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Mechanically Robust, Sodium-Ion Conducting Membranes for Nonaqueous Redox Flow Batteries Rose Ruther, Guang Yang, Frank M. Delnick, Zhijiang Tang, Michelle Lehmann, Tomonori Saito, Yujie Meng, Thomas A. Zawodzinski, and Jagjit Nanda ACS Energy Lett., Just Accepted Manuscript • DOI: 10.1021/acsenergylett.8b00680 • Publication Date (Web): 18 Jun 2018 Downloaded from http://pubs.acs.org on June 20, 2018

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ACS Energy Letters

Mechanically Robust, Sodium-Ion Conducting Membranes for Nonaqueous Redox Flow Batteries Rose E. Ruther,a Guang Yang,a Frank M. Delnicka,z Zhijiang Tang,a,b Michelle L. Lehmann,a Tomonori Saito,a Yujie Meng,c Thomas A. Zawodzinski Jr.,a,b and Jagjit Nanda,a,b,z a Oak Ridge National Laboratory, Oak Ridge, TN 37831 USA b Department of Chemical and Biomolecular Engineering, University of Tennessee, Knoxville, TN 37996, USA cNanomechanics, Inc., Oak Ridge, TN 37830, USA z

Corresponding Authors: E-mail address: [email protected] Telephone :(865)241-8361; E-mail address: [email protected] Telephone; (865) 576-8352

Abstract Sodium-based batteries are promising for grid storage applications due to significantly lower cost compared to lithium-based systems. The advancement of solid state and redox-flow sodium-ion batteries requires sodium-ion exchange membranes with high conductivity, electrochemical stability, and mechanical robustness. This study demonstrates that membranes based on poly(ethylene oxide) (PEO) can meet these requirements. Membranes plasticized with tetraethylene glycol dimethyl ether (TEGDME) achieve high ionic conductivity. Plasticized PEO membranes containing sodium triflate salt (NaTFS) show about 2 orders of magnitude higher conductivity compared to non-plasticized PEO membranes. Results from vibrational spectroscopy and differential scanning calorimetry (DSC) describe the coordination chemistry in these multi-phase materials and explain the mechanisms behind the increased conductivity. The mechanical properties of the membranes improve by addition of 5 wt.% sodium carboxymethyl cellulose (CMC) without compromising the conductivity or electrochemical stability against sodium metal. The optimized membrane is an excellent candidate for low-cost energy storage systems that operate over a wide voltage window near ambient temperature. TOC Figure

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Batteries that use sodium-ions to carry charge are becoming increasingly competitive with Libased systems for grid-scale energy storage.1-7 Sodium is abundant, inexpensive, and possesses a suitably low redox potential to enable high voltage cells. Many different types of sodium-based batteries have been developed or proposed including high-temperature secondary batteries (Na/S and Na/NiCl2),8 redox-flow batteries,9 and solid state batteries.10-12 While high temperature sodium batteries use a ceramic alumina membrane, the development of sodium-based redox-flow and solid state batteries has been slowed by the lack of suitable membranes that operate near ambient temperature. Sodium-based, nonaqueous redox-flow batteries are especially promising for grid storage. The use of organic electrolytes with a large electrochemical window could double or triple energy density compared to aqueous systems.13 While the vast majority of membranes for redox-flow batteries have been developed for aqueous electrolyte,14-16 there is increasing recognition of the need to develop special polymers specifically for nonaqueous systems.16-20 Our strategy is to modify poly(ethylene oxide) (PEO) membranes to improve Na-ion conductivity and mechanical strength for grid storage applications. While PEO has been extensively studied as a polymer electrolyte for batteries,21 the conductivity is generally too low (10-7 – 10-6 S/cm) for practical applications near ambient temperature. The introduction of plasticizers (typically solvent additives) can significantly improve conductivity.22-24 Tetraethylene glycol dimethyl ether (TEGDME), a linear oligoether with the same ethylene oxide (EO) repeat unit as PEO, was chosen as the plasticizer for this study because it does not reduce the electrochemical window of the membranes. Both PEO and TEGDME have excellent electrochemical stability near the sodium potential.25 This enables the use of redox-active couples with very negative potentials, which is desirable for high-energy batteries. The addition of plasticizers improves conductivity but worsens the mechanical properties of the membranes, which are critical for large-scale deployment. Blending PEO with carboxymethyl cellulose (CMC) strengthens the polymer electrolyte film, providing an inexpensive and environmentally friendly alternative to ceramic nanofillers.26 CMC is also stable at very negative potentials (near Na/Na+) and, therefore, the addition of CMC does not reduce the voltage window of the membrane.25 Herein, the conductivity and apparent activation energy for ion transport in PEO membranes made with sodium triflate (NaTFS) salt, PEO, TEGDME, and CMC are reported. DSC, Raman spectroscopy, and Fourier transform infrared spectroscopy (FTIR) distinguish different phases, describe the ion coordination, and explain trends in membrane conductivity. Nanoindentation evaluates the mechanical properties of the different membranes. Finally, device testing in symmetric cells with sodium metal electrodes establishes the chemical and electrochemical stability of the membranes.

NaTFS-PEO-TEGDME membranes were fabricated by casting aqueous solutions containing NaTFS, PEO, and TEGDME onto a Teflon surface or directly onto stainless steel electrodes and evaporating the water. Na-CMC (5 wt.% based on dry film weight) was added to the casting solutions to provide structural reinforcement to the dry films. PEO (M.W. 600,000), Na-CMC (M.W. 700,000), NaTFS, and TEGDME were purchased from Sigma-Aldrich and used without further purification. Three different membranes were studied: 1:8 NaTFS:PEO, 1:8:2 NaTFS:PEO:TEGDME, and 95 wt.% 1:8:2 NaTFS:PEO:TEGDME + 5 wt.% CMC. The film 2 ACS Paragon Plus Environment

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compositions are expressed as molar ratios. For example, 1:8:2 refers to 1 mole of NaTFS: 8 moles of the EO repeat unit in PEO: 2 moles of TEGDME. Note that each mole of TEGDME has five EO repeat units. After drying in air, the membrane films were vacuum dried for ~48 hours and then heated at 50 °C in an argon-filled glove box (< 1 ppm H2O and < 1 ppm O2) for 4 hours. FTIR spectra of the membranes showed no presence of water after this drying procedure (data not shown). All subsequent measurements were made under rigorous control of the environment to exclude water. Complex impedance measurements of the polymer membranes were performed in symmetric cells with stainless steel blocking electrodes under inert atmosphere. Impedance was measured from 106 – 1 Hz with a spacing of six frequency increments per decade. The films were 12.7 mm in diameter with thicknesses in the range of 80 to 150 µm, and a small spring maintained pressure on the cell. Prior to the complex impedance measurements, the cell was slowly heated to the maximum temperature intended for subsequent impedance measurement and then cooled back to the lowest temperature for the impedance measurement. One full thermal cycle required ~10 hours. These thermal cycles were required to relieve stresses in the film and allow for a better contact to both electrodes. Each cell was thermally cycled twice prior to impedance measurement. Attempts to measure the impedance of these cells prior to thermal cycling yielded irreproducible results. The thickness of the membranes was measured before and after thermal cycling with a micrometer. Thickness did not change significantly with thermal cycling, but the thickness values measured after cycling were used to determine the conductivity of the membranes. Glass transition temperature (Tg), melting temperature (Tm), and crystallization temperature (Tc) of each membrane and reference materials were measured using DSC (TA instruments Q2500). Samples were sealed in aluminum DSC pans under an inert atmosphere prior to measurements. The samples were cycled at a rate of 10 °C/min from -90 to 90 °C for 2 cycles. The degree of crystallinity of PEO, Xi, was calculated by integrating the melting endotherms in the DSC thermograms: = where ∆Hm is the specific melting enthalpy of PEO in each sample and ∆Hmo = 197 J/g is the specific enthalpy of melting of PEO crystals of 100% crystallinity.27 Plasticized membranes had two overlapping melting peaks around 40 °C, which were attributed to crystalline PEO and a coordination complex of NaTFS and TEGDME. Peak fitting and deconvolution were necessary to distinguish the contribution of each to the endothermal enthalpy. PEO crystallinity is reported as the weight percent of the entire sample (mPEO,crystalline/msample) and as the weight percent of all PEO in the sample (mPEO,crystalline/mPEO,total) (Table 1). FTIR spectra were collected in attenuated total reflectance (ATR) mode with a diamond crystal. The FTIR instrument (Bruker Alpha) was housed in an argon-filled glove box. For analysis by Raman spectroscopy, membranes were sealed under glass in a special cell to prevent air exposure. Raman spectra were acquired with an Alpha 300 confocal Raman microscope (WITec, 3 ACS Paragon Plus Environment


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GmbH) using a solid-state 532 nm excitation laser, a 20x objective, and a 1800 grooves per millimeter grating. The laser spot size was approximately 1 µm, and the laser power was attenuated to 1 mW to minimize laser-induced heating. Young’s modulus of the membranes was characterized by nanoindentation technique with a target indentation strain rate of 0.2 1/s. An iNano nanoindenter (Nanomechanics, Inc., Oak Ridge, TN), configured with a Berkovich tip, was used to perform indentations. A dynamic oscillation is superimposed upon the semi-static loading in order to measure hardness and Young's modulus as a continuous function of force.28 The membranes for nanoindentation measurement were over 200 µm thick. The maximum indentation depth was 3 µm or 60 °C.43 In addition, they proposed the formation of several metastable complexes with stoichiometric ratios of TEGDME to salt ranging from 1.21.5. The melting temperatures of these phases were lower than the 1:1 compounds, ranging from -5 to 60 °C depending on the stoichiometry and the salt. The melting points of the NaTFSTEGDME complexes in the membranes are below the melting point for the 1:1 complex, which suggests that they are metastable phases with a TEGDME to salt ratio >1. Table 1. Melting (Tm), crystallization (Tc), and glass transition (Tg) temperatures for the different membrane materials measured by DSC. The degree of crystallinity of PEO is also reported in weight percent. Tm (°C)

Sample

TEG DME

NaTFSTEGDME complex

Tc (°C)

PEO

TEG DME

NaTFSTEGDME complex

Tg (°C)

PEO

PEO crystallinity wt.% of sample

wt.% of PEO

PEO

66

46

-56

72%

72%

NaTFS:PEO 1:8

56

34

-30

29%

43%

NaTFS:PEO: TEGDME 1:8:2

-36

31

42

-50

-7

18

-39

19%

53%

NaTFS:PEO: TEGDME 1:8:2 + 5 wt.% CMC

-35

31

44

-49

-15

20

-39

15%

44%

1 m NaTFSTEGDME

-29

15

-49

-40

TEGDME

-28

-36

Table 1 summarizes the melting (Tm), crystallization (Tc), and glass transition (Tg) temperatures for the different membranes and reference materials. The addition of NaTFS salt to PEO lowers Tm of PEO, which is further depressed in the plasticized membranes with TEGDME. The melting point depression confirms the miscibility of PEO, NaTFS, and TEGDME in the molten 9 ACS Paragon Plus Environment


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state.33, 36, 50 Tg of the NaTFS-PEO membrane is higher than pure PEO, because the salt promotes the formation of physical cross-links between the PEO chains.51 Plasticizing the membrane with TEGDME lowers Tg relative to the unplasticized membrane. The increase in segmental motion of the polymer will contribute to higher conductivity for ions that are coordinated by amorphous PEO.33-34, 52 The addition of CMC to the NaTFS-PEO-TEGDME membrane does not significantly change the thermal properties over the measured range (-90 to 90 °C). The measured range is below the Tg and Tm for CMC, which are around 135 and 270 °C, respectively. In addition to lowering Tg, adding TEGDME as a plasticizer changes the fraction of crystalline PEO in the samples (Table 1). The pure PEO used in this study is 72% crystalline. Adding NaTFS salt reduces the amount of crystalline PEO to 43%. The remaining 57% of the PEO in the 1:8 NaTFS-PEO membrane forms the crystalline (PEO)NaTFS compound or contributes to an amorphous form of PEO with dissolved salt. Plasticizing the membranes with TEGDME increases the PEO crystallinity to 53%, but the addition of CMC reduces PEO crystallinity to 44%. By coordinating with the NaTFS salt, TEGDME lowers the salt interaction with the polymer matrix. The decrease in salt interaction could explain the increase in PEO crystallinity in the plasticized membranes without CMC. While TEGDME does not contribute to a reduction in PEO crystallinity, it effectively dilutes the amount of crystalline polymer in the membranes (Table 1). Crystalline PEO is 29 wt.% of the unplasticized membrane and only 19 wt.% of the plasticized membrane. The overall reduction in crystallinity favors higher conductivity since most conduction occurs through connected regions of amorphous electrolyte.33-39

Figure 4. Raman spectra of NaTFS, TEGDME, 1 m NaTFS in TEGDME, PEO, NaTFS-PEO membrane, NaTFS-PEO-TEGDME membrane, and NaTFS-PEO-TEGDME-CMC membrane. (a) Region for coupled CH2 rocking (ρ CH2) and C-O stretching vibrations (ν CO). (b) Region for SO3 symmetric stretch (νs SO3). (c) Region for CF3 in-plane bend (δs CF3). Vibrational spectroscopy provides further insights into the coordination chemistry of the different phases present in the membranes. To understand the solvation of NaTFS and the impact on ionic conductivity, Raman spectra were acquired for membranes with and without TEGDME 10 ACS Paragon Plus Environment

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plasticizer and compared to a 1 m NaTFS solution in TEGDME (Figure 4). Pure TEGDME and PEO have Raman-active modes from 800-860 cm-1 that arise from coupled CH2 rocking and C-O stretching vibrations (Figure 4a).37, 40-41, 44-45, 53-54 The coordination of TEGDME with Na+ gives rise to a new vibrational mode near 867 cm-1, which is assigned to a symmetric Na-On breathing mode (Figure 4a).40-45 This new band is most distinct in the NaTFS-TEGDME solution, but also appears as a shoulder in the PEO membranes plasticized with TEGDME. Coordination of Na+ by PEO does not result in this new band because the structure of the PEO backbone is not distorted by the Na+ cation.45 The plasticized membranes show a new mode at 833 cm-1 that has been assigned to the formation of sodium complexes with glymes at high salt concentrations (≤ 2 moles of tetraglyme per mole of salt).43 The appearance of this mode indicates that most of the Na+ cations are coordinated by TEGDME and not PEO in the plasticized membranes. The 833 cm-1 mode does not appear as a sharp band in the 1 m solution of NaTFS in TEGDME because the salt is too dilute. The addition of CMC to the plasticized membrane has no impact on the Na+ coordination (Figure 4a). Preferential coordination of NaTFS by TEGDME over PEO is supported by studies from Frech and co-workers who compared the (PEO)NaTFS compound to a crystalline complex of NaTFS with TEGDME.46 They noted that packing requirements in the (PEO)NaTFS crystal impose more high energy gauche conformations within the EO units. In contrast, the shorter-chain TEGDME can wrap around the sodium cation in a more energetically favorable configuration. Raman-active modes for the triflate anion are also sensitive to changes in coordination environment (Figure 4b, c). The symmetric stretch of the SO3 group shifts towards higher frequency with increasing ion association (Figure 4b).45-46, 54-55 Based on Raman studies of LiTFS and NaTFS, bands at 1036, 1044, and 1059 cm-1 are assigned to ion pairs, and two different aggregates, respectively.45-46, 54-57 In particular, the νs(SO3) mode at 1059 cm-1 band is a good match for the (PEO)NaTFS crystalline compound that is predicted from the NaTFS-PEO phase diagram.47 In (PEO)NaTFS the TFS anion is highly coordinated as the [Na4TFS]3+ aggregate species.48 Crystallization of the compound depends on the molecular weight of PEO and thermal history of the sample, and we cannot distinguish the crystalline (PEO)NaTFS compound from local structures with similar coordination.45, 57 For simplicity, we attribute this highly aggregated form of TFS to (PEO)NaTFS with the understanding that similar aggregates could form in local structures that mimic the compound. The unplasticized membrane contains NaTFS predominantly in the form of (PEO)NaTFS, in agreement with the phase diagram.47 The Raman spectrum also shows evidence for small amounts of triflate anions with lower degrees of cation coordination. These triflate anions are dissolved in the amorphous PEO polymer responsible for most of the conductivity in the unplasticized membrane. The addition of TEGDME dramatically reduces the amount of (PEO)NaTFS in the membrane. The shift of the SO3 symmetric stretch indicates that the triflate anion forms aggregate species and/or ion pairs with lower degrees of cation coordination (Figure 4b). This reinforces the concept that NaTFS interacts preferentially with TEGDME over PEO, forming coordination complexes with much higher conductivity than (PEO)NaTFS. The average cation coordination number for triflate anions in the plasticized membranes is higher than in the 1 m solution of NaTFS in TEGDME, which shows mainly ion pairs. The band for free anions at 1032 cm-1 is not prominent in the solution or membranes.



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The in-plane bend (scissoring) of the CF3 group in the triflate anion also shifts towards higher frequency with increasing ion association.37, 45-46, 55 Modes at 756, 762, and 769 cm-1 are assigned to ion pairs, aggregates, and the (PEO)NaTFS compound, respectively.45-46, 54-55, 57 Consistent with the analysis of νs(SO3), the addition of TEGDME greatly reduces (PEO)NaTFS in favor of triflate ion pairs and aggregates with lower cation coordination numbers (Figure 4c). The CF3 scissoring mode also confirms that the NaTFS solution in TEGDME has mainly ion pairs. Based on both νs(SO3) and δs(CF3), the addition of CMC to the plasticized membrane does not change the coordination environment of the triflate anions (Figure 4b, c).

Figure 5. FTIR spectra of NaTFS, TEGDME, 1 m NaTFS in TEGDME, PEO, NaTFS-PEO membrane, NaTFS-PEO-TEGDME membrane, and NaTFS-PEO-TEGDME-CMC membrane. (a) Region for SO3 asymmetric stretch (νas SO3) and CF3 symmetric stretch. (b) Region for CF3 in-plane bend (δs CF3). FTIR provides further evidence that the plasticized membranes have anions with lower cation coordination number than the unplasticized membrane (Figure 5). The asymmetric stretch of the SO3 group for free triflate appears as a single stretch around 1270 cm-1. The formation of ion pairs and aggregates splits the νas(SO3) mode into multiple bands.55 The NaTFS solution in TEGDME has two main bands at 1265 and 1291 cm-1 in this region assigned to ion pairs. The PEO membranes have bands at even higher frequency, consistent with the formation of aggregates. The CF3 in-plane bend (scissoring) shows the same characteristic shifts in the IR spectra as in the Raman spectra of the membranes (Figure 5b). The band at 766 cm-1, which is attributed to (PEO)NaTFS, dominates the spectrum of the unplasticized membrane but is largely absent in the plasticized membrane. Consistent with the analysis of the Raman spectra, the addition of CMC to the plasticized membrane does not change the anion coordination. Overall, the vibrational spectra of the triflate anion show that the average cation coordination number in the plasticized membrane is intermediate between a standard liquid electrolyte and the unplasticized membrane, in agreement with their intermediate conductivity and apparent activation energy (Figure 1b).


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The addition of 5 wt.% CMC does not significantly alter the conductivity of the films (Figures 1b) or the ion coordination (Figures 4 and 5). However, CMC addition increase the mechanical strength of the film. The reported Young’s moduli are ~0.3 GPa for pure PEO (M.W. 600,000)58 and ~1.4 GPa for pure CMC.59 The stiffness of composite membranes was measured by nanoindentation (Figure 6). The membrane without TEGDME plasticizer (NaTFS:PEO 1:8) has a Young’s modulus of 0.48±0.04 GPa. The addition of TEGDME (NaTFS:PEO:TEGDME 1:8:2) reduces the modulus to 0.32±0.01 GPa. However, all the decrease in mechanical strength can be recovered by the addition of 5% CMC, which increases the Young’s modulus to 0.47±0.04 GPa. Although this value is relatively low, it is within the range of Young’s moduli typically reported for Nafion™ (0.03 – 3 GPa depending on the level of hydration),60-61 which is the most commonly used membrane in aqueous redox flow batteries and fuel cells. Given the fact that the addition of CMC did not significantly reduce the conductivity of membrane, this strategy is a viable way to reinforce the composite for grid-storage applications.

Figure 6. Young’s modulus measured by nanoindentation as a function of indentation depth. The reported bulk modulus of the film is represented by the measured modulus at a depth of 2,000 nm.



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Figure 7. Overvoltage for sodium plating and stripping in symmetric cells with sodium working and counter electrodes at a current density of 39 µA/cm2. The electrolytes are NaTFS-PEO-TEGDME membranes with and without CMC. The membrane without CMC was 320 µm thick, and the membrane reinforced with CMC was 270 µm thick. The electrochemical performance of the membranes was further evaluated by plating and stripping sodium metal in symmetric cells (Figure 7). The NaTFS-PEO-TEGDME membrane failed after passing only 3.1 C/cm2 of charge, due to the growth of sodium dendrites and subsequent short-circuit. The addition of CMC significantly increased the lifetime of the membrane by more than a factor of four to 14.8 C/cm2. The greater resistance to penetration by sodium dendrites is consistent with the higher modulus measured by nanoindentation (Figure 6). While significant improvements are still needed before these membranes could be used in solid state batteries with sodium metal anodes, the NaTFS-PEO-TEGDME-CMC membrane is an excellent candidate for redox flow batteries. Importantly, the overvoltage did not increase during sodium plating/stripping, which confirms the high stability of this system at extremely negative potentials. A robust membrane with high conductivity has been developed for Na-ion batteries. Standard PEO membranes have poor conductivity near ambient temperature (10-7 – 10-6 S/cm), but the addition of TEGDME as a plasticizer increases the conductivity by two orders of magnitude. Analysis of Raman spectra and DSC thermograms indicates that the increase in conductivity is due to three coactive effects of the plasticizer: (1) TEGDME preferentially solvates the sodium salt to form NaTFS-TEGDME coordination complexes. This prevents the formation of the (PEO)NaTFS compound or similar local structures with highly aggregated ions. (PEO)NaTFS has negligible conductivity but sequesters a large fraction of salt in the unplasticized membranes.


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(2) Plasticizing with TEGDME reduces the overall fraction of crystalline PEO in the membranes and favors the formation of interconnected regions of amorphous polymer and TEGDME with dissolved salt. The amorphous electrolyte forms the pathways for ionic conduction. (3) The addition of TEGDME increases PEO chain mobility in the amorphous polymer, as reflected in the decrease in Tg. While the introduction of TEGDME plasticizer reduces the Young’s modulus of the membrane, the mechanical properties recover by adding a small amount of CMC. The addition of CMC does not significantly change the conductivity or the apparent activation energy for conduction. Importantly, all the materials selected for the membrane (PEO, NaTFS, TEGDME, and CMC) are highly stable at very negative voltages (near Na/Na+), and therefore compatible with sodium metal anodes or strongly reducing redox couples. Such a membrane can enable high voltage, high energy density batteries for grid storage applications. ACKNOWLEDGEMENT This work is supported by Dr. Imre Gyuk, Manager, Energy Storage Program, Office of Electricity Delivery and Reliability, Department of Energy. The spectroscopy effort is supported by the Laboratory Directed Research and Development Program of Oak Ridge National Laboratory, managed by UT-Battelle, LLC, for the U. S. Department of Energy. We thank Xi Chelsea Chen for valuable discussions during the preparation of this manuscript. This manuscript has been authored by UT-Battelle, LLC under Contract No. DE-AC05-00OR22725 with the U.S. Department of Energy. The United States Government retains and the publisher, by accepting the article for publication, acknowledges that the United States Government retains a non-exclusive, paid-up, irrevocable, world-wide license to publish or reproduce the published form of this manuscript, or allow others to do so, for United States Government purposes. The Department of Energy will provide public access to these results of federally sponsored research in accordance with the DOE Public Access Plan (http://energy.gov/downloads/doe-public-access-plan

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ACS Energy Letters

17. Small, L. J.; Pratt, H. D.; Fujimoto, C. H.; Anderson, T. M., Diels Alder Polyphenylene Anion Exchange Membrane for Nonaqueous Redox Flow Batteries. J. Electrochem. Soc. 2016, 163 (1), A5106-A5111. 18. Maurya, S.; Shin, S. H.; Sung, K. W.; Moon, S. H., Anion Exchange Membrane Prepared from Simultaneous Polymerization and Quaternization of 4-Vinyl Pyridine for Non-Aqueous Vanadium Redox Flow Battery Applications. J. Power Sources 2014, 255, 325-334. 19. Kim, D. H.; Seo, S. J.; Lee, M. J.; Park, J. S.; Moon, S. H.; Kang, Y. S.; Choi, Y. W.; Kang, M. S., Pore-Filled Anion-Exchange Membranes for Non-Aqueous Redox Flow Batteries with Dual-Metal-Complex Redox Shuttles. J. Membr. Sci. 2014, 454, 44-50. 20. Wei, X.; Pan, W.; Duan, W.; Hollas, A.; Yang, Z.; Li, B.; Nie, Z.; Liu, J.; Reed, D.; Wang, W.; Sprenkle, V., Materials and Systems for Organic Redox Flow Batteries: Status and Challenges. ACS Energy Lett. 2017, 2 (9), 2187-2204. 21. Jr., D. T. H.; Balsara, N. P., Polymer Electrolytes. Annu. Rev. Mater. Res. 2013, 43 (1), 503-525. 22. Kim, Y. T.; Smotkin, E. S., The Effect of Plasticizers on Transport and Electrochemical Properties of PEO-Based Electrolytes for Lithium Rechargeable Batteries. Solid State Ion. 2002, 149 (1-2), 29-37. 23. Bhide, A.; Hariharan, K., Ionic Transport Studies on (PEO)6:NaPO3 Polymer Electrolyte Plasticized with PEG400. Eur. Polym. J. 2007, 43 (10), 4253-4270. 24. Stephan, A. M., Review on Gel Polymer Electrolytes for Lithium Batteries. Eur. Polym. J. 2006, 42 (1), 21-42. 25. Ruther, R. E.; Sun, C. N.; Holliday, A.; Cheng, S. W.; Delnick, F. M.; Zawodzinski, T. A.; Nanda, J., Stable Electrolyte for High Voltage Electrochemical Double-Layer Capacitors. J. Electrochem. Soc. 2017, 164 (2), A277-A283. 26. Croce, F.; Appetecchi, G. B.; Persi, L.; Scrosati, B., Nanocomposite Polymer Electrolytes for Lithium Batteries. Nature 1998, 394 (6692), 456-458. 27. Buckley, C.; Kovacs, A., Melting Behaviour of Low Molecular Weight Poly (EthyleneOxide) Fractions. In Polymere Aspekte, Springer: 1975; pp 44-52. 28. Oliver, W. C.; Pharr, G. M., An Improved Technique for Determining Hardness and Elastic-Modulus Using Load and Displacement Sensing Indentation Experiments. J. Mater. Res. 1992, 7 (6), 1564-1583. 29. ISO 14577-1: Metallic Materials- Instrumented Indentation Test for Hardness and Materials Parameters. International Organization for Standardization, 2015.



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30. Berthier, C.; Gorecki, W.; Minier, M.; Armand, M. B.; Chabagno, J. M.; Rigaud, P., Microscopic Investigation of Ionic-Conductivity in Alkali-Metal Salts Poly(Ethylene Oxide) Adducts. Solid State Ion. 1983, 11 (1), 91-95. 31. Marzantowicz, M.; Dygas, J. R.; Krok, F.; Nowinski, J. L.; Tomaszewska, A.; Florjanczyk, Z.; Zygadlo-Monikowska, E., Crystalline Phases, Morphology and Conductivity of PEO:LiTFSI Electrolytes in the Eutectic Region. J. Power Sources 2006, 159 (1), 420-430. 32. Irvine, J. T. S.; Sinclair, D. C.; West, A. R., Electroceramics: Characterization by Impedance Spectroscopy. Adv. Mater. 1990, 2 (3), 132-138. 33. Kelly, I. E.; Owen, J. R.; Steele, B. C. H., Poly(Ethylene Oxide) Electrolytes for Operation at Near Room-Temperature. J. Power Sources 1985, 14 (1-3), 13-21. 34. Wang, C. H.; Liu, Q. G.; Cao, Q. J.; Meng, Q. G.; Yang, L. L., Investigation on the Structure and the Conductivity of Plasticized Polymer Electrolytes. Solid State Ion. 1992, 53, 1106-1110. 35. Nicotera, I.; Ranieri, G. A.; Terenzi, M.; Chadwick, A. V.; Webster, M. I., A Study of Stability of Plasticized PEO Electrolytes. Solid State Ion. 2002, 146 (1-2), 143-150. 36. Kelly, I.; Owen, J. R.; Steele, B. C. H., Mixed Polyether Lithium-Ion Conductors. J. Electroanal. Chem. 1984, 168 (1-2), 467-478. 37. Chintapalli, S.; Frech, R., Effect of Plasticizers on Ionic Association and Conductivity in the (PEO)9LiCF3SO3 system. Macromolecules 1996, 29 (10), 3499-3506. 38. Chintapalli, S.; Frech, R., Effect of Plasticizers on High Molecular Weight PEOLiCF3SO3 Complexes. Solid State Ion. 1996, 86-8, 341-346. 39. Frech, R.; Chintapalli, S., Effect of Propylene Carbonate as a Plasticizer in High Molecular Weight PEO-LiCF3SO3 Electrolytes. Solid State Ion. 1996, 85 (1-4), 61-66. 40. Boschin, A.; Johansson, P., Characterization of NaX (X: TFSI, FSI) - PEO Based Solid Polymer Electrolytes for Sodium Batteries. Electrochim. Acta 2015, 175, 124-133. 41. Papke, B. L.; Ratner, M. A.; Shriver, D. F., Vibrational Spectroscopy and Structure of Polymer Electrolytes, Poly(Ethylene Oxide) Complexes of Alkali-Metal Salts. J. Phys. Chem. Solids 1981, 42 (6), 493-500. 42. Mandai, T.; Yoshida, K.; Tsuzuki, S.; Nozawa, R.; Masu, H.; Ueno, K.; Dokko, K.; Watanabe, M., Effect of Ionic Size on Solvate Stability of Glyme-Based Solvate Ionic Liquids. J. Phys. Chem. B 2015, 119 (4), 1523-1534. 43. Mandai, T.; Nozawa, R.; Tsuzuki, S.; Yoshida, K.; Ueno, K.; Dokko, K.; Watanabe, M., Phase Diagrams and Solvate Structures of Binary Mixtures of Glymes and Na Salts. J. Phys. Chem. B 2013, 117 (48), 15072-15085.


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ACS Energy Letters

44. Dong, H. T.; Hyun, J. K.; Rhodes, C. P.; Frech, R.; Wheeler, R. A., Molecular Dynamics Simulations and Vibrational Spectroscopic Studies of Local Structure in Tetraglyme : Sodium Triflate (CH3O(CH2CH2O)4CH3 : NaCF3SO3) Solutions. J. Phys. Chem. B 2002, 106 (18), 48784885. 45. Rhodes, C. P.; Frech, R., Cation-Anion and Cation-Polymer Interactions in (PEO)nNaCF3SO3 (n=1-80). Solid State Ion. 1999, 121 (1-4), 91-99. 46. Rhodes, C. P.; Khan, M.; Frech, R., Crystalline Phases of Poly(Ethylene Oxide) Oligomers and Sodium Triflate: Changes in Coordination and Conformation with Chain Length. J. Phys. Chem. B 2002, 106 (40), 10330-10337. 47. Besner, S.; Vallee, A.; Bouchard, G.; Prudhomme, J., Effect of Anion Polarization on Conductivity Behavior of Poly(Ethylene Oxide) Complexed with Alkali Salts. Macromolecules 1992, 25 (24), 6480-6488. 48. Frech, R.; Rhodes, C. P.; York, S. S., A Comparative Study of Ionic Association in Poly(Ethylene Oxide)-MCF3SO3 Systems (M=Lithium and Sodium). In Solid State Ionics V, Nazri, G. A.; Julien, C.; Rougier, A., Eds. Materials Research Society: Warrendale, 1999; Vol. 548, pp 335-345. 49. Martinvosshage, D.; Chowdari, B. V. R., X-Ray Photoelectron-Spectroscopy Studies on Poly-(Ethylene Oxide) with Sodium Triflate. J. Electrochem. Soc. 1993, 140 (12), 3531-3536. 50.

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