Letter pubs.acs.org/macroletters
Poly(terphenylene) Anion Exchange Membranes: The Effect of Backbone Structure on Morphology and Membrane Property Woo-Hyung Lee,† Eun Joo Park,†,‡ Junyoung Han,† Dong Won Shin,† Yu Seung Kim,‡ and Chulsung Bae*,† †
Department of Chemistry and Chemical Biology, Rensselaer Polytechnic Institute, Troy, New York 12180, United States MPA-11: Materials Synthesis and Integrated Devices, Los Alamos National Laboratory, Los Alamos, New Mexico 87545, United States
‡
S Supporting Information *
ABSTRACT: A new design concept for ion-conducting polymers in anion exchange membranes (AEMs) fuel cells is proposed based on structural studies and conformational analysis of polymers and their effect on the properties of AEMs. Thermally, chemically, and mechanically stable terphenyl-based polymers with pendant quaternary ammonium alkyl groups were synthesized to investigate the effect of varying the arrangement of the polymer backbone and cation-tethered alkyl chains. The results demonstrate that the microstructure and morphology of these polymeric membranes significantly influence ion conductivity and fuel cell performance. The results of this study provide new insights that will guide the molecular design of polymer electrolyte materials to improve fuel cell performance.
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However, when all relevant factors including hydroxide conductivity, alkaline stability, and synthetic versatility are considered, QA moieties remain as the most widely used hydroxide-conducting headgroup for AEMs. Designing a new polymer architecure for AEMs to improve ion conductivity is a complicated task because it requires profound understanding of the mechanisms of ion transport and their relationship to polymer structure. Various PEMs with flexible tethered side chains and comb-shaped polymer architectures (e.g., Nafion) have been synthesized to enhance the degree of microphase separation in morphology. These materials have demonstrated superior proton conductivity compared with that of their analogues in which sulfonic acid groups are attached directly along the rigid aromatic chains.6 QA groups on AEMs are ordinarily prepared via halomethylation followed by quaternization, which produces AEMs with QA groups closely attached to the aromatic polymer backbones. These AEMs tend to give poor alkaline stability (particularly above 80 °C), because the presence of a positively charged nitrogen cation at the activated benzylic position promotes nucleophilic attack of hydroxide anions. The close proximity of the QA group to the polymer backbone also induces high swelling of the membrane and less-ordered, phase-separated morphologies, which result in low hydroxide conductivity.
uel cells are among the most promising candidate technologies for addressing the impending energy crisis. Anion exchange membrane (AEM) fuel cells have significant advantages over proton exchange membrane (PEM) fuel cells, including faster kinetics of oxygen reduction reaction and the capacity to use nonprecious metal catalysts (e.g., silver, cobalt, nickel).1 However, the development of commercially viable AEM fuel cells remains challenging because of the lack of highly conductive, durable AEMs under basic operating conditions. First, hydroxide conductivity in AEM fuel cells is generally lower than proton conductivity in PEM fuel cells because of the inherently lower mobility of hydroxide ions and the insufficient dissociation of hydroxide ions from the anion conductive group.2 Second, compared to perfluorosulfonated ionomer in PEMs, most aromatic AEMs have less developed ion channels in microstructure, which tends to retard ion transport.3 Third, most aroamtic AEM materials prepared by condensation polymers contain aryl ether bonds in their backbone, thus, they cannot survive in the most desirable AEM fuel cell operating environment (i.e., highly alkaline conditions above 80 °C) in which strongly basic hydroxide ions are abundant.4 These limitations of AEMs are key bottlenecks in the development of high-performance AEM fuel cells. Numerous strategies have been used to improve hydroxide conductivity and thermochemical stability of AEMs. These strategies can be classified into two main types: (a) the synthesis of new anion conductive groups, and (b) the design of new polymer architectures.5 To date, various ion conductive moieties have been investigated as new tethered ionic groups to replace conventional quaternary ammonium (QA) in AEMs.2 © XXXX American Chemical Society
Received: February 25, 2017 Accepted: April 28, 2017
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DOI: 10.1021/acsmacrolett.7b00148 ACS Macro Lett. 2017, 6, 566−570
Letter
ACS Macro Letters Recently, long flexible side chains terminated by QA groups have been used to facilitate the formation of microphaseseparated morphology in AEMs.7 Our group also reported high-molecular-weight, solvent-processable AEMs based on poly(biphenyl alkylene)s (BPN1-m, where m indicates the mol % of the ionic repeating unit; Figure 1).7b Compared with other
Bromoalkyl-tethered precursor polymers (p-TPBr and mTPBr) were synthesized via one-pot, acid-catalyzed Friedel− Crafts polycondensation (Scheme S1, Figure S1). The bromopentyl group on the side chain of these precursor polymers was subsequently converted to trimethylpentylammonium bromide via a substitution reaction with trimethylamine, affording p-TPN1 and m-TPN1. The weight-average molecular weights (Mw) of p-TPBr and m-TPBr were determined with gel permeation chromatography (Table 1). Table 1. Properties of Bromoalkyl-Tethered Precursor Polymers polymer
Mna
Mwa
PDI
DFb (%)
p-TPBr m-TPBr BPBr-65c
30.9 62.2 72.1
61.3 126.2 138.8
1.98 2.02 1.93
100 100 65
a In kg/mol. bDegrees of functionalization based on 1H NMR (Figure S1). cLee et al. ACS Macro Lett. 2015, 4, 814.
The molecular weight of p-TPBr is lower than that of m-TPBr, likely because the former has lower solubility owing to the more rigid para-terphenyl moieties. Compared with the terphenyl polymers, biphenyl-based BPBr-65 had a higher molecular weight owing to the smaller and less rigid biphenyl in the repeating unit and better solubility. The chemical structures of the ionic polymers were analyzed with 1H NMR spectroscopy (Figure S2). Comparison of the integral ratio of trimethylammonium protons at 3.02 ppm and aromatic proton peaks at 7.30−7.80 ppm indicated that the quaternization reaction gave quantitative conversion. IECs calculated from the relative intensity of the repeating units were 2.12 and 2.15 mequiv/g for p-TPN1 and m-TPN1, respectively, which are in good agreement with the IECs determined with titration (2.16 and 2.18 mequiv/g; see Tables 2 and S1).
Figure 1. Chemical structures of para-terphenyl (p-TPBr, p-TPN1), meta-terphenyl (m-TPBr, m-TPN1), and biphenyl (BPBr, BPN1) polymers where R = (CH2)5Br and R′ = (CH2)5N+(CH3)3Br−.
AEM materials prepared from typical base-catalyzed condensation polymers, these aryl ether-free aromatic AEMs exhibit superior chemical and mechanical stabilities. Although additional study is needed, we believe that the incorporation of QA at the end of long alkyl chains in aromatic polymers improves the mobility of hydroxide ion. When the morphology of these BPN1 membranes was examined with small-angle Xray scattering (SAXS), however, no obvious scattering peak was observed. While investigating membrane properties of different backbone structures, we discovered that a proper arrangement of polymer backbone is beneficial in creation of ion-conductive polyelectrolyte membrane morphology with enhanced hydroxide ion conductivity. Thus, we herein report the results of our study of the influence of backbone structure of anionconducting polymers on morphology and the overall performance of AEMs. We synthesized terphenyl-based hydroxide ionconducting polymers using either para-terphenyl (p-TPN1) or meta-terphenyl (m-TPN1) as an aromatic monomer (see Figure 1). When biphenyl was used as an aryl monomer,7b because the acid-catalyzed polycondensation is highly regioselective only for 4- and 4′-positions,8 the resulting BPN1 polymers did not allow further structural variations in the polymer backbone. However, the use of two different terphenyl monomers may demonstrate how changing the orientation of the backbone influences the properties of these side chainfunctionalized polymers; p-TPN1 and m-TPN1 have the same concentration of ionic side chains and differ only in the substitution of the middle aromatic ring (para vs meta). To investigate the effect of changes in polymer backbone structure, we analyzed the physical and morphological properties of these two terphenyl polymers and compared them with a biphenyl polymer (i.e., BPN1-65) with a similar ion exchange capacity (IEC). Our results provide new insights for the design of highperformance anion-conducting polymers.
Table 2. IECs, Water Uptake, Hydration Number, and Hydroxide Conductivity of Ionic Polymers σe (mS/cm) a
polymer
IEC
p-TPN1 m-TPN1 BPN1-65 BPN1-100
2.12 2.15 1.94 2.70
IEC
b
2.15 2.13 1.93f 2.80f
c
WU (%) 17 25 34 61
(65) (70) (85) (124)
−
λ (OH )
30 °C
80 °C
17 18 24 26
43 54 41 62
81 112 88 122
d
a IEC values were determined by 1H NMR analysis (mequiv/g). bIECs by 1H NMR analysis after alkaline test at 1 M NaOH at 95 °C for 30 days (mequiv/g) with the error range of ±0.05. cWater uptake was measured in both Cl− and OH− form at 80 °C. The values in brackets represent water uptake in OH− form. dBased on water uptake value at 80 °C. eHydroxide conductivity. fAlkaline test at 80 °C (from Lee et al. ACS Macro Lett. 2015, 4, 814).
The thermal stability, chemical stability under alkaline conditions, and mechanical durability of the membranes fabricated from these polymers were evaluated for practical applications in fuel cells. Both p-TPN1 and m-TPN1 showed good thermal stability up to 250 °C, at which the QA group started to decompose, as shown with TGA (Figure S3). Chemical stability testing of these membranes in an alkaline environment was conducted in 1 M NaOH at 80 °C, and experimental IECs were measured with back-titration and 1H NMR spectra after 30 days of treatment (Table S1 and Figure 567
DOI: 10.1021/acsmacrolett.7b00148 ACS Macro Lett. 2017, 6, 566−570
Letter
ACS Macro Letters
substituted foldamers,9 while the backbone chains of paraterphenyl tend to spread out to prevent such conformation (Figure 3). The difference in morphological behavior of these
S4). Similar to the results of a previous study of BPN1 polymer,7b the differences in IEC values and 1H NMR spectra of terphenyl ionic polymers were negligible, indicating excellent long-term thermochemical stability. Even after treatment in 1 M NaOH at 95 °C for 30 days, the decrease in IEC was negligible (Table 2). We conclude that a polymer design based on a flexible alkyl chain-tethered QA headgroup and aryl etherfree bonds on backbone is a key to the excellent alkaline stability of these aromatic AEMs. When the mechanical properties were examined under conditions of 50 °C with 50% relative humidity, the terphenyl polymers had elongations at break that were lower than that of BPN1-65 because their additional aromatic ring along the polymer backbone structure makes them less flexible than their biphenyl counterpart (Figure S5). Compared with p-TPN1, m-TPN1 showed greater stress and elongation at break owing to a combination of a more flexible backbone and higher molecular weight. The development of ionic clusters via nanoscale phase separation of hydrophobic−hydrophilic domains helps to facilitate ion transport in fuel cell membranes.5b,d The microstructure of membranes was examined with small-angle X-ray scattering (SAXS). As shown in Figure 2, biphenyl
Figure 2. (a) SAXS and (b) WAXS data of ionic polymers and TEM images of (c) p-TPN1 and (d) m-TPN1.
Figure 3. 3D model structure of the ionic polymers. (a) p-TPN1 (10 repeating units), (b) p-TPN1 (48 repeating units), (c) m-TPN1 (10 repeating units), and (d) m-TPN1 (48 repeating units). Color code: white for hydrogen, gray for carbon, red for nitrogen, blue for fluorine.
polymer BPN1-65 showed no scattering peak in the SAXS spectrum. However, Figure 2a shows large-width scattering peaks at 0.96 and 1.2 nm−1 corresponding to d-spacing of approximately 6.5 and 5.0 nm, respectively, for the m-TPN1 membrane. By contrast, those two peaks appeared with very low intensities in the p-TPN1 membrane, indicating that the level of microphase separation for this polymer is insignificant. The difference in their morphologies arose from the different connectivity of repeating units. A 3D model structure study suggests that meta-connectivity of terphenyl units allows the polymer chains to fold back to maximize the interaction between hydrocarbon backbones and promote the peripheral formation of ion aggregates, similar to the case of rigid meta-
two terphenyl membranes is also evidenced in TEM images in which m-TPN1 membrane sample stained with tetrachloroplatinate ions shows greater contrast indicating more obvious hydrophobic−hydrophilic phase separation (Figure 2c,d). The effect of terphenyl arrangement on the polymer chain packing was further examined by wide-angle X-ray scattering (WAXS; Figure 2b). While there is no obvious peak observed in WAXS for p-TPN1, m-TPN1 showed a strong peak at 2θ angle of 8.35°, corresponding to interchain spacing of 10.6 Å. This spacing is believed to correlate with the intermolecular distance between the self-assembled polymer structures, as graphically illustrated in Figure S6. A similar regular arrangement of side chains has been frequently observed for planar 568
DOI: 10.1021/acsmacrolett.7b00148 ACS Macro Lett. 2017, 6, 566−570
Letter
ACS Macro Letters conjugated polymer systems.10 It is likely that the presence of additional kinked structure from meta-substituted polymer could contribute the polymer chains to fold back and allow to form more compact polymer structure and promote selfassembly of ionic side chains. Therefore, we concluded that the difference in the backbone structure (para- vs meta-terphenyl) has influenced the self-assembly of polymer chains, which further affects the ion conductivity and morphology of the polymer membranes. For AEM fuel cell applications, water uptake and hydroxide conductivity are critical. An ideal material for AEMs should have good conductivity with manageable water uptake. The water uptake values and swelling ratios of terphenyl ionic polymers in chloride or hydroxide ion forms were measured after immersing the membranes in water at 30 and 80 °C for 24 h (Tables 2 and S2). Biphenyl AEMs, one with a similar IEC (i.e., BPN1-65) and one with a higher IEC (i.e., BPN1-100), were studied for comparison. As expected, the water uptake values of the membranes increased with increasing temperature because the mobility of the alkyl side chains containing QA groups increased at elevated temperature and provided a larger space for water absorption. Compared with BPN1-65, the terphenyl-based ionic polymers exhibited lower water uptakes at 80 °C despite having slightly higher IEC values (Table 2). This result may be due to their more hydrophobic rigid backbone structures composed of three aromatic rings. The more hydrophobic nature of the terphenyl polymers compared to biphenyl polymers is also reflected by the lower hydration number (λ, see Table 2). Accordingly, terphenyl AEMs showed lower swelling ratios than biphenyl AEMs (Table S2). Between the two terphenyl AEMs, p-TPN1 showed slightly lower water uptake and a smaller swelling ratio, which may be attributed to the more rigid structure of its polymer backbone chains. The anion conductivities of these polymers were measured with a four-point probe method. The chloride and hydroxide ion conductivities of the p-TPN1 and m-TPN1 membranes both increased as temperature rose from 30 to 80 °C (Tables 2 and S2, and Figure 4b). At room temperature, p-TPN1 and m-
close to that of BPN1-65 (88 mS/cm), although the water uptake of the p-terphenyl polymer is much lower than that of the biphenyl polymer (65% vs 85% in OH− form; see Table 2). Although the two terphenyl-based AEMs have similar IEC values (2.12−2.15 mequiv/g), m-TPN1 has much higher hydroxide ion conductivity. The hydroxide conductivity of mTPN1 nearly reaches that of BPN1−100 (122 mS/cm), which has much higher IEC (2.70 mequiv/g) and hydration number (λ = 26), demonstrating more favorable morphology for hydroxide ion transportation. As indicated by the morphology data in Figure 2, unlike BPN1-65 and p-TPN1, which show no or less ordered morphology, the higher ion conductivity of mTPN1 can be attributed to better formation of ionic clusters by self-assembly of hydrophobic polymer backbone and iontethered side chains. Figure S7 shows the fuel cell performance of p-TPN1 and mTPN1 as AEMs at 80 °C. The open-circuit voltage was 1.01 V, which is close to the theoretical value in hydrogen-supplied fuel cells. The maximum power density obtained from cells using mTPN1 was 196 mW cm−2 at a current density of 358 mA cm−2. This power density was higher than that achieved with p-TPN1 (154 mW cm−2 at a current density of 309 mA cm−2) under the same conditions despite the similarity in IEC values of the polymers. The higher power density from m-TPN1 may be attributable to the (a) greater hydroxide conductivity derived from the ordered morphology structure and (b) higher molecular weight of the polymer, which may help to build a more effective three-phase boundary in the catalyst layer. Membrane electrode assemblies using the meta-terphenyl membrane also showed low high-frequency resistance (HFR;
