Hyperbranched Polystyrene Copolymer Makes Superior Anion

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Hyper-Branched Polystyrene Copolymer Makes Superior Anion Exchange Membrane Yazhi Liu, Jiahui Zhou, Jianqiu Hou, Zhengjin Yang, and Tongwen Xu ACS Appl. Polym. Mater., Just Accepted Manuscript • DOI: 10.1021/acsapm.8b00058 • Publication Date (Web): 10 Dec 2018 Downloaded from http://pubs.acs.org on December 14, 2018

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ACS Applied Polymer Materials

Hyper-Branched Polystyrene Copolymer Makes Superior Anion Exchange Membrane Yazhi Liua, Jiahui Zhoua, Jianqiu Houa, Zhengjin Yanga*, Tongwen Xua aCAS

Key Laboratory of Soft Matter Chemistry

iChEM (Collaborative Innovation Center of Chemistry for Energy Materials) School of Chemistry and Material Science University of Science and Technology of China Hefei 230026, P.R. China E-mail: [email protected]*

Keywords: hyper-branched polymers; RAFT polymerization; chain entanglement; morphology; fuel cells; anion exchange membranes.

Abstract: To meet the ever-increasing demand of energy conversion/storage devices on anion exchange membrane (AEM) conductivity, it is urgent to explore alternate polymers in terms of the backbone architecture. In contrast to linear polymers, hyperbranched polymers exhibit unexpected properties because of the unique backbone architecture. Unfortunately, reports on hyper-branched polymer based AEM are rare because weaker chain entanglement makes it hard to cast free-standing membranes. We 1

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proposed here two strategies, i.e., cross-linking and implementing flexible segments to strengthen chain entanglement, to fabricate free-standing hyper-branched AEMs. Our results confirm that implementing flexible segments imparts hyper-branched AEMs with superior properties compared with the ones fabricated via cross-linking and AEMs made from linear polymers. By exploiting the RAFT polymerization, we synthesized a hyper-branched poly(vinyl benzyl chloride) core and implemented a polyisoprene shell (the flexible segment to strengthen chain entanglement). After quaternization, a freestanding highly conductive AEM is acquired, with a conductivity of 85.1 mS cm-1 at 80 oC and limited water swelling. Such gratifying properties originate from the unique membrane morphology, and by proper control, the self-assembled pattern could be directed in the hyper-branched membranes. The obtained results prove the effectiveness of our strategy, resulting in superior AEMs for energy conversion or energy storage.

Introduction Because of the ever-increasing power our modern lifestyle demand, the awareness of environment protection as well as fossil fuel depletion, the conversion and storage of energy from renewable sources have attracted worldwide attention.1 Among the various energy conversion processes, alkaline fuel cell (AFC) can directly convert energy stored in chemicals (for instance, H2, MeOH) into electricity without relying on the noble catalyst (for example, Pt) and exhibits faster electrode reaction kinetics compared to other type of fuel cells.2 As for energy storage, aqueous organic redox flow battery 2


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(AORFB) based on water-soluble organics represents the most advanced system, especially those working in neutral pH.3 All currently reported AORFB operated in neutral pH consists of redox active materials bearing positively charged ammonium groups to increase water solubility.4 In both AFCs and neutral pH AORFBs, anion exchange membrane crucially determines the internal resistance (conversion efficiency), working lifespan of the former and the coulombic efficiency, cycle life of the latter.5, 6 However, properties of conventional AEMs can hardly meet the demand of practical application, especially in ion conductivity, which is critical to the efficiency of those systems.7 An AEM consists the polymer backbone, providing mechanical strength and positively charged ion exchange groups, along with counter-ions, forming anion conductive “channels”.8 Before carrying on our adventure towards better AEMs, we summarized and compared state-of-the-art AEMs (Figure 1).9-16 More than 50% of them is main-chain type (Figure 1a, 1c and 1g, ion exchange groups are directly attached to the polymer backbones), leading to the fact that the increase in ion conductivity can only be achieved by increasing the ion exchange capacity (IEC, representing the amount of ion exchange groups). Higher IEC value however would lead to severe water swelling and decrease the mechanical strength of AEMs, resulting in potential failure of cell assembly. This dilemma is mitigated by constructing wellconnected ion channels, increasing the efficiency of anion transport. Strategies in building more efficient anion transport channels include inducing micro-phase 3



separation by designing side-chain type AEMs (Figure 1d, 1e and 1h, even with side chains bearing densely concentrated cationic groups17) and driving block copolymers to self-assemble into ordered structures.

Figure 1．Structures and schematic illustrations of the state-of-the-art polymeric anion exchange membrane, representing major breakthrough or big advancement in the field. They are selected based on the development of polymer backbone architecture and the advanced membrane conductivity. (a), (c) and (g) are main-chain type AEMs; (b): Tröger's base AEM; (d), (e) and (h) are side-chain type AEMs; (f): N-Spirocyclic AEM. Although tuning membrane morphology seems to be an effective approach in constructing more conductive AEMs, the improvement is yet too limited for practical application and we speculate that stems intrinsically from the linear polymer chain architecture. Linear polymers are preferred and show better membrane-forming properties (Figure 2) because of stronger chain entanglement compared with for example, dendritic, star-shaped and hyper-branched polymers. Turning to a different 4


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polymer architecture would pave way to better AEMs and offer insights into the limitations brought by the currently preferred linear polymers. Nevertheless, it seems especially challenging to fabricate an AEM from hyper-branched polymers, although they are extensively investigated in nanomaterials,18 drug delivery,19 light-emitting materials,20 coatings,21 etc. This is majorly ascribed to the much weaker chain entanglement in hyper-branched polymers,22 leading to difficulties in fabricating freestanding AEMs (Figure 2). We proposed herein two synthetic strategies towards freestanding AEMs from hyper-branched polymers (As shown in Figure 3). We firstly demonstrated here cross-linking as a convenient approach to make free-standing AEMs from hyper-branched polymers. Cross-linking however requires redundant agents and makes the membranes insoluble, thus difficult to reprocess. We then exploited the reversible addition-fragmentation transfer (RAFT) polymerization, implementing flexible segments in the weakly-tangled hyper-branched polymers and obtained a freestanding AEM from hyper-branched polymer, which shows satisfying conductivity, good mechanical strength and can be repeatedly reprocessed.

Figure 2. Illustrations showing the film-forming behavior of hyper-branched polymers (a) and linear polymers (b) during solvent evaporation process. Linear polymers show better membrane-forming properties because of greater extent of chain entanglement. In contrary, hardly could a membrane be obtained from hyper-branched polymers because of limited chain entanglement. 5



With the latter strategy, we synthesized a hyper-branched polymer (HPVI), which can be well-dissolved in NMP (5 wt%) and the resulting solution was directly used to cast a transparent, free-standing and flexible membrane. After quaternization, the AEM (QHPVI) exhibited both high hydroxide conductivity (85.1 mS cm-1) and appropriate mechanical strength (6.77 MPa). Results and Discussion Hyper-branched polystyrene derivatives and hyper-branched AEMs Polystyrene and derivatives remain a most important type of polymers ever since the invention of free radical polymerization and living polymerization. They are inexpensive and commercially available, remaining however the most challenging in making flexible AEMs in hyper-branched chain architecture, because of the weak chain entanglement. We therefore chose polystyrene and its derivatives as examples to prove the effectiveness of our proposed strategies in turning common but “tough” materials into AEMs. Both hyper-branched PS-b-PVBC (HPSV) and hyper-branched PVBC-bPI (HPVI) were synthesized via RAFT polymerization with S-(4-vinyl) benzyl S’propyltrithiocarbonate (VBPT, for synthetic details see supporting information and Figure S1, S2) as macro-molecule chain transfer agent and no catalyst is required (as described in Figure 3). We chose RAFT polymerization simply because RAFT polymerization is a facile and versatile approach in synthesizing hyper-branched 6


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polymers with advantages such as mild reaction conditions, a wide range of monomers, tolerance of functionalities, and no metal catalyst required.

Figure 3. Synthetic details of QHPSV (a), QHPVI (b) and digital images of the asprepared QHPSV (A), QHPVI (B) anion exchange membranes via solution casting.

As shown in Figure 3a, we firstly synthesized hyper-branched PS (HPS) via RAFT polymerization initiated by VBPT. HPS then served as a macro-molecule chain transfer agent to graft vinyl benzyl chloride (VBC) segments, leading to HPSV. The structure of HPS and HPSV was confirmed by 1H NMR (Figure S3, Figure S4). Peaks at 4.50 ppm belong to the methylene protons in benzyl chloride group and indicate the 7



successful incorporation of VBC segments, which was implemented as an active segment to introduce quaternary ammonium functionalities. The ratio of styrene segment to VBC segment (St/VBC) is 1.7 according to the ratio of methylene proton signal integration to the integration of proton signals from the benzyl ring (6.30 ppm7.42 ppm). We have tried to dissolve HPSV in NMP (5%) and cast the solution on a clean glass plate. However, after solvent evaporation, broken pieces were obtained instead of a free-standing membrane. We then prepared an NMP solution (10%) of HPSV and N,N,N’,N’-Tetramethyl-1,6-hexanediamine (4:1) and the same casting procedure was conducted. During solvent evaporation, N,N,N’,N’-Tetramethyl-1,6hexanediamine reacts with benzyl chloride segments, resulting in a cross-linked network with positively charged quaternary ammonium groups. Eventually, a freestanding, flexible and yellowish AEM was obtained (Figure 3a, right-hand side). However, it is not easy to control such process since sometimes gels formed before the solution could be cast (even worse when the room temperature is high) and the resulting AEMs cannot be dissolved again. Although this could be considered as an effective strategy to fabricate AEMs from a hyper-branched polymer and as we have proved in a previous publication23, the sophisticated process and uncontrollability make it hard for large-scale application and recycling the membrane waste after its service lifetime would raise further issues. Alternatively, we synthesized hyper-branched PVBC (HPV) in a VBPT initiated RAFT polymerization. The resulting HPV is used as macro-molecule chain transfer 8


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agent, initiating isoprene polymerization and a flexible polyisoprene (PI) segment was implemented, leading to HPVI (Figure 3b). The structure of HPVI was confirmed via 1H

NMR (Figure S5, Figure S6) and signals at 5.10 ppm belongs to the protons in

CH=CH group, implying the successful incorporation of isoprene segment. By comparing the integration of proton peaks at 5.10 ppm with integration of proton signals from the benzyl ring (6.30 ppm-7.42 ppm), we found the ratio of isoprene units to VBC units (PI/PVBC) is 0.7. The molecular weight of HPVI measured by gel permeation chromatography(GPC)is Mn=36500 g mol-1 or Mw=75800 g mol-1, against standard polystyrene samples and the PDI is 2.23. The as-prepared HPVI was dissolved in NMP and a free-standing, flexible and transparent membrane was obtained by directly casting the solution on a clean glass plate (Figure 3b, right-hand side). Immersing the resulted HPVI membrane in aqueous trimethylamine solution initiates the reaction between benzyl chloride in VBC units and trimethylamine, resulting in positively charged functional groups in the membrane (an AEM is therefore obtained, QHPVI). Theoretically, if all VBC units are quaternised, IEC value of the resulting AEM is 2.43 mmol g-1. And the titrated value is 1.26 mmol g-1. The discrepancy is caused by the heterogeneity of the quaternization reaction. To fully understand the impact of polymer backbone architecture on properties of the resulting AEM, we thoroughly investigated the properties of QHPVI, in terms of conductivity, morphology and water swelling behavior, etc. The reported linear AEMs, QPPO and QPSV, were selected as control. 9



Membrane conductivity and alkaline stability The conductivity of AEMs will determine the efficiency of, for instance, anion exchange membrane fuel cells and neutral aqueous redox flow batteries. We therefore measured the anion (OH−, Cl−) conductivity of QHPVI membrane (IEC=1.26 mmol g1)

in water at various operating temperature. And its conductivity was compared with

the benchmark linear QPSV, with similar IEC of 1.36 mmol g-1 (Figure 4). 24

Figure. 4 Cl− and OH− conductivity of QHPVI membrane and OH− conductivity of linear QPSV membrane as a function of operating temperature. The chloride conductivity of QHPVI at 30 oC is 19.8 mS cm-1, which is steadily increased to 38.4 mS cm-1 at 80 oC. The hydroxide conductivity of QHPVI at 30 oC is 50.8 mS cm-1, which climbs up to 85.1 mS cm-1 at 80 oC. QHPVI shows higher hydroxide conductivity because hydroxide ion moves intrinsically faster than chloride ion under the same condition (OH−/Cl−=2.5).25 As for the benchmark linear QPSV with 10


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similar IEC value, a much lower hydroxide conductivity was reported.25 The hydroxide conductivity of linear QPSV at 30 oC is ~7 mS cm-1 and remains low (~15 mS cm-1) even at 80 oC.25 Comparatively, the hydroxide conductivity of QHPVI at 30 oC (50.8 mS cm-1) is >7 times that of linear QPSV (6.9 mS cm-1), implying an obviously improved hydroxide conducting efficiency in hyper-branched polystyrene. In addition, what we learned here is that different anions appear to have different mobility/conductivity in the membrane and we think that to some extent could reflect the selectivity of the membrane. But without more information, we cannot give a solid conclusion here. We are currently considering evaluating the membranes for ion separations. Noting that to avoid the interference of ambient CO2, the deionized water used for conductivity measurements was Ar purged prior to use. One would argue that OHanions present in the membrane can be chemically transformed to bicarbonate or carbonate ions and the latter ions have less mobility than OH- ions. To prove the accuracy of our results, we compared the conductivity of the as-prepared membrane in both HCO3- form and OH- form. . The OH- conductivity we obtained is ~4 times higher than that of HCO3-, as shown in Figure S7, which matches the results reported in the references.26 And we have found the improved hydroxide conductivity is probably caused by the formation of patterned morphology within QHPVI, a unique phenomenon brought by the hyper-branched polymer architecture (discussed in a later section).

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The alkaline stability of QHPVI membrane was investigated by immersion in 1 M NaOH aqueous solution at 80 oC for 10 days. During alkaline treatment, the membrane samples were taken out and the hydroxide conductivity were measured. As shown in the Figure 5, it turns out that the hydroxide conductivity remains almost unchanged over a prolonged period of time.

Figure 5. OH- conductivity variation of the QHPVI membrane during the alkaline treatment for 10 days.

For a more intuitive comparison of the performance, the hydroxide conductivity and water uptake (WU) of QHPVI at 80 oC are compared with the reported ones (Table 1), with various polymer backbone architecture. And by comparison, we are convinced that hyper-branched polymer architecture would pave the way towards the advanced AEMs, especially when we can finely tune its morphology.

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Table 1. Ion exchange capacity (IEC), water uptake (WU), and OH− conductivity of QHPVI compared with those of representative anion exchange membranes. IEC Membrane

(mmol g-1)

σOH−

WU Configuration

(mS

cm-1)

Ref

(wt%)

PPO-1Q-1.5

1.5

11.9

31.7

27

PPO-7Q-1.8

1.8

18.7

60.1

27

gQAPPO

1.78

52.5

53.6

28

QSPES-40

1.56

42.5

37

29

QPMBV-APE

1.28

43

30

30

XPP-DMHDA

1.42

52

21

31

PEEK Q100

1.43

41

41.7

32

PES-B80-C16

1.18

29.0

8.2 33

PES-B100-C16

1.43

44.0

12.6

X60Y15

2.76

28

31

X60Y60

2.42

18

18 34

X80Y20

3.20

40

38

13



QHPVI

1.26

85.1

51.2

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This work

Membrane morphology The difference between hydrophilic segments and hydrophobic segments could drive the self-assembly of membranes after the annealing process, resulting in obvious microphase separation and various nanostructures.35 QHPVI contains hydrophilic segments (QPV) and hydrophobic segments (PI) and because of the difference in hydrophilicity and hydrophobicity, a patterned morphology could be thus obtained via polymer selfassembly, which could hardly be fulfilled in linear homopolymers. We monitored the transition of QHPVI surface morphology in SEM after solution casting and after further solvent annealing. QHPVI membrane prepared via solution casting method shows patterned surface morphology, consisting alternating ridges and valleys (Figure 6a, 6b). Although quantitative description the depth of the valley and width of the ridge seems hard, the patterned structure on membrane surface could be observed. We believe the pattern membrane surface morphology is caused by the self-assembly of QHPVI chain segments because of the difference in hydrophilicity/hydrophobicity. The ridges or valleys could probably constitute the ion conduction channels. To make such pattern

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more obvious, the QHPVI membrane sample was further annealed in chloroform vapor for 40s.

Figure 6. SEM surface images of QHPVI (a, b) and solvent annealed QHPVI (c, d). Solvent annealing of QHPVI was conducted in chloroform vapor for 40s.

A more evident surface pattern was observed in its SEM images (Figure 6c, 6d). Such transition is triggered by further self-assembly of the hyper-branched QHPVI chain segments. When treated with solvent vapor, the ionic segments (QHPV cores) are forced to aggregate, facilitating the conduction of anions. The driving force of this process is provided by the solvent (chloroform) induced self-aggregation of non-ionic segments (flexible PI shells). Compared with AEMs from linear polymers with similar or even higher IEC,36 the patterned surface morphology of QHPVI due to chain segment self-assembly is a major cause of its superior conductivity. By finely tuning surface 15



morphology, better AEMs could be obtained and that is what we are going to devote all our research efforts to. Water swelling and thermal stability The hyper-branched AEM exhibits a water uptake of 38.6 wt% and a swelling ratio of 36.3 wt% at 30 oC with an IEC of 1.26 mmol g-1 (Table S1). As the temperature is elevated from 30 oC to 80 oC, the water uptake of QHPVI was slightly increased from 38.6 wt% up to 51.2 wt%, while the linear swelling ratio remains almost unchanged (See Supporting Discussions). Thermogravimetric analysis suggests that QHPVI has sufficient thermal stability for practical application (Figure S8). QHPVI also shows excellent mechanical properties, with tensile strength of 6.77 MPa and elongation at break of 4.48% (Table S2). All the properties satisfy requirements of membrane electrode assembly in a fuel cell apparatus.

Conclusions In summary, we have prepared hyper-branched type anion exchange membranes (AEMs) from cheap monomers. We proved here the challenge in fabricating freestanding hyper-branched AEMs could be solved by two synthetic strategies, diamine cross-linking and implementing soft segments to enhance chain entanglement. The latter strategy is more promising because the resulting AEMs could be reprocessed; the quaternization process is easy to control and the resulting AEM provides more efficient 16


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hydroxide conductive channels, while in the meantime maintaining suppressed water swelling even at elevated temperature. By proper control of the membrane formation process or by post-annealing, an ordered morphology could be easily induced. The results we obtained here indicate that hyper-branched AEMs would open up a new avenue for improved anion conductivity and may eventually lead to improved efficiency for practical energy conversion/storage devices. Supporting Information. Scheme S1. The synthesis of S-(4-vinyl) benzyl S’-propyltrithiocarbonate (VBPT) Figure S1. The 1H NMR spectra of VBPT. Figure S2. The 1H NMR spectra of HPS polymers. Figure S3. The 1H NMR spectra of HPSV polymers. Figure S4. The 1H NMR spectra of HPV polymers. Figure S5. The 1H NMR spectra of HPVI polymers. Table S1. Water uptake (WU) and swelling ratio (SR) of OH－ formed QHPVI AEMs (IEC=1.26 mmol g-1) at varied temperatures. The QHPVI membrane exhibits good swelling resistance at elevated operating temperature while increased water content. Table S2. Molecular weight, polydispersity index (PDI), ion exchange capacity (IEC), hydroxide conductivity (30 oC), water uptake (WU, 30 oC), yielding stress and elongation at break of QHPVI membrane and QPPO membrane. Figure S7 OH- conductivity and HCO3- conductivity of QHPVI (IEC=1.26 mmol g-1). Figure S8. The thermogravimetric analysis and DTG curves of QHPVI (a) and HPV (b), heating from room temperature to 900 °C at a heating rate of 10 °C min-1 under N2. Author Contributions Zhengjin Yang and Yazhi Liu conceived the idea and designed the experiments; 17



Yazhi Liu, Jiahui Zhou and Jianqiu Hou. conducted the experiments and analyze the data; Yazhi Liu, Zhengjin Yang, Tongwen Xu wrote the paper. Acknowledgments Financial support from the National Science Foundation of China (Nos. 21506201, 21720102003, 91534203), Key Technologies R&D Program of Anhui Province (No. 17030901079), K.C.Wong Education Foundation (2016-11) and International Partnership Program of Chinese Academy of Sciences (No. 21134ky5b20170010) is acknowledged. References 1.

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TOC:


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Figure 1 150x99mm (600 x 600 DPI)


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Figure 2 148x65mm (600 x 600 DPI)



Figure 3 152x119mm (300 x 300 DPI)


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Figure 4 79x68mm (300 x 300 DPI)



Figure 5 80x60mm (300 x 300 DPI)


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Figure 6 120x83mm (300 x 300 DPI)



TOC 99x61mm (300 x 300 DPI)


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Hyperbranched Polystyrene Copolymer Makes Superior Anion

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