Fig. 5. a) The temperature-dependence ionic conductivity of the AEMs in OH- form, b) the Arrhenius plots of the semi-IPN and semi-IPN MOF membranes, c) the alkaline stability of the semi-IPN and semi-IPN MOF membranes in a 2 M KOH solution at 80 ºC, and d) the Methanol crossover of semi-IPN MOF membrane at 70 ºC.

Fig. 6. The SEM images of the (a) surface of semi-IPN MOF, and (b) surface of the semi-IPN MOF membrane after the stability test

Fig. 7. The single cell performance of semi-IPN MOF membrane at 70 °C


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Table. 1. The BET measurement results Samples

BET surface area Total pore volume (m2 g-1)

(cm3 g-1)

Cr-MIL-101-NH2

1975

1.019

semi-IPN MOF

0.209

0.002

Table. 2. The water uptake, swelling ratio, and ion exchange capacity of the membranes ( 𝑚𝑒𝑞.𝑔 ―1)

WU (%)

SR (%) ∆T SR (%) ∆L

aIEC

polyamide

29 ± 2.14

17 ± 3.12

11 ± 2.10

0.85± 0.16

-

0

semi-IPN

48 ± 2.80

7 ± 1.40

2 ± 0.25

1.92 ± 0.05

1.83± 0.04

55 ± 3

semi-IPN MOF

73 ± 3.41

7 ± 0.85

2 ± 0.32

2.40 ± 0.08

2.36± 0.05

68 ± 3

aIEC

before alkaline stability test

bIEC

after alkaline stability test

Table. 3. Mechanical properties of membranes Samples

Stress (MPa) Wet

Elongation at break (%) Dry

Wet

Dry

polyamide

17.16 ± 1.30 25.50 ± 1.41

25.68 ± 1.52 18.50 ± 1.21

semi-IPN

21.85 ± 1.26 28.30 ± 1.83

13.32 ± 0.91 15.05 ± 1.10

semi-IPN MOF

23.25 ± 1.43 31.75 ± 1.35

12.82 ± 0.78 13.72 ± 0.81


bIEC

( 𝑚𝑒𝑞.𝑔 ―1)

Samples

GF (%)

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Table. 4. Comparative study of IEC, swelling ratio, water uptake, and hydroxide conductivity of semi-IPN MOF membrane in this work and other AEMs in literature. Samples

Membrane type

WU

SR

IEC ( 𝑚𝑒𝑞. Conductivity

(%)

(%)

𝑔 ―1)

(S cm-1)

Ref This

semi-IPN MOF

Semi-IPN

73

7(Δt)

2.40

0.12 (80 ºC)

PP–60

Semi-IPN

29.5

22.4

2.63

0.035 (60 ºC)

66

QPPT-20

Semi-IPN

82

-

2.31

0.06 (70 ºC )

69

PIL-MIL-101-Cr

polymer hybrid

-

-

-

0.138 (80 ºC)

25

pc-MBPPO

polymer hybrid

27

5

-

0.145 (80 ºC)

26

IL@ZIF-8/IL/PVA

polymer hybrid

-

-

-

0.09×10-2 (60 ºC)

70

SAN50-[DMVIm][OH]50 copolymer

135

43.3

1.98

0.05 (90 ºC )

59

QCPPAE-4/1

cross-linked

40

16 (Δt) 1.75

0.065 (60 ºC)

71

M3

cross-linked

70.4

20

1.33

0.028 (60 ºC)

72

RC-QPPO-2.13

cross-linked

70

21.8

2.13

0.070 (80 ºC)

73

AEM-0.36M-80

composite

22.1

10.6

2.13

0.05 (80 ºC)

74

PSU/SiO2-QPSt-70

composite

94.12

11.26

2.31

0.188 (80 ºC)

75

QPAES/10% nano-ZrO2

composite

28.1

7.4

1.82

0.036 (60 ºC)

19


study

Enhancement of hydroxide conduction by incorporation of metal

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