Effect of Plasticization on Ionic Conductivity Enhancement in Relation

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Effect of Plasticization on Ionic Conductivity Enhancement in Relation to Glass Transition Temperature of Crosslinked Polymer Electrolyte Membranes Ruixuan He and Thein Kyu* Department of Polymer Engineering, University of Akron, Akron, Ohio 44325, United States ABSTRACT: The relationship between glass transition (Tg) and ionic conductivity (σ) of an amorphous crosslinked polymer electrolyte membrane (PEM) was examined based on ion−dipole complexation between dissociated lithium cations and ether oxygen of poly(ethylene glycol diacrylate) and plasticization by succinonitrile (SCN). In a binary PEM consisting of a lithium salt/polymer network, Tg increased due to a strong ion−dipole interaction, whereas σ declined due to lower ion mobility coupled to reduced chain mobility. Above the threshold salt concentration of 7 mol %, dual loss tangent peaks were observed in dynamic mechanical studies, which may be ascribed to segmental relaxations of ion−dipole complexed networks and that of polymer chains surrounding the undissociated lithium salt acting like “fillers”. Upon SCN plasticization, these two peaks merged into one that was further suppressed below Tg of the pure network, whereas σ improved to the superionic conductor level. The role of plasticization on the ionic conductivity enhancement is discussed.

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INTRODUCTION Polymer electrolytes for lithium ion batteries have been extensively investigated since the early 1970s.1,2 In general, polymer electrolytes can be classified into gel polymer electrolyte and solid polymer electrolyte (SPE). In gel polymer electrolyte, organic solvents (i.e., ethylene carbonate with dimethyl or diethyl carbonate) not only dissociate lithium ions from their salts but also plasticize polymer networks, thereby affording high ionic conductivity. However, these mixed solvents are notoriously known for volatility and flammability due to low flash points. Consequently, heavy and robust casing is needed to contain these toxic liquid electrolytes during battery assembling. SPEs, being solvent-free and flexible, are promising alternatives that eliminate most if not all of the aforementioned deficiencies without the need for a membrane separator. SPEs are primarily composed of low lattice energy lithium salt (e.g., lithium hexafluorophosphate (LiPF6), lithium bis(trifluoromethanesulfonyl) imide (LiTFSI)) dissolved in high molecular weight polymers such as poly(ethylene oxide) (PEO). By virtue of its low glass transition temperature (Tg) of approximately −60 °C,3−5 PEO has been regarded as a potential candidate because of its rapid segmental motion that facilitates the transport of lithium ions between ion−dipole complexation sites. However, the ionic conductivity of the PEO-based solid polymer electrolytes is rather poor, i.e., in the range of 10−9 to 10−6 S/cm.6 Among various strategies for improving ionic conductivity to realize commercialization, plasticizer addition is the easiest yet most effective one. Superionic conductivity, i.e., >10−3 S/cm, can be achieved by simply adding organic small molecule plasticizers.7 © XXXX American Chemical Society

The correlation between ion transport and polymeric segmental motion in plasticized polymer electrolytes has been extensively studied based on molecular simulations8−10 and also using a variety of spectroscopic characterization techniques,11,12 including Fourier-transform infrared (FT-IR),13 nuclear magnetic resonance (NMR),14 dielectric,15 and impedance13,16 spectroscopy in conjunction with differential scanning calorimetry (DSC) and dynamic mechanical analysis (DMA).17 Suppression of Tg in polymeric networks upon plasticizer addition is a well-known phenomenon,18 but there are a few systematic studies10,13,16 on its relationship with ionic conductivity enhancement in plasticized SPE systems. The present article demonstrates (i) the role of molecular complexation such as ion−dipole interactions in polymer chain relaxation in the crosslinked poly(ethylene glycol) diacrylate (PEGDA)/LiTFSI network and (ii) the effect of SCN plasticizer addition on ionic conductivity enhancement of the crosslinked PEGDA/LiTFSI/succinonitrile (SCN) solid-state polymer electrolyte membrane (PEM). Crosslinked PEGDA networks provide mechanical support as well as conducting medium for lithium ion transport. By virtue of its low lattice energy, LiTFSI can be easily dissociated into Li cations and TFSI anions upon contact with carbonate, ester, or ether groups of polymer electrolytes. Plastic crystalline SCN is chosen here as a solid plasticizer to afford mechanically sturdy, all solid-state PEM as opposed to mechanically fragile “gel electrolyte” that usually contains carbonate derivatives. MoreReceived: May 2, 2016 Revised: July 8, 2016

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DOI: 10.1021/acs.macromol.6b00918 Macromolecules XXXX, XXX, XXX−XXX

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out at a frequency of 1 Hz at an amplitude of 0.2% relative to the initial sample length. FT-IR spectra of the crosslinked membranes of various compositions were collected using an FT-IR spectrometer (Nicolet 380 Thermo Scientific) in attenuated total reflectance (ATR) mode. The acquired spectra were the average of 32 scans with a spectral resolution of 4 cm−1. Ionic conductivity of PEM samples was determined by means of impedance spectroscopy (SI 1260 Impedance/Gain Phase Analyzer and SI 1287 Electrochemical Interface; Solartron Analytical Inc.). A stainless steel (SS)/PEM/SS cell was tested in a frequency range from 1 MHz to 0.1 Hz at an amplitude of 10 mV. A 1 mm thick PEM was sandwiched between the parallel SS electrodes having an area of 10 × 10 mm2. To determine the temperature-dependent ionic conductivity, a homemade heating chamber equipped with a temperature controller (Omron E5AK) was used. Room temperature transference number (t+) of lithium ion was measured on the Li/PEM/Li coin cells on the basis of the potentiostatic polarization method.22,23 Interfacial resistance of the cells before (R0) and after (Rs) polarization were acquired by impedance spectroscopy. The transference number was computed in accordance with the equation t+ = Is(ΔV − I0R0)/Is(ΔV − IsRs), where ΔV is the applied potential difference, and I0 and Is represent the initial and steady state currents of the potentiostatic measurement, respectively.

over, the dielectric constant of SCN is comparable to that of carbonates19 (the dissociation capability should be at least equal to or higher than the mixture of cyclic and acyclic carbonate), and thus, comparable or better ionization (or dissociation) of LiTFSI salt can be anticipated. As demonstrated in our previous paper, completely amorphous PEMs can be fabricated through ultraviolet (UV) crosslinking in the isotropic region of the ternary phase diagrams.20 In this paper, the implication of molecular complexation on the glass transition temperature and ionic conductivity of the unplasticized PEMs will be demonstrated on the basis of DSC, DMA, FT-IR, and impedance spectroscopy. Subsequently, the role of plasticization on ionic conductivity enhancement will be elucidated in relation to the polymer segmental relaxation of the crosslinked PEM network.

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EXPERIMENTAL METHODS

Materials and Sample Preparation. PEGDA with a numberaveraged molecular weight of 700 g/mol, LiTFSI (99.95% purity), and SCN (99%) were purchased from Sigma-Aldrich. Photoinitiator for UV crosslinking is Irgacure 819 (bis(2,4,6-trimethylbenzoyl)-phenylphosphine oxide) obtained from Ciba Co. PEGDA was vacuum-dried at ambient temperature. LiTFSI solid powder was dried at 120 °C prior to use, and SCN was used as received. LiTFSI was soluble in liquid PEGDA at ambient temperature up to 30 wt %. Above 30 wt %, cosolvent of 40/1 (v/v) methyl dichloride/ acetonitrile was needed to ensure complete mixing. Various PEGDA/ LiTFSI membranes were fabricated by curing in a nitrogen-filled glovebox with the aid of the photoinitiator. A detailed procedure of polymer electrolyte membrane (PEM) fabrication by the UV crosslinking procedure can be found in our previous paper.21 For the solvent cast samples containing high LiTFSI concentrations of >30 wt %, the UV-cured membranes were vacuum-dried at 110 °C for at least 24 h to remove any residual solvent. Spacers of different thicknesses were used to prepare film samples for different characterization methods, e.g., the film thickness was varied from 100 μm for FT-IR to 1 mm for DMA studies. Likewise, a binary PEGDA/SCN mixture with an SCN concentration up to 40 wt % and ternary PEGDA/LiTFSI/SCN mixture were prepared in a nitrogen-filled glovebox. Three different ternary mixtures having various SCN amounts were prepared by keeping fixed PEGDA/LiTFSI ratios of 3:1, 1:1, and 1:2 by weight. After vigorously stirring at room temperature, the amorphous melt blend became completely mixed and transparent without requiring any solvent. The photoinitiator amount of 2 wt % relative to the PEGDA weight was added to the isotropic mixtures, and then UV curing was undertaken. Transparent solvent-free solid PEM films were obtained. Sample Characterizations. Differential scanning calorimetry (DSC, TA Q200, TA Instruments, Inc.) was employed to determine the crystal melting temperature (Tm) and glass transition temperature (Tg) of the binary and ternary mixtures under nitrogen gas circulation. Samples weighing the recommended amount of approximately 10 mg were placed in aluminum pans and hermitically sealed. In the case of PEGDA/LiTFSI liquid mixtures, the samples were cooled to −90 °C, equilibrated for 10 min, and then ramped up to 40 °C at a heating rate of 3 °C/min. For the cured PEGDA/LiTFSI, PEGDA/SCN binary membranes, and PEGDA/LiTFSI/SCN ternary PEMs, DSC scans were conducted from −90 to 80 °C at the same heating rate unless indicated otherwise. Dynamic mechanical analysis (DMA, TA Q800, TA Instruments, Inc.) was conducted on the UV-cured samples from −65 to 80 °C at 3 °C/min after keeping isothermally at −65 °C for 10 min. For the DMA measurement, samples were fabricated by UV curing of the starting liquid mixtures in a spacer with dimensions of 20.0 mm in length × 7.0 mm in width × 1.0 mm in thickness. The rectangular cured strip was stretched for a small static strain by applying a preload force of 0.01 N, and the dynamic oscillatory measurement was carried

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RESULTS AND DISCUSSION 1. Molecular Complexation in Binary PEGDA/LiTFSI Mixtures. 1.1. DSC Investigation of PEGDA/LiTFSI Mixtures before and after Crosslinking. In panels a and b of Figure 1 are shown the DSC thermograms of binary PEGDA/LiTFSI mixtures before and after photo-crosslinking, respectively. As shown in Figure 1a, the melting temperature (Tm) of neat

Figure 1. Variations of DSC thermograms of PEGDA/LiTFSI mixtures as a function of LiTFSI salt concentration ranging from 0 to 67 wt % (33/67 equiv to 28 mol %) (a) before and (b) after photocrosslinking. B


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increases by approximately 40 °C upon chemical crosslinking, but its extent in regime I is less clear due to the presence of the PEGDA crystal. The observed Tg increment of ∼40 °C can be unambiguously assigned to the restricted chain motions caused by the chemical junctions. As manifested by the parallel slopes in regime II, the composition-dependent trend of Tg in the crosslinked PEM is seemingly governed by the strong ion− dipole coupling. The combination of both transient “physical” (ion−dipole complexation) and chemical crosslinking synergistically suppresses the segmental relaxation of polymer chains in the crosslinked PEM system. A similar observation was made by Lascaud et al. in a supercooled linear-PEO/LiTFSI blend, namely, Tg increases linearly with increasing salt concentration up to a threshold concentration of 11 mol % of LiTFSI, beyond which the slope becomes smaller.25 Because segmental motion of polymer chains can exert a strong influence on ion transport, it is interesting to probe the effect of lithium salt concentration on ionic conductivity of PEMs. 1.2. Ionic Conductivity of Crosslinked PEGDA/LiTFSI PEMs. Complex plane impedance plots of PEMs with lithium salt concentration up to 40 wt % are depicted in Figure 3a. Low frequency intersection of the semicircle to the real axis signifies bulk resistance of PEM, from which ionic conductivity can be evaluated. In Figure 3b is shown the decay of the room temperature ionic conductivity with lithium salt concentration of the crosslinked PEMs, which varies from 10−6 to 10−9 S/cm with increasing salt concentration, which is far below the

PEGDA is located at 13 °C. With the addition of LiTFSI salt, Tm is depressed and eventually vanishes at concentrations above 30 wt % of lithium salt, indicating the amorphous nature of these mixtures. The melting point depression behavior was already discussed in our previous paper20 that showed the eutectic phase diagram, and thus, it will not be elaborated here. Of particular interest is that a single glass transition temperature (Tg) can be seen clearly in all intermediate concentrations, which systematically shifts from −59 °C at 30 wt % (i.e., 7 mol %) to −21 °C at 67 wt % (or 28 mol %) with increasing salt concentration. The elevation of the Tg implies the formation of the molecularly complexed network due to the strong ion− dipole interaction between the dissociated lithium cations of the salt and ether oxygen of the PEGDA network chains. However, for mixtures below 30 wt % of salt concentration, Tg remains virtually invariant, which may be attributed to the presence of PEGDA crystals. Of particular importance is that adding salt not only elevates the Tg of the PEGDA/LiTFSI mixtures but also broadens the glass transition zone. The broadening of the DSC-Tg zone is suggestive of various chain dynamics resulting from the network heterogeneity, implying that multiple relaxations might be taking place, e.g., the segmental relaxation of polymer chains in the immediate vicinity of the ion−dipole complexed site may behave differently from that in the bulk network. Unlike the liquid mixture of PEGDA/LiTFSI, the crystal melting peak is no longer discernible in the crosslinked PEM system (Figure 1b), where only a single Tg is evident that increases from −41 °C at 0 wt % to 14 °C at 67 wt % with increasing lithium salt concentration. The lack of crystal melting in combination with the systematic movement of the single Tg indicates completely amorphous character of the crosslinked PEMs. As shown in Figure 1b, the temperature range corresponding to the Tg curve (hereafter called Tg transition zone) becomes broader at approximately 10 °C, especially at the high salt concentrations relative to the low salt concentrations. The observed trend of broad Tg elevation due to ion−dipole complexation is consistent with that reported by Seki et al.24 for their crosslinked P(EO/PO)/LiTFSI system. In Figure 2, the DSC-Tg is further plotted against lithium salt concentration for a better comparison of the trends of the uncrosslinked versus crosslinked systems in which two regimes were identified for the present UV-crosslinked PEM. It can be noticed in regime II that Tg of all uncrosslinked compositions

Figure 3. (a) Complex plane impedance plots of various crosslinked PEGDA/LiTFSI PEMs and (b) room temperature ionic conductivity of PEMs with various lithium salt concentrations ranging from 10 (90/ 10) to 67 wt %.

Figure 2. DSC-Tg of PEGDA/LiTFSI liquid mixtures (square) and crosslinked PEMs (circle) at various LiTFSI concentrations ranging from 0 to 67 wt % (28 mol %). C


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Macromolecules targeted value of 10−3 S/cm of the superionic conductor range. Possible reasons for the rapid decline of ionic conductivity with increasing salt concentration will be further discussed in section 1.3. Regardless, the observed low value of ionic conductivity suggests the need for accelerating the polymer chain mobility, which may be achieved through plasticization. We are pleasantly surprised that two regimes having two different slopes can be clearly identified below and above the threshold concentration of 30 wt % (or 7 mol %), which coincides with that determined from the DSC-Tg in Figure 2. In regime I, the ionic conductivity decreases with increasing salt concentration, whereas it drops much faster in regime II with increasing salt loading. In these PEGDA/LiTFSI binary PEMs, the opposite trend between Tg versus salt concentration and ionic conductivity versus salt concentration indicates their reciprocal relationship, implying strong coupling between lithium ionic transport and the segmental motion of polymeric network chains. According to the ab initio calculations by Johansson et al.,8 one lithium ion can form four to six coordination bonds with polyether oxygens simultaneously, which may occur either within a single PEG chain (intrachain) or multiple neighboring chains (interchains). The molecular complexation via ion− dipole interactions presumably serves as a transient “physical” crosslinking in the present polymer electrolyte network, wherein lithium ion can move from one complexed site to another through conformational “trans−gauche” transformation. As manifested by the elevation of Tg, the transient physical ion−dipole complexation sites behave like the chemical junctions (i.e., crosslinked points), whereby the average chain lengths between the adjacent chemical junctions become shorter and less mobile due to the restricted segmental motions imposed by the ion−dipole interaction. The reduced ionic conductivity in both regimes is undoubtedly coupled to the reduced chain mobility as manifested by the increasing trend of Tg with salt concentration. 1.3. Dynamic Mechanical Analysis of Crosslinked PEGDA/ LiTFSI PEMs. In support of the DSC-Tg observation, DMA was performed to probe the mechanical relaxations of PEMs. It should be pointed out that DMA is rarely employed in the characterization of polymer “gel” electrolytes because of their inherently weak mechanical strength, whereas PEO-based thermoplastic PEMs are too fragile. In panels a and b in Figure 4 are shown the variations of storage modulus and tan δ curves of the crosslinked PEGDA/LiTFSI PEMs as a function of LiTFSI concentration ranging from 0 to 67 wt %. As depicted in Figure 4a, the storage modulus for membranes in the glassy state appears more or less comparable for all lithium salt concentrations. With increasing temperature, the storage moduli of the PEM in the low salt composition drops approximately 2 orders of magnitude, which is characteristic of the glass transition. A systematic movement of Tg can be seen with salt up to 30 wt %, beyond which there appears dual drops in the modulus, which is suggestive of additional relaxation. In the DMA experiments, it is customary to analyze Tg behavior based on the variation of tan δ peaks with compositions as depicted in Figure 4b. The Tg of the neat PEGDA network is located at −16 °C, which increases to 4 °C at 30 wt % of LiTFSI loading. Consistent with the DSC observation, the trend of Tg elevation with the addition of lithium salt may be attributed to lithium cation−ether oxygen complexation. Of particular interest is that two relaxation peaks

Figure 4. (a) Storage modulus and (b) tan δ vs temperature of crosslinked PEGDA/LiTFSI PEM with the LiTFSI fraction ranging from 0 (100/0) to 67 wt % (33/67 equiv to 28 mol %) obtained by dynamic mechanical temperature sweep measurement at a frequency of 1 Hz.

corresponding to different relaxation processes may be operating at the 40 wt % salt (or 11 mol %) and higher concentrations. For instance, in the 60/40 PEGDA/LiTFSI PEM, a shoulder peak appears at 9 °C on the lower temperature side of the main tan δ peak that is located at 15 °C. Upon further increasing the lithium salt concentration, the two peaks gradually separate further from each other. The high temperature peak continues to elevate, and the low temperature (shoulder) peak tends to level off (suggestive of saturation of ion−dipole complexes) with a further increase in the lithium salt concentrations. The occurrence of dual relaxations is clearly discernible in the DMA experiment, but the low temperature Tg can hardly be noticed in the DSC thermograms due to the broad nature of the Tg zone in high salt concentrations, reflecting different, albeit complementary, sensitivities between the two techniques. For the purpose of completeness, Tg as determined by tan δ peaks from DMA is plotted against the LiTFSI fraction in comparison with those of DSC-Tg and ionic conductivity in Figure 5. Note that the Tg values as determined from DMA and DSC are different due to the well-known frequency effect, e.g., tan δ peak (or DMA-Tg at −16 °C) of the neat PEGDA network is approximately 25 °C higher than that at −41 °C from DSC. However, the dual slopes can be observed in both DMA and DSC, except that the DMA-Tg further splits into two slopes in regime II; the higher slope exhibits strong dependence on salt concentration and the other shows little or no dependency on salt concentration, suggesting the occurrence of an additional relaxation mechanism. It should be emphasized that the higher slope of DMA-Tg in regime II is parallel to that D


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scattering (SAXS, data not shown) in the present unplasticized crosslinked PEM system. The change of slope from regime I to regime II may be attributed to the change in the mechanism of segmental motion of the network chains contributed by the ion−dipole complexation to that of the PEG chains absorbed on the undissociated salt surface. Recognizing how complexation of dissociated lithium cations and undissociated lithium salts with polyether chains can profoundly affect the polymer chain motions in the crosslinked PEMs, the declining trends of ionic conductivity in Figure 5 (and also Figure 3b) may be interpreted in terms of both the number of dissociated lithium ions and their mobility. In regime I, the reduction of ionic conductivity with LiTFSI may be attributed to the complexation between the lithium cations and polyether oxygen driven by the ion−dipole interaction that lowers the mobility of ions. Whereas in regime II, the accelerated reduction is primarily contributed by the hightemperature relaxation of complexed PEG chains on the undissociated salt surface. As mentioned earlier, increasing lithium salt is analogous to reducing polyether chains, and thus, there is no sufficient ether oxygen available above this threshold salt concentration to further dissociate the lithium cations from the salt. Consequently, in regime II, the number of dissociated lithium cations as well as the ion mobility is expected to decline, although the amount of undissociated lithium salt will increase with a further increase in the salt concentration that eventually acts like “fillers”. Hence, the observed rapid decline in the ionic conductivity with increasing salt concentration in regime II may be produced by the reduction in the number of dissociated lithium ions as well as the ion mobility coupled to the restricted segmental chain motions at the undissociated salt surface. 1.4. FT-IR Investigation of Crosslinked PEGDA/LiTFSI PEMs. For the lithium ion−polyether oxygen complexations to be verified, the FT-IR spectra of a binary PEGDA/LiTFSI PEM are depicted in Figure 6 corresponding to three wavelength regions. The IR bands that correspond to the functional groups on PEGDA, including nonresolved CH2 symmetric and asymmetric stretching (νCH2),12 overlap of C−O−C symmetric and asymmetric stretching (νCOC),28 and CH2 rocking (ρCH2),13,29,30 can be observed in the wavelength ranges of 3000−2800, 1110−1080, and 1000−820 cm−1, respectively. As shown in Figure 6a, the CH2 stretching band located at 2864 cm−1 shifts to a higher wavenumber with increasing LiTFSI salt concentration. At 67 wt % of LiTFSI, a drastic shift of 25 cm−1 can be noticed. As salt concentration increases, a new shoulder peak at 2915 cm−1 appears and concurrently moves to a higher wavenumber. This observation is consistent with an earlier FTIR study on the PEO/LiTFSI system,12 indicating the formation of lithium ion−polyether oxygen complexation. The formation of ion−dipole complexation can be also confirmed via the shift of other bands in Figure 6b. The C− O−C band at 1091 cm−1 of the neat PEGDA moves to a lower wavenumber for 10 cm−1 at the 67 wt % LiTFSI loading, suggesting the occurrence of an ion−dipole interaction between lithium ion and ether oxygen.31 Transport of lithium ion is generally facilitated via the conformational transition of polyether segments from trans to gauche and vice versa; thus, it is of paramount importance to examine the peaks corresponding to CH2 rocking vibrations that are sensitive to different conformations. Figure 6c exhibits a peak at 848 cm−1 corresponding to the trans conformation of the O−CH2−CH2−O bond (please see the molecular structure depicted at the top inset of Figure 6c) and a peak at 944 cm−1

Figure 5. Tg of crosslinked PEGDA/LiTFSI PEMs determined by the tan δ peaks (triangle) in comparison with the DSC- Tg (circle) and room temperature ionic conductivity (star) of PEMs as a function of the LiTFSI salt concentration.

obtained in DSC, although the minor (or almost invariant) slope was not discernible in the DSC experiment due to the broad nature of the Tg zone. Moreover, the threshold concentration of DMA-Tg for regimes I and II coincides with the change of slope of room temperature ionic conductivity versus salt concentration plot of the PEM, suggesting a strong correlation between the ion transport and the segmental chain mobility of the PEM network. It should be noted that increasing LiTFSI concentration is equivalent to reducing the concentration of the ether oxygen. One may conjecture that, up to the threshold concentration, all ether oxygens are fully capable of complexing with the dissociated lithium cations from the LiTFSI salt. At the threshold concentration, all dissociated lithium cations are presumably consumed in the complexation with the available ether oxygens, thereby resulting in the stationary Tg upon further increasing LiTFSI. It may be hypothesized that, beyond the threshold salt concentration, there will be an insufficient amount of ether oxygen, and thus, there will be some lithium salt left undissociated. The undissociated salt amount is expected to increase with further addition of LiTFSI, which could thus form salt aggregates that probably act like “fillers”. It may be envisaged that polymer chains adjacent to these salt aggregates may be absorbed on their surfaces, and thus, their segmental chain motions may be restricted, resulting in further increase of the Tg as manifested in regime II (Figure 5). In ion-containing polymers or ionomers, ion clusters (or ion aggregates) are known to be microphase separated,26,27 exhibiting dual Tgs attributable to the backbone chain motion and ionic cluster domains with the latter (Tg of ionic cluster) being higher than the former (Tg of backbone chains). In the present DMA investigation, splitting of the tan δ peak into two peaks was found in regime II above the threshold concentration (Figure 5), and the two peaks further show different dependences on LiTFSI concentration. Therefore, it may be hypothesized that, in regime II, the seemingly invariant lowtemperature relaxation may be ascribed to segmental relaxation of chains in the bulk network. On the other hand, the hightemperature relaxation peak may be attributed to segmental relaxation of the PEG chains surrounding the undissociated “filler-like” salts that occur in regime II. These assignments are consistent with the mechanical relaxations of ion-containing polymers or ionomers, except that there are no identifiable microphase-separated cluster domains by small-angle X-ray E


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gradually to a higher wavenumber of 950 cm−1 at the LiTFSI concentration of 67 wt %. This lithium salt concentration range in Figure 6d is consistent with that of regime II, which was previously obtained by both Tg and ionic conductivity measurements. Because intensity of both trans and gauche peaks have decreased, it is reasonable to infer that, for PEM having a concentration higher than 40 wt %, the amount of available ether oxygen is too low to dissolve all the lithium salts. Hence, some excess salts will be left undissociated at such high salt concentrations, which in turn would act like “fillers”. Moreover, the shift of the gauche peak at 944 cm−1 toward a higher wavenumber of 950 cm −1 with increasing salt concentration implies that some undissociated salts may be interacting with lithium cations of the complexed sites, thereby weakening the complexation. Although the observed correlation between the Tg of the PEM network and ionic conductivity is noteworthy, the extremely low ionic conductivity of these PEMs makes it impractical to be applied to any solid-state batteries. For a higher ionic conductivity to be achieved, it is of paramount importance to adopt a strategy of adding solid plasticizer such as SCN to efficiently ionize the lithium salt and concurrently plasticize the PEM networks. 2. Effect of Plasticization on the Glass Transition of the Crosslinked PEGDA Network. Prior to discussing the plasticization effect of SCN on molecularly complexed, crosslinked PEGDA/LiTFSI network, it is essential to investigate the influence of SCN addition on the crosslinked neat PEGDA membrane. As shown in Figure 7a and b, the variation of Tg with SCN loading can be captured in both DSC thermograms and DMA tan δ curves of the UV-crosslinked

Figure 6. FT-IR spectra for crosslinked PEGDA/LiTFSI PEM as a function of LiTFSI salt concentration ranging from 0 (100/0) to 67 wt % (33/67 equiv to 28 mol %): (a) CH2 symmetric and asymmetric stretching at 2864 cm−1, (b) C−O−C symmetric and asymmetric stretching at 1091 cm−1, (c) CH2 rocking of both trans (848 cm−1) and gauche (944 cm−1) modes of the O−CH2−CH2−O bond at 0−40 wt % salt concentrations and (d) 40−67 wt % salt concentrations.

representing the gauche conformation (please see the molecular structure depicted at the bottom inset of Figure 6c). With increasing LiTFSI concentration from 0 to 40 wt %, the intensity of the 848 cm−1 peak becomes smaller. However, the intensity of the gauche peak of the O−CH2−CH2−O bond keeps increasing despite the fact that the amount of PEG has reduced in the PEM, indicating the occurrence of conformational transformation from trans to gauche of the O−CH2− CH2−O bond with increasing lithium salt concentrations up to 40 wt %. In the gauche conformation, two adjacent oxygen atoms are pulled closer to each other due to ion−dipole complexation, and thus there is higher energy cost in the gauche conformation. Hence, the gauche state may be relatively unstable, which tends to revert back to the trans state by releasing the coordination-bonded lithium cations. It may be envisaged that Li ion transport may be further influenced by the segmental mobility of the PEG chains. In Figure 6d are shown the CH2 rocking peaks of the PEMs having lithium salt concentration from 40 to 67 wt % in which the trans peak at 848 cm−1 continues to decline due to the lesser amount of PEG. On the other hand, the gauche peak at 944 cm−1 shows the reverse trend of decreasing peak intensity with increasing salt concentration, while the peak shifts

Figure 7. (a) DSC thermograms and (b) tan δ vs temperature for various UV-crosslinked PEGDA/SCN membranes with SCN concentration varying from 0 to 40 wt % (i.e., 60/40). F


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Figure 8. DSC thermograms of three plasticized PEM systems. Each system has an individual fixed PEGDA/LiTFSI weight ratio (e.g., 3:1, 1:1, and 1:2) yet the SCN amount is varied from 0 to 40 wt % (i.e., 60/40).

binary PEGDA/SCN membranes. As can be seen in the DSC thermograms (Figure 7a), only a single glass transition is evident that moves systematically from −40 °C of neat PEGDA (0% SCN) to −47 °C at 30% SCN loading. The Tg reduction with SCN loading implies faster polymer chain relaxation afforded by effective plasticization of SCN. According to the binary phase diagram reported previously,20 at ambient temperature the PEGDA/SCN mixtures having SCN up to 53 wt % are completely amorphous before photopolymerization. Although most of the membranes remained amorphous after crosslinking, in our recent paper it was found that polymerization-induced phase separation (PIPS) and polymerization-induced crystallization (PIC) can occur at high SCN concentration in close proximity to the coexistence line of isotropic + plastic crystal,21 including the present 60/40 PEGDA/SCN mixture (Figure 7a). Consequently, two endothermic peaks corresponding to SCN crystalplastic crystal transition at −40 °C and SCN plastic crystal melting temperature at 58 °C can be discerned in the DSC thermogram of 60/40 PEGDA/SCN PEM while the Tg continues to decline to −50 °C. These phase transitions can be discerned in the variation of tan δ with temperature plot. In Figure 7b, the tan δ peak at −25 °C represents the crystalplastic crystal transition; a broader loss tangent peak at a higher temperature is associated with the segmental chain motions within the plastic crystal phase that melts away at 66 °C. 3. Effect of Plasticizer on the Complexed, Crosslinked PEGDA/LiTFSI Network. 3.1. DSC Investigation of Crosslinked PEGDA/LiTFSI/SCN PEMs. For the influence of SCN plasticization on Tg of complexed PEMs to be illustrated, Figure 8 shows DSC thermograms of three PEGDA/LiTFSI systems having various LiTFSI amounts of 25, 50, and 67 wt % that correspond to PEGDA/LiTFSI weight ratios of 3:1, 1:1, and 1:2. In each system, the PEGDA/LiTFSI weight ratio was kept constant such that the only variable is SCN concentration in the ternary crosslinked PEMs ranging from 0 to 40 wt % of SCN at an increment of 10 wt %. In Figure 8a, the (3:1 PEGDA/LiTFSI)/SCN system reveals one single Tg that shifts from −25 to −47 °C as SCN concentration increases from 0 to 20 wt %. The Tg continues to decline to −57 °C with further addition of SCN to 40 wt %, whereas two endotherms can be discerned clearly in the thermograms corresponding to SCN crystal-plastic crystal transition at −40 °C and SCN plastic crystal melting at 50 °C. As reported previously by us, for those PEMs close to the coexistence envelop of the isotropic + plastic crystalline region, polymerization-induced crystallization can take place during the PEM fabrication via photo-crosslinking.21

As shown in Figure 8b and c, a similar declining trend of the single Tg can be noticed in the other two PEM systems. Of particular interest is that higher LiTFSI concentration in the plasticizer-free PEGDA/LiTFSI PEM results in larger reduction of Tg at the same SCN loading. Likewise, the crystal-plastic crystal phase transition of SCN and plastic crystal melting transition can be observed. With increasing LiTFSI, the crystalplastic transition remains virtually invariant, but at 40 wt % of SCN, the plastic crystal melting temperature moves to 22 °C at a PEGDA/LiTFSI weight ratio of 1:1 (Figure 8b) and −5 °C at a PEGDA/LiTFSI weight ratio of 1:2 (Figure 8c). The observed trends of the systematic Tg shift and melting point depression of the SCN plastic crystal may be attributed to the plasticization effect, which is commonly observed in crystalline polymer blends. The stationary crystal-plastic crystal phase transition at all PEGDA/LiTFSI ratios implies that SCN crystals are no longer soluble in the PEM network. 3.2. Ionic Conductivity of Crosslinked PEGDA/LiTFSI/SCN PEMs. In the examination of the relationship between Tg of the PEM network and ionic conductivity, the observed DSC-Tg of various PEMs and the corresponding room-temperature ionic conductivities are plotted in Figure 9a and b, respectively. It can be clearly seen in Figure 9a that Tg of the SCN-plasticized, complexed PEM is even lower than that of the plasticized, uncomplexed PEGDA/SCN network. That is to say, in the (1:2 PEGDA/LiTFSI)/SCN PEM system, Tg decreases from 14 °C at 0 wt % to −83 °C at 40 wt % with increasing SCN concentration as compared to only a minor drop of 10 °C in the case of the crosslinked PEGDA/SCN membrane. The drastic reduction of Tg may be attributed not only to plasticization by SCN, but also additional plasticization afforded by the further dissociation of lithium ions from their salts by the nitrile group of SCN. In Figure 9b, on the other hand, is shown drastic improvement of ionic conductivity for several orders of magnitude with increasing SCN concentration in all three systems, whereas ionic conductivity tended to level off due to the existence of a plastic crystal phase caused by PIC at higher SCN concentrations. It should be noted that, for the third PEM system having the lowest amount of the PEGDA matrix, i.e., (1:2 PEGDA/LiTFSI)/SCN, ionic conductivity even reaches the superionic conductivity level (10−3 S/cm) at 40 wt % SCN loading relative to ∼10−9 S/cm in the unplasticized PEM. This superionic conductive PEM was tested in Li-ion half-cells, which showed high capacity and satisfactory retention of ∼80% after 50 cycles.21 G


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further plasticization beyond that afforded by −CN, leading to enhanced chain mobility as manifested by the dramatic Tg reduction in Figure 9a. In practice, lithium ion batteries are expected to withstand high temperature applications. Hence, it is worth evaluating the temperature dependence of ionic conductivity of PEMs of the plasticized PEMs. As shown in Figure 10, for all three PEM systems investigated in a temperature range from 20 to 100 °C, the ionic conductivity increases with temperature. However, the increase of ionic conductivity was found to be larger (i.e., 3−4 orders of magnitude) for the unplasticized PEM and smaller (i.e., approximately 1 order of magnitude) for the highly plasticized PEM. It should be noted that having a plastic crystal melting point at −5 °C, the 60/40 (1:2 PEGDA/LiTFSI)/SCN (corresponding to 20/40/40 PEGDA/LiTFSI/SCN) is amorphous within the investigated temperature range, and its ionic conductivity increases from the superionic value of 1.0 × 10−3 S/cm at 26 °C to 5.8 × 10−3 S/cm at 91 °C. In most amorphous PEMs, the ionic conductivity versus reciprocal absolute temperature plots exhibit nonlinear curvature trends implying the departure from the Arrhenius type. However, a sudden slope change occurs in some PEMs having high SCN loadings (e.g., see (1:1 PEGDA/LiTFSI)/SCN with 40 wt % of SCN in Figure 10b) attributable to the melting of SCN plastic crystals that formed via PIC during photo-crosslinking. Next, the Vogel−Tammann−Fulcher (VTF) empirical equation was employed to analyze the relationship between ionic conductivity and glass transition temperature in plasticized PEMs because the VTF equation is capable of elucidating temperature-dependent trends of the ionic conductivities of various polymer electrolyte systems.24,33 However, at high salt concentrations where ion pairing occurs, the VTF equation may not be operative because ionic concentration also varies significantly with temperature.34 The original VTF equation, describing temperature-dependent viscosity based on the free volume theory,35,36 may be extended to ionic conductivity of a dilute electrolyte system, i.e.,

Figure 9. (a) DSC-Tg and (b) room temperature ionic conductivity of various plasticized PEM systems (PEGDA/LiTFSI weight ratios are fixed at 3:1, 1:1, and 1:2) and the plasticized neat PEGDA network as a function of SCN concentration.

The reverse trend can be noticed in Figure 9b in comparison to that in Figure 9a, implying a reciprocal relationship between ionic conductivity and Tg upon SCN plasticization. It is noteworthy that greater improvement of ionic conductivity was also achieved for the system with higher lithium salt contents or lower polymer precursor amounts. Recalling that nitrile (CN)32 has a higher dielectric constant than that of ether oxygen (C− O−C),16 it is plausible that in the plasticized PEM SCN may not only plasticize polymer segments but also dissociate lithium cations from the lithium salt as well as from the lithium ion− polyether complexes. Hence, the number of mobile lithium ions is anticipated to rise after the additional dissociation afforded by the −CN, thereby contributing to a higher ionic conductivity. On the same token, dissociated TFSI bulky anions could offer

σ = σ0e−B / R(T − T0)

(1)

The same expression can be also deduced on the basis of the configurational entropy model.37−39 In eq 1, σ is ionic conductivity. T0 represents the reference temperature, which is approximately 40−50 °C lower than Tg. A = σ0 is the empirical prefactor. B is the empirical exponent, representing apparent activation energy, i.e., B = RB0 normalized by the gas

Figure 10. Ionic conductivity versus reciprocal temperature for three different plasticized PEM systems (PEGDA/LiTFSI weight ratios of 3:1, 1:1, and 1:2) with SCN concentration up to 40 wt % of SCN (i.e., 60/40). H


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Macromolecules constant (R) in eq 1. Note that B0 is generally expressed in kelvin. In the present case, T0 is taken as 50 °C lower than DSC-Tg. Because the original VTF equation on viscosity is only valid for amorphous polymers, only those data belonging to amorphous PEMs are analyzed as tabulated in Table 1.

at ambient temperature; thus, it is of particular interest to evaluate the influence of SCN concentration on transference number of PEMs having a PEGDA/LiTFSI ratio of 1:2; which is shown in Table 2. Table 2. Room Temperature Ionic Conductivity (σRT) and Transference Number (tLi+) of Lithium Ion in (1:2 PEGDA/ LiTFSI)/SCN PEM for Various SCN Loadings

Table 1. DSC-Tg, Room Temperature Ionic Conductivity (σRT), and VTF Parameter B for Various Amorphous Plasticized PEMs with different PEGDA/LiTFSI Weight Ratiosa PEGDA/ LiTFSI

SCN (wt %)

Tg (°C)

3:1

10 20 10 20 30 10 20 30 40

−36.0 −47.3 −38.6 −55.5 −71.1 −31.2 −51.5 −78.4 −84.1

1:1

1:2

σRT (S/cm) 1.3 6.0 1.1 1.1 4.0 2.4 5.1 3.4 1.0

× × × × × × × × ×

10−5 10−5 10−5 10−4 10−4 10−6 10−5 10−4 10−3

B (kJ/mol)

R

8.6 8.1 9.1 9.0 8.9 9.8 6.7 4.7 3.8

0.9995 0.9990 0.9993 0.9998 0.9981 0.9996 0.9961 0.9995 0.9996

PEGDA/LiTFSI 1:2

2

SCN (wt %) 20 30 40

σRT (S/cm) −5

5.1 × 10 3.4 × 10−4 1.0 × 10−3

tLi+ 0.26 0.38 0.69

In typical carbonate-containing liquid electrolyte systems, the transference number is approximately 0.3−0.4.40 In the present solid PEMs in which all ether oxygens in the polymer network are used up in the ion−dipole complexation before plasticization, the transference number would be low. With increasing SCN concentration, it increases from 0.26 at 20 wt % of SCN to 0.70 at 40 wt % of SCN, suggesting dominant lithium cation transport over anion transport. Contrary to the aforementioned simulation results, this enhanced cation transport may be ascribed to the dissociation of the lithium ions from both the lithium ion−ether oxygen complexes and the residual undissociated lithium salts, which are further afforded by the nitrile group of the SCN plasticizer. Therefore, more lithium cations are freed up from the complexation sites, which enable them to contribute to the enhanced ionic conductivity while some bulky TFSI anions may be trapped in the crosslinked network. 3.3. FT-IR Results of Crosslinked PEGDA/LiTFSI/SCN PEMs. For the interaction between SCN and the crosslinked PEGDA/ LiTFSI PEM network to be examined, the FT-IR spectra of plasticized PEMs were acquired as depicted in Figure 11. In an effort to analyze various interactions such as lithium ion− polyether oxygen complexation, nitrile-ethylene glycol hydrogen bonding,41 and the dissociation capability of SCN, the IR bands belonging to PEGDA in conjunction with those of the TFSI anion may be probed. However, the IR bands associated with PEGDA are rather complex in comparison with the latter TFSI, which is mainly contributed by ionic interaction. Figure 11 represents IR spectra of three PEM systems having different PEGDA/LiTFSI weight ratios. The broad peak at 1080 cm−1 is the overlap of the symmetric and asymmetric C− O−C stretching (νCOC) of PEGDA, which is evident in the high PEGDA contents, but the peak position remains unchanged with increasing SCN. However, the C−O−C stretching band of PEGDA diminishes (see Figure 11b) and eventually disappears with a further decrease of PEDGA (see Figure 11 c). The IR peaks at 1180, 1130, and 1050 cm−1 corresponding to CF3 asymmetric stretching (νaCF3), SO2 symmetric stretching (νsSO2), and S−N−S asymmetric stretching (νaSNS) of TFSI anion, respectively,42,43 belong to the TFSI anion. These peaks shift to a higher wavenumber for approximately 2−5 cm−1 with SCN addition, indicating that TFSI anions are dissociated from their counter lithium cations as a result of the −CN of SCN competing with the TFSI anion for the lithium ions. It may be inferred that SCN is capable of dissociating not only the lithium ion from its salt but also from the lithium ion−ether oxygen complexes of the crosslinked PEGDA/LiTFSI. As a consequence, more mobile lithium ions would be available to contribute to higher ionic conductivity.

The upper limit of temperature in the VTF test was set to 100 °C above DSC-Tg.

a

Parameter B and the coefficient of determination (i.e., R2) obtained by VTF analysis, in conjunction with DSC-Tg and room temperature ionic conductivity of various plasticized PEMs, are compared in Table 1. It was mentioned earlier that reduction of Tg and improvement of σ with SCN addition were seen in all three systems. In the VTF analysis, B also decreases with increasing SCN loading, implying the apparent energy barrier of ionic transport is lowered due to SCN plasticization. Among the three systems, both B and Tg decrease with increasing SCN loading, and concurrently, σ is enhanced. The relative change of these three parameters upon SCN addition are most pronounced for the 1:2 PEGDA/LiTFSI system in which B decreases from 9.8 kJ/mol at 10 wt % to 3.8 kJ/mol at 40 wt % with increasing SCN concentration. Furthermore, it can be noted in Figure 10 that the ionic conductivity trend deviates from the typical curve of VTF to a straight line with increasing SCN concentration, implying lithium ion transport seemingly decouples from the network chain motion, which presumably comes from the dissociation of lithium salt and/or lithium ion−polyether complex by the nitrile groups of SCN. The reduction of B may be attributed to decoupling of ionic transport from polymer chain dynamics. Besides ionic conductivity, the transference number of lithium ion (tLi+) is another important transport parameter of the electrolyte, which is defined as the ionic conductivity contributed by the lithium cation relative to the total sum of ionic conductivity due to the transport of both cations and anions. An ideal transference number of unity represents exclusive cation transport, which is highly desirable for high power battery applications. In the case of liquid carbonate plasticized polymer electrolyte (PEO/LiTFSI) (with EO/Li = 15:1, the LiTFSI mole fraction is approximately 6%, which thus corresponds to complexation-unsaturated regime I), molecular dynamics simulation by Wu et al.10 revealed that it is the faster TFSI diffusion that mainly contributes to enhanced ionic conductivity with carbonate addition, leading to the reduction of transference number. As mentioned earlier, the 60/40 (1:2 PEGDA/LiTFSI)/SCN PEM showed superionic conductivity I


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Figure 11. FT-IR spectra for SCN-plasticized crosslinked PEGDA/LiTFSI PEMs with fixed weight ratios of 3:1, 1:1, and 1:2 and SCN concentrations ranging from 0 to 40 wt %.

Figure 12. Loss tan δ vs temperature for three PEM systems with weight ratios of PEGDA/LiTFSI fixed at (a) 3:1, (b) 1:1, and (c) 1:2 with varying SCN concentration up to 40 wt %, and (d) the corresponding tan δ-Tg of plasticized complexed PEMs and plasticized neat PEGDA network (note: the higher temperature Tg was used if two relaxations exist).

3.4. Dynamic Mechanical Analysis of Crosslinked PEGDA/ LiTFSI/SCN PEMs. Next, the effect of SCN plasticization on polymer chain relaxations was investigated by means of DMA. In panels a−c of Figure 12 are shown tan δ peaks of the three aforementioned PEM systems analyzed as a function of ascending order of temperature. The tan δ peak corresponding to the glass transition of the unplasticized 3:1 PEGDA/LiTFSI PEM is located at 1 °C. Upon increasing the PEGDA/LiTFSI ratio to 1:1, the tan δ peak splits into two peaks with the minor peak located at 9 °C and the major peak at 23 °C. These dual tan δ peaks further move to higher temperatures of 13 and 48 °C, respectively, as the PEGDA/LiTFSI ratio increases to 1:2. With increasing SCN content to 10−30 wt %, the tan δ peaks shift systematically to lower temperatures, which is attributable to the SCN plasticization. However, when the SCN

concentration is equal to or above 30 wt %, multiple transitions can be noticed in Figure 12a and b corresponding to the crystal-plastic crystal transition of SCN at −25 °C in both PEGDA/LiTFSI ratios. The very broad relaxation peak corresponding to the SCN plastic crystal phase eventually melts away at a higher temperature of 65 °C for 60/40 (3:1 PEGDA/LiTFSI)/SCN PEM in Figure 12a and 42 °C for 60/ 40 (1:1 PEGDA/LiTFSI)/SCN PEM in Figure 12b. The tan δTg of various plasticized PEMs were further plotted in Figure 12d, which shows a similar trend as that for DSC-Tg in Figure 9a, indicating enhanced polymer chain motions with SCN addition, especially for those PEMs having a high amount of LiTFSI. As pointed out earlier, the unplasticized (1:2 PEGDA/ LiTFSI) PEM exhibited two Tgs that correspond to segmental J


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Macromolecules relaxation of the bulk polymer network (referred to as the first peak) and segmental relaxation of polymer chains adjacent to undissociated “filler-like” salt (denoted as a second peak). Therefore, it is of particular importance to track the movement of the dual relaxation peaks in this plasticized PEM system (Figure 12c). Upon 10 wt % of SCN addition, both peaks shifted to lower temperatures, indicating not only bulk polymer chains but also that those chains surrounding undissociated ionic salts can be further plasticized by SCN. When the SCN concentration increases to 20 wt %, the two peaks merge into one, implying a further dissociation of lithium salt due to the extra SCN, which is consistent with the previous finding in the FT-IR investigation. With SCN concentration increasing to 40 wt %, the single Tg moves to a much lower temperature. These dual Tg reductions and merging of the two Tgs indicate that the constrained polymer chain motions due to lithium cation− ether oxygen complexation as well as those chains on the undissociated salt, i.e., “filler” effect, are gradually overcome by the SCN addition, which is taken into effect through profound plasticization of polymer network chains as well as dissociation of the lithium ions. These observations made by DMA and FT-IR characterizations suggest that it is the high polarity (or dielectric constant) of SCN that makes SCN not only capable of dissociating the lithium salt but also weakening the ion−dipole interactions between lithium ion−polyether oxygen, thereby leading to an increasing number of uncomplexed mobile lithium cations. On the same token, the number of dissociated TFSI anions also increases, which in turn acts as complementary plasticizer to SCN for polymer chains, as manifested by further reduction of Tg in the plasticized, complexed PEM network. Therefore, the pronounced improvement of ionic conductivity is not solely due to plasticization of SCN to the polymer chains but also arises from the consequence of ion dissociation from its salt as well as from the lithium ion−ether oxygen complexes by weakening the ion−dipole interaction.

decay of ionic conductivity in regime II may be due to constrained polymer chain motions involving lesser numbers of mobile lithium ions in the unplasticized PEM system. Upon plasticization, polymer chain motion in the plasticized binary PEGDA/SCN network is expedited as manifested by the drastic Tg reduction, which becomes more dramatic in the lithium ion-complexed PEGDA/LiTFSI/SCN network. The drastic suppression of Tg in the ion−dipole complexed network may be ascribed not only to the plasticization effect exerted by SCN but also plasticization by TFSI anion, albeit bulky, which is freed up as manifested by the blue shifts of the anion IR bands such as CF3 asymmetric stretching, SO2 symmetric stretching, and S−N−S asymmetric stretching of the TFSI anion. Moreover, SCN can further dissociate the ionic salt left undissociated by ether oxygen of PEG as well as the ion−ether oxygen complexes. Hence, the lithium cation transference number improves, whereas ionic conductivity of the plasticized PEM increases over several orders of magnitude, even reaching the level of superionic conductivity in several compositions. The analysis of temperature-dependent ionic conductivity by the VTF equation indicates a considerable decrease of the B value (i.e., apparent activation energy) with decreasing Tg. Therefore, ionic conductivity enhancement of the SCNplasticized PEM system may be afforded by both plasticization and dissociation capabilities of SCN, which results in enhanced ion mobility as well as an increase in ion populations. It may be inferred that the highly polar SCN molecule affords not only dissociation of lithium ion−polyether complexes but also can further dissociate lithium salt left undissociated due to the limited availability of ether oxygen at high LiTFSI loadings, which eventually leads to decoupling of lithium ion transport from polymer chain motions.

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AUTHOR INFORMATION

Corresponding Author

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*Department of Polymer Engineering, Polymer Engineering Academic Center, The University of Akron, 250 South Forge Street, Akron, OH 44325. E-mail: [email protected]. Tel.: +1 330-972-6672. Fax: +1 330-972-3406.

CONCLUSIONS The relationship between glass transition temperature (Tg) and ionic conductivity of crosslinked solid PEM, comprised of PEGDA prepolymer and lithium salt LiTFSI, was demonstrated. In the crosslinked PEGDA/LiTFSI PEM system, Tg increases with LiTFSI salt addition. This Tg elevation arises from both lithium ion−polyether oxygen coordination bonding, which acts as transient “physical” crosslinking in addition to the chemical crosslinking of the PEGDA network. Above the threshold concentration of 7 mol % LiTFSI (i.e., 30 wt %), the Tg versus salt concentration trend changes its slope, as a lesser amount of ether oxygen is available to further dissociate the increasing lithium salt concentration. This is to say, all available ether oxygens are used up in the ion−dipole complexation, implying the saturation of polyether oxygen in the complexed network. At the high salt concentration regime (regime II), tan δ peak, corresponding to DMA-Tg, splits into two slopes. Of particular interest is that the lower temperature Tg peak in regime II, exhibiting little or no change with increasing salt concentration, is due to the chain relaxation of the bulk polymer matrix, whereas the higher temperature Tg, attributable to constrained segmental relaxation of polymer chains surrounding the “filler-like” undissociated ionic salt, continues to rise. As a consequence of constrained chain motion, the ionic conductivity of the unplasticized PEM system further declines from 10−6 to 10−9 S/cm. It is reasonable to infer that the faster

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Support of this work by the National Science Foundation (NSF) through NSF-DMR 1161070 is gratefully acknowledged.

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REFERENCES

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