Effect of Side-Chain Branching on Enhancement of Ionic Conductivity

Subscriber access provided by READING UNIV

Article

Effect of side chain branching on enhancement of ionic conductivity and capacity retention of copolymer electrolyte membrane Guopeng Fu, and Thein Kyu Langmuir, Just Accepted Manuscript • DOI: 10.1021/acs.langmuir.7b03449 • Publication Date (Web): 17 Nov 2017 Downloaded from http://pubs.acs.org on November 18, 2017

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a free service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are accessible to all readers and citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

Langmuir is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

1

Effect of side chain branching on enhancement of ionic conductivity and capacity retention of solid copolymer electrolyte membrane Guopeng Fu and Thein Kyu* Department of Polymer Engineering, University of Akron, Akron, OH, 44325

Abstract Low current drain driven by the low ionic conductivity of solid polymer electrolyte is one of the major obstacles of solid state battery. In an effort to improve the ionic conductivity of solid polymer electrolyte membrane (PEM), polyethylene glycol diacrylate (PEGDA) and monofunctional polyethylene glycol methyl ether acrylate (PEGMEA) were copolymerized via photopolymerization to afford PEGDA network with dangling PEGMEA side chains. By attaching PEGMEA side branches to the PEGDA network backbone, the glass transition temperature (Tg) was found to decrease, which may be controlled by relative amounts of PEGMEA and PEGDA. Concurrently, the ionic conductivity of co-polymer electrolyte membrane (co-PEM) consisting of lithium tetrafluorosulfonylimide (LiTFSI) salt and succinonitrile (SCN) plasticizer in the PEGMEA-co-PEGDA copolymer network was enhanced with increasing PEGMEA side branching. The relationship between the network Tg and ionic conductivity of the branched co-PEM was analyzed in the context of Vogel-Tammann-Fulcher (VTF) equation. The plasticized branched co-PEM network exhibited room temperature ionic conductivity at a superionic conductor level of 10-3 S/cm. Of particular importance is that excellent capacity retention at a high current rate (2C) in charge/discharge cyclings of Li4Ti5O12/co-PEM/Li and LiFePO4/co-PEM/Li half-cells was achieved. This improved charge retention may be attributed to lower frictional surfaces of the electrodes afforded by side brushes, which probably alleviates formation of irreversible reaction by-products at the electrode/electrolyte interface. Keywords: Flexible polymer electrolyte membrane, copolymer network, side-chain branching, superionic conductivity

ACS Paragon Plus Environment

Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 2 of 32

2 *Corresponding author: [email protected] Introduction Solid polymer electrolyte (SPE), which is a solid solution of alkali metal salt in polymer, has been extensively explored for potential applications in lithium ion batteries. By virtue of non-flammability and chemical stability, polyethylene oxide (PEO) has been considered as a potential candidate for fabricating SPE1-4. However, the ionic conductivity of these thermoplastic PEO based SPEs is inherently poor (i.e., in the range of 10-9 to 10-6 S/cm at ambient temperatures). The low ionic conductivity has led to low current drains or low power density in battery performance, which hinders the application of solid-state battery. Moreover, the SPE membrane, albeit self-standing, is very fragile. To overcome the aforementioned deficiencies, thermoset polymers such as crosslinkable polyethylene glycol diacrylate (PEGDA) has been employed as a polymer precursor for fabricating solid state polymer electrolyte membrane (PEM) networks

5-7

. The mechanical strength of the cured PEM consisting of PEGDA/lithium salt mixture has

improved considerably in a manner dependent on molecular weights between the crosslinked points (or chemical junctions). However, the ionic conductivity of such binary PEM system remains low due to strong ion-dipole complexation between lithium cation and ether oxygen of PEG that reduces the ion transport. In solid-state PEM, the ion transport depends not only on the mobility of the polymer host, but also on the strength of the ion-dipole complexation as the lithium cation traverses through its complexation sites with the ether oxygen of the PEG segments that undergo large Brownian motions8, 9. To enhance the polymer chain motion, a conventional strategy is to incorporate a carbonate solvent such as ethylene carbonate (EC) and/or dimethyl carbonate (DEC) organic liquids into the PEM system to form a gel-type electrolyte.10-13 However, the addition of organic solvents tends to not only reduce the mechanical strength and integrity, but also raise the flammability issue associated with low flash points of these mixed liquid electrolytes, thereby hindering its application in lithium ion battery. Attaching side brunches to the polymer backbone is an alternative strategy to enhance the ionic conductivity.14,15 To afford a more flexible system, polysiloxane having PEO side chains was first


Page 3 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

3 synthesized by Hall et al.16-18 These branched polymers based PEM systems were highly viscos liquids, the ionic conductivity of which reached a reasonable level of approximately 10-4 S/cm at ambient temperature. Obviously, some improvements in regard to ionic conductivity may have been achieved, but the lack of mechanical strength and their complex chemistry still remain as the drawbacks of such viscos liquids. Recently, polyethylene glycol diacrylate (PEGDA) network, which offers the advantage of outstanding flexibility, extensibility, and strength, becomes attractive for fabricating all solid-state PEM. The ionic conductivity has been elevated to the level of superionic conductor (i.e., 10-3 S/cm at ambient temperature) by plasticizing the network’s Tg using a solid plasticizer such as succinonitrile (SCN), which is also known to be an effective ionizer that weakens the ion-dipole interaction between the lithium cation and the ether oxygen19-21. Since then, it has become a common practice that the ionic conductivity may be improved by lowering the Tg as well as weakening the ion-dipole complexation or combination of both. However, the addition of plasticizer enhances the extensibility, but it reduces the mechanical strength and modulus of the solid PEM. The role of plasticization on electrochemical performance of the solid PEM needs to be understood in order to further improve the ionic conductivity and capacity retention. In this article, monofunctional oligomer, viz., polyethylene glycol methyl ether acrylate (PEGMEA) was copolymerized in conjunction with PEGDA via UV photopolymerization to afford the PEGMEA-co-PEGDA co-networks having dangling PEGMEA side branches22, 23. Subsequently, solidstate PEM co-networks were fabricated by UV-crosslinking the PEGMEA/PEGDA in their pseudoternary mixtures containing SCN plasticizer and lithium bis-(trifluoromethane) sulfonimide (LiTFSI) salt. The effects of polymer side branches on Tg and ionic conductivity of the co-PEM network were systematically investigated as a function of copolymer composition and interpreted in the framework of Vogel-Fulcher-Tammann (VTF) equation. The effects of current rates on the cycling performance and improved charge retention in the anode (Li4Ti5O12/co-PEM/Li) and cathode (LiFePO4/co-PEM/Li) halfcells were discussed.


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 4 of 32

4

Experimental Materials Lithium bis-(trifluoromethane) sulfonimide (LiTFSI) having a purity of 99.9%, succinonitrile (SCN) (>99%), polyethylene glycol diacrylate (PEGDA, average Mn≈700) and polyethylene glycol methyl ether acrylate (PEGMEA, average Mn≈480) were purchased from Sigma-Aldrich Co. Photoinitiator (Irgacure 819, i.e., bis(2,4,6-trimethylbenzoyl)-phenylphosphine oxide) was obtained from Sigma-Aldrich for photo-curing of acrylate groups. The electrode materials such as lithium iron phosphate (LiFePO4) and lithium titanium oxide (Li4Ti5O12) and carbon blacks were purchased from MTI Corp. Electrode binder, i.e., poly(vinylidene fluoride) (PVDF, having an average Mw≈543,000) and 1methyl-2-pyrrolidinone (NMP) solvent were obtained from Sigma-Aldrich. Sample Preparation To eliminate any moisture absorption, the LiTFSI salt was dried at 170 °C under vacuum for 24 h prior to melt blending. By virtue of the similarity of chemical structures of PEGMEA and PEGDA precursors, their binary blends were miscible. Upon further mixing with SCN plasticizer and LiTFSI salt, these PEGMEA/PEGDA/SCN/LiTFSI mixtures remained miscible in a wide composition region as manifested by their pseudo-ternary phase diagram (data not shown). In the fabrication of solid PEM films, appropriate compositions of PEGMEA/PEGDA/SCN/LiTFSI were mixed at room temperature by mechanically stirring in the melt mixtures. A small amount of photoinitiator (Irgacure 819), i.e., 2 wt% with respect to the weight of PEGMEA/PEGDA polymer precursor blend, was added to the SCN/LiTFSI mixtures during melt blending. The homogeneous melt mixture thus obtained was spread on a glass slide within a tape frame spacer having a square shape with a dimension of 10 mm×10 mm. A transparent cover glass was carefully placed on top of the spacer, then lightly pressed, and then performed UV


Page 5 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

5 polymerization by illuminating with a UV lamp (Bondwand 350 nm) at an intensity of 5 mWcm-2 for 15 min in a glovebox under argon atmosphere. Upon UV photopolymerization, transparent solid-state PEM films having various PEGMEA-co-PEGDA/SCN/LiTFSI ratios were obtained. The average thickness of these PEM films was approximately 200 µm. For battery performance evaluations, the working electrodes were prepared by mixing active materials (LiFePO4 or Li4Ti5O12), carbon black, and PVDF in NMP solvent at a weight ratio of 8:1:1. The slurry mixture was coated onto current collectors such as aluminum or copper foil, and then dried under vacuum at 160 oC for 24 h. The 2032 type coin cell was used for half-cell assembly. Lithium metal foil (Alfa Aesar Co.) acted as the counter electrode. The Li4Ti5O12 anode and LiFePO4 cathode half-cells were assembled in a glovebox under argon gas circulation. Sample characterization Complex moduli and loss tangent peaks corresponding to the glass transition temperatures of the PEGDA-co-PEGMEA copolymers were determined as a function of their copolymer compositions using a dynamic mechanical analyzer (DMA Q800 TA instruments) from −80 oC to 10 oC at a heating rate 3 oC. The dynamic oscillatory measurement was carried out at a frequency of 1 Hz and amplitude of 0.1 % relative to the initial sample length. Similarly, the DMA measurements of the UV-cured PEM films containing succinonitrile (SCN) and LiTFSI salt were analyzed from −95 oC to 10 oC at the heating rate of 3 oC/min under nitrogen gas circulation. The PEGMEA/PEGDA copolymer ratios were varied from 0/100 to 75/25 by weight while fixing the equal ratio of SCN plasticizer and LiTFSI salt at 40 wt% each, such that PEGMEA-co-PEGDA/SCN/LiTFSI ratio became 20/40/40 by weight. Fourier transform infrared (FTIR) spectra were acquired by means of an FTIR spectrometer (Nicolet 380) in a transmission mode. The acquired IR spectra represented the average of 64 scans with a resolution of 2 cm-1. The extent of UV photopolymerization of the PEGMEA-co-PEGDA was evaluated


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 6 of 32

6 as a function of copolymer concentration by mimicking the C=C twisting of acrylate group in the range of 780 to 830 cm-1 and the corresponding C=C stretching of acrylate in the range of 1580 to1660 cm-1. Ionic conductivities were determined by using an AC impedance analyzer (HP4192A LF, Hewlett Packard). Samples were sealed between two parallel polished stainless-steel plates with an area of 10 mm × 10 mm and a gap of 1 mm. Sample loading was carried out in a glovebox to prevent moisture absorption, if any. The frequency sweep experiment was carried out from 13 MHz to 5 Hz at a voltage of 10 mV in amplitude. Temperature ramping was performed from room temperature to 120 oC using a homemade heating chamber with the aid of a temperature controller (E5AK, Omron). Galvanostatic charge/discharge cycling tests of the half-cells were conducted using MTI 8channel battery cycler (MTI Corp.). The operation voltages for the LiFePO4/co-PEM/Li and Li4Ti5O12/co-PEM/Li cells were in the ranges of 2.5 ~ 4.2 V and 1.0 ~ 2.5 V, respectively. All battery tests were performed at room temperature (i.e., 22 oC).

Results and Discussion Synthesis and characterization of PEGDA co-network with polyethylene glycol methyl ether acrylate (PEGMEA) side branches The monofunctional oligomer (or precursor), viz., polyethylene glycol methyl ether acrylate (PEGMEA) and low molecular weight PEGDA precursors were mixed at various ratios in the liquid state. Subsequently, following the procedure reported earlier22,

23

, these isotropic liquid mixtures were

copolymerized by irradiating with uniform UV light at 5 mWcm-2 for 15 min. Infrared spectra were acquired in a transmission mode to determine the extent of acrylate double bonds being reacted under the above UV-curing condition. Upon UV polymerization, a chain-growth polymerization between the C=C bonds of PEGDA and PEGMEA has occurred leading to the formation of co-network membranes, which are flexible and stretchable. It should be pointed out that the photopolymerization reaction is


Page 7 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

7 instantaneous and thus the reaction rates of difunctional and monofunctional acrylates may be comparable at least in the initial period UV irradiation. However, given the large steric hindrance of difunctional acrylates, the reaction rates of first and second groups of the diacrylates are expected to be different, which in turn would affect subsequent reactions including the dark reaction. By virtue of similarity of the chemical structures and functionality of the PEGMEA versus PEGDA constituent, the precursor mixture is completely miscible before the photopolymerization and also remains miscible after the curing reaction as manifested by the optical clarity of the PEM films, although by no-means a proof. As pointed earlier, the fast initial photopolymerization rate in the miscible state of the precursor mixtures probably affords a co-network formation, whereby PEGMEA chains are attached as dangling side branches onto the PEGDA network in a manner dependent on the relative copolymer ratios as depicted in the reaction scheme in Figure 1(a).

Figure 1. (a) Schematic representation of UV-crosslinking of PEGMEA-co-PEGDA network, exhibiting PEGMEA side branching on the PEGDA network. The Infrared spectra of neat PEGDA oligomers, PEGMEA oligomers and the PEGMEA-co-PEGDA network with various compositions: (b) the enlarged region between 780 and 830 cm-1 associated with the C=C twisting of acrylate group and (c) the enlarged region between 1580 and 1660 cm-1 corresponding to the C=C stretching of acrylate group, showing the disappearance of the C=C double bonds upon photopolymerization. Figures 1 (b) and 1(c) exhibit the enlarged regions of the infrared spectra, where the characteristic peaks corresponding to the C=C bonds of the acrylate group in the PEGDA oligomers were fully consumed, suggestive of the PEGDA network formation during UV crosslinking. Similarly, the IR


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 8 of 32

8 band at 1635 cm-1 belonging to the acrylate group of PEGMEA disappeared indicating that most of the C=C bonds of the acrylate end groups were used up in forming the PEGMEA-co-PEGDA network, i.e., the PEGMEA side groups may be covalently linked to the PEGDA network, thereby forming dangling side chains as depicted in Figure 1(a).

Relationships between the glass transition temperature and ionic conductivity Dynamic mechanical analysis (DMA) experiments were performed as a function of the PEGMEA-co-PEGDA network compositions over a temperature range of −75 oC to 10 oC. Figure 2 exhibits the variation of storage modulus and loss tan δ with temperature for various copolymer ratios. It should be pointed that the pure PEGMEAs by themselves cannot form a solid film during photoreaction, the DMA experiment of the PEGMEA/PEGDA 100/0 was not performed. At low temperatures, the neat PEGDA network (indicated by 0/100 in Figure 2 behaved like a glassy solid with a very high storage modulus (i.e., over 1 GPa)), but the storage modulus dropped dramatically for two orders of magnitude when the temperature reached around −20 oC. The corresponding loss tangent showed a pronounced peak approximately at −20 oC attributable to the transformation from a glassy state to a rubbery state,24, 25 which may be regarded as the glass transition temperature (Tg) of the crosslinked PEGDA network26. The same phenomena can be witnessed in other PEGMEA-co-PEGDA compositions, showing the systematic movement of both storage moduli and tan δ peaks to lower temperatures with increasing PEGMEA side branches.


Page 9 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

9

Figure 2. Dynamic mechanical analysis of crosslinked PEGMEA-co-PEGDA film showing the variation of (a) storage modulus and (b) tan δ as a function of copolymer composition As shown in Figures 2(a) and (b), the Tg exhibited a systematic decline from −20 oC to −39 oC with increasing PEGMEA to 75%. As manifested in the systematic shift of tan δ with composition corresponding to the Tg, it may be hypothesized that PEGMEA/PEGDA mixtures are probably miscible in the amorphous state (i.e., before crosslinking) as well as after crosslinking, suggestive of the formation of a random co-network. In the cured 75/25 PEGMEA/PEGDA composition, there appears a minor shoulder in the vicinity of −10 oC, which may be ascribed to the relaxation of the PEGMEA dangling side chains.27,28 The Tg depression trend indicates increasing segmental motions with increasing PEGMEA side branches of the crosslinked PEGDA network.


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 10 of 32

10

Figure 3. The variations of (a) storage moduli and (b) tan δ versus temperature of the UV-crosslinked 20/40/40 PEGMEA-co-PEGDA/SCN/LiTFSI PEM for various PEGMEA/PEGDA weight ratios as measured by DMA. Note that the ratios of SCN and LiTFSI were both fixed at 40 wt% each. As mentioned above, PEGMEA and PEGDA pre-polymers were further mixed with SCN solid plasticizer and LiTFSI salt. Subsequently, these pseudo-ternary blends were crosslinked under uniform UV-light exposure to form PEM co-network films. Figure 3 exhibits the variations of the storage modulus and tan δ of the PEMs as a function of PEGMEA/PEGDA weight ratios. Compared to the PEGMEA-coPEGDA copolymer network, the moduli of the PEGMEA-co-PEGDA/SCN/LiTFSI PEMs were reduced for approximately one-half at the glassy state of –80 oC and one order of magnitude at the Tg of about −20 o

C. Moreover, the Tgs of all PEMs were depressed to a lower temperature for approximately 20 oC due to

the plasticization effect afforded by SCN.10, 29 Similar to the trend of the neat co-polymer networks, the Tg


Page 11 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

11 of the PEMs shifted to a lower temperature with increasing PEGMEA concentration. The Tg of PEGDA/SCN/LiTFSI (PEM) was located initially at −37oC that shifted to −55 oC when the amount of PEGMEA was increased to 15 wt%. Next, ionic conductivities were determined from the complex plane impedance plots of PEGMEA/PEGDA/SCN/LiTFSI PEMs, i.e., imaginary impedance Z” versus real component Z’ plots for the PEM with different PEGMEA/PEGDA ratios, while the weight ratios of SCN and LiTFSI were both fixed at 40%. The bulk resistance of the PEM was obtained at Z”=0 on the real component Z’-axis of the Nyquist plot (as shown in Figure 4(a)). With increasing amount PEGMEA of the PEM, the bulk resistance of the PEM declined from 152 Ω for the 20/40/40 PEGDA/SCN/LiTFSI to 83 Ω for the (15/5)/40/40 PEGMEA-co-PEGDA/SCN/LiTFSI.

Figure 4. (a) Nyquist plots of PEGMEA/PEGDA/SCN/LiTFSI with various PEGMEA/PEGDA ratios measured at 22 oC; (b) Variation of ionic conductivity of the PEMs with reciprocal absolute temperature for various PEGMEA/PEGDA ratios. Note that 0/20 represents the neat PEGMEA matrix, whereas 20/0 represents the neat PEGDA network. The experimental data points of each composition were fitted by the solid curves calculated in the framework of the VTF equation.


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 12 of 32

12 Figure 4(b) exhibits the plot of ionic conductivity versus reciprocal absolute temperature as a function of PEGMEA/PEGDA ratio in the PEM containing SCN plasticizer and lithium salt. Note that the SCN plasticizer and LiTFSI salt concentrations were both fixed at 40 wt% while varying only the conetwork compositions. The ionic conductivity of 20/40/40 PEGDA(0/20)/SCN/LiTFSI at room temperature is 0.9×10-3 S/cm at 33 oC and it reaches 7.7×10-3 S/cm at 115 oC. With the addition of PEGMEA, the ionic conductivity of the PEM increases dramatically, i.e., the ionic conductivity of 15/5 PEGMEA/PEGDA ratio reached the value of 1.7×10-3 S/cm at 36 oC and the superionic conductor level of 10-2 S/cm at 120 oC. The relationship between Tg and the temperature-dependent ionic conductivity of the PEM was analyzed in the context of the empirical Vogel-Tammann-Fulcher (VTF) equation 30-32, viz., = − / −

(1)

where σ0 is a prefactor; T0 is the ideal glass transition temperature. Either B or B0=B/R, i.e., the apparent activation energy B is normalized by a gas constant (R) may be used. The fitted curves using various VTF parameters were shown in Figure 4 (b). Table 1 summarizes the best-fit VTF parameters obtained from temperature-dependent ionic conductivity plots in Figure 4 (b). Table 1. Tg, room temperature ionic conductivity, and fitting parameters of the VTF equation to account for the temperature-dependent ionic conductivity of PEGMEA/PEGDA/SCN/LiTFSI for various PEGMEA/PEGDA ratios (Figure 4). Both concentrations of SCN and LiTFSI were fixed at 40 wt%. r2 stands for the standard of deviation from the VTF regression analysis. PEGMEA/ PEGDA 0/20

Tg (oC)

σRT (S/cm)

T0 (oC)

σ0 (S/cm)

B0(K)

-36.6

6.17×10-4

-65.9

0.230

5/15

-46.2

8.42×10-4

-73.2

10/10

-48.8

1.16×10-3

15/5

-55.2

20/0

r2

560.7

B (KJ/mol) 4.66

0.999

0.191

543.2

4.52

0.997

-76.5

0.142

530.7

4.41

0.999

1.25×10-3

-79.7

0.136

523.1

4.35

0.998

1.54×10-3

-81.9

0.103

477.1

3.96

0.999


Page 13 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

13 In Table 1, the value of T0 is found to be lower for 20 oC ~ 30 oC than that of the Tg as measured by the peak position of tan δ. The ideal glass transition temperature, T0, is the reference temperature at which the configurational entropy is zero and also the ionic conduction via the PEM network is completely frozen, and thus in literature, it is usually observed to be lowered for 30 oC ~ 50 oC below the experimental Tg 30,33. Both Tg and T0 decrease with increasing PEGMEA concentration, as discussed previously. The apparent activation energy B (or B0) can be regarded as the energy barrier for the rotational motion of the polymer segments. The apparent activation energy B is estimated to be 4.66 KJ/mol

for

the

neat

PEGDA

based

PEM

labeled

as

the

20/40/40

PEGMEA-co-

PEGDA(0/20)/SCN/LiTFSI composition. By increasing the weight ratio of PEGMEA/PEGDA, the activation energy B decreases and reaches 3.96 KJ/mol for the neat PEGMEA based non-solid PEM denoted as PEGMEA-co-PEGDA(20/0)/SCN/LiTFSI. This implies that increasing PEGMEA loading in the copolymer ratio of the PEM reduces the crosslinked density of the polymeric networks, and thus lowers the activation energy for the polymer segmental motions (see Table 1). Consequently, with increasing side chain branches of the PEM, the ionic transport is expedited which may be attributed to the lowering of the activation energy for polymer segmental motions.

Figure 5. The variation of room temperature ionic conductivity of the PEM as a function PEGMEA/PEGDA ratios and Tg of the PEM as determined by the peak positions of tan δ.


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 14 of 32

14

Figure 5 exhibits the Tg of the PEMs obtained from the peak positions of tan δ in comparison with room temperature ionic conductivity as a function PEGMEA weight percent. The inverse relationship between Tg and ionic conductivity can be established, i.e., the ionic conductivity of PEMs increases with increasing amount of PEGMEA concentration, whereas Tg decreases. As previously pointed out, introducing PEGMEA side branches into the PEGDA network loosens the network and concurrently boosts the polymer segmental motion, which results in the suppression of Tg to lower temperatures. According to the molecular dynamic simulations, cations are transported through the complexation sites between the lithium ions and the ether oxygen of PEO chains.34 According to our previous work35, given the strong coupling between the lithium cations and the ether oxygen, the Tg of the binary crosslinked PEG/lithium salt (LiTFSI) mixture has increased for about 50 oC (from −40 to 10 oC), which makes the ionic conductivity to be extremely low, i.e., 10-9 S/cm (at 30% LiTFSI) ~ 10-6 S/cm (at 67% LiTFSI). As demonstrated by He et al., upon plasticization with SCN, not only the PEGDA network is being plasticized, but also the ion-dipole complexation gets weakened or dissociated.35 Consequently, the Tg is reduced to −60 oC while the ionic conductivity has increased to the superionic conductor level of 10-3 S/cm at ambient temperature. Hence it may be inferred that the lithium ion transport is governed by the strength of the ion-dipole complexation assisted by the segmental motions of the PEG network mixture.35 The enhanced polymer chain segmental motion due to attaching PEGMEA side branches to the PEM network is certainly beneficial to the enhancement of the ionic conductivity. Stability analysis in half-cells by cyclic voltammetry We shall now turn our attention to the stability analysis of cathode and anode half-cells based on the PEM containing 10 wt% PEGMEA, 10 wt% PEGDA, 40 wt% SCN and 40 wt% LiTFSI, denoted as (10/10)/40/40 PEGMEA-co-PEGDA/SCN/LiTFSI, against Li4Ti5O12 and LiFePO4 electrodes in a halfcell with the Li metal as the counter electrode. This PEM composition was chosen because of the balance between the high ionic conductivity and reasonable mechanical properties (see Figures 2 and 3).


Page 15 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

15 Figure 6(a) shows the variation of capacity and Coulombic efficiency as a function of number of cycles for the (10/10)/40/40 PEGMEA-co-PEGMEA/PEGDA/SCN/LiTFSI against the Li4Ti5O12 anode at various current densities. The half-cell battery was cycled at C/4 initially and the capacity of approximately 140 mAh/g was obtained suggestive of the compatibility of the PEM co-network to the Li4Ti5O12 anode.

Figure 6. Cyclability of the PEM (PEGMEA/PEGDA/SCN/LiTFSI 10/10/40/40) in (a) Li4Ti5O12/PEM/Li half-cell (b) LiFePO4/PEM/Li half-cell at various current densities. Note that the denominator of the C stands for the time (in hour) required completing 1 charge/discharge cycle. The battery test was carried out at room temperature of 22 oC.

The rate dependent capability test was performed from 6th and 30th cycles at various current densities. For C/2, 1C and 2C the half-cell delivered the reversible capacity of 127, 114 and 98 mAh/g, respectively. After the C-rate test, the current was switched back to C/4 and the capacity value was recovered to the initial value of 140 mAh/g and stabilized thereafter. Similar to Figure 6 (a), Figure 6 (b) exhibits the cycling behavior of PEM containing 10 wt% PEGMEA, 10 wt% PEGDA, 40 wt% SCN and 40 wt% LiTFSI against LiFePO4 electrode at various current densities. The initial capacity of the battery was 126 mAh/g. With increasing current density, the specific capacity of the battery was reduced to 120, 105, 80 mAh/g. After 25th cycle, the current density was reverted back to C/4 and the initial capacity of 135 mAh/g was recovered, suggesting the compatibility between the present PEM and the cathode.


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 16 of 32

16 Benefited from the high ionic conductivity of the PEM, both half-cells reveal high capacity retention at various current rates, thereby making the present co-PEM a good candidate for high power-density battery applications. To evaluate the electrochemical performance of PEM in lithium ion battery, prototype half-cells were fabricated using the same co-PEM composition containing 10 wt% PEGMEA, 10 wt% PEGDA, 40 wt% SCN and 40 wt% LiTFSI. Figure 7(a) exhibits the charge-discharge profiles of the Li4Ti5O12/coPEM/Li half-cell, obtained at a current density of 140 mAh/g corresponding to the current rate of C/4 within the cut-off voltages of 1.0 V and 2.5 V. The observed stable voltage plateaus at 1.63 V and 1.55 V correspond to the lithium ion insertion and desertion processes, respectively. Figure 7(b) shows the behavior of specific capacity retention of the PEM ((10/10)/40/40 PEGMEA-co-PEGDA/SCN/LiTFSI) against the Li4Ti5O12 anode half-cell. The discharge specific capacity at the first cycle was approximately 140 mAh/g and remained stable with increasing number of cycles up to 50 cycles tested. The Coulombic efficiency is virtually invariant at approximately 100% for 50 cycles, indicating that the aforementioned PEM appears compatible with the Li4Ti5O12 electrode.


Page 17 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

17

Figure 7. Galvanostatic charge-discharge measurement of the PEM (PEGMEA/PEGDA/SCN/LiTFSI 10/10/40/40) at the current rate of C/4. The battery test was undertaken at room temperature (22 oC). (a) Voltage as a function of the specific capacity of Li4Ti5O12/PEM/Li half-cell; (b) Variation of specific capacity as a function of cycles for Li4Ti5O12/PEM/Li half-cell; (a) Voltage as a function of the specific capacity of LiFePO4/PEM/Li half-cell; (b) Variation of specific capacity as a function of cycles for LiFePO4/PEM/Li half-cell.

Figure 7 (c) and (d) exhibit the corresponding charge-discharge profiles of the LiFePO4/PEM/Li half-cell containing the co-PEM ((10/10)/40/40PEGMEA-co-PEGDA/SCN/LiTFSI). The half-cell was cycled at a current rate of C/4 with the cut-off voltages of 2.5 V and 4.0 V. The discharging and charging voltages show the stable plateaus at approximately 3.38 V and 3.48 V, respectively, which may be attributed to lithium ion insertion and desertion. As shown in Figure 8(a), the specific capacity remains invariant at about 140 mAh/g, suggesting excellent capacity retention up to 50 cycles tested. The corresponding Columbic efficiency remains virtually invariant at about 100%, which implies that the present co-PEM is compatible to the LiFePO4 cathode. It may be reasonable to infer that the plasticizing


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 18 of 32

18 strategy via attaching the PEGMEA side chain branches to the PEGDA network in the co-PEM affords both high room temperature ionic conductivity in excess of 10-3 S/cm and excellent capacity retention at low and high current rates during charge/discharge cyclings of Li4Ti5O12/co-PEM/Li and LiFePO4/coPEM/Li half-cells. As pointed above, the capacity retention of the present co-PEM system is stable at almost 100% for both Li4Ti5O12 and LiFePO4 electrodes, therefore the attachment of dangling PEGMEA side branching to the PEGDA network has helped in preventing irreversible solid electrolyte interface (SEI) layer formation due to highly mobile nature of the side branches of the PEM network at the electrode interface. The Fourier transformed infrared (FTIR) spectra of the surfaces of the present co-PEM system are virtually the same before and after the charge/discharge cycling and also the corresponding scanning electron microscopy (SEM) images appear to be similar (data not shown) suggestive of the lack of interface reaction, although by no means a proof. It is reasonable to infer that the chemical composition and/or morphology of the PEM/electrode interface are likely to remain unchanged before and after cycling. It may be contrasted with our previous findings in a complementary PEGDA based PEM (i.e., 20/40/40 PEGDA/SCN/LiTFSI) system,36 wherein the capacity retention has reduced to 80% after 50 cycles in the LiFePO4 cathode and lithium foil electrode half-cell tests. In that case, one can clearly identify the irreversible reaction at the interface between the electrode and solid PEM during the cycling at an elevated temperature of 60 oC, based on energy dispersive x-ray (EDX) spectroscopy and SEM investigations. However, this capacity retention can be improved to over 95% by merely adding 0.4 wt% lithium bis(oxalato)borate (LiBOB) to the PEM. With this LiBOB modification, the interface morphology before and after cycling becomes similar and also there is no more enrichment of nitrogen and carbon in the EDX investigations. Although the lack of irreversible interface reaction can be confirmed unambiguously in the case of LiBOB modification,36 it is difficult to confirm in the present co-PEM system having dangling side branches, especially when there is no noticeable change in surface morphology or chemical composition before and after cycling. One can only hypothesize that the


Page 19 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

19 dangling side branches of the co-PEM network probably contribute not only to higher chain mobility, but also improves the ionic conductivity as well as capacity retention by preventing deposition of reaction-byproducts at the interface between the electrode and solid PEM. Conclusions In this article, the effect of PEGMEA side chain branching on the enhancement of ionic conductivity and improvement of capacity retention of the PEGMEA-co-PEGDA copolymer electrolyte membranes has been demonstrated. By controlling the relative ratio of di-functional PEGDA and monofunctional PEGMEA oligomers, the network elasticity of the co-PEM in relation to electrochemical properties can be manipulated. Of particular interest is that the glass transition temperature of polymer matrix has decreased with increasing PEGMEA side branches, which in turn expedite the ionic conductivity. The relationship between the glass transition temperature of the copolymer based PEM and ionic conductivity was analyzed in the context of VTF equation and found to be satisfactory. The plasticization strategy through the side chain branching of the PEGDA-PEGMEA copolymer - network affords both high ionic conductivity in excess of 10-3 S/cm at room temperature and excellent capacity retention at various low and high current densities during the charge/discharge cyclings of Li4Ti5O12/coPEM/Li and LiFePO4/co-PEM/Li half-cells. It may be reasonable to infer that the attachment of dangling PEGMEA side branching to the PEGDA network of the PEM has not only improved the ionic conductivity relative to the PEM without the side branches, but also retained the cyclic stability at nearly 100% in the Li4Ti5O12/co-PEM/Li as well as of LiFePO4/co-PEM/Li half-cells. It may be inferred that the dangling side-branches have imparted higher chain mobility at the electrode and electrolyte interfaces; thereby preventing deposition of any reaction by-products at the interface between the electrode and solid electrolyte, which in turn improves the capacity retention. It should be emphasized that the ionic conductivity as well as capacity retention of the present PEGMEA-co-PEGDA PEM have improved considerably over the PEM containing no PEGMEA side chain branches.36 Acknowledgment: Support of this work by NSF-DMR 1502543 is gratefully acknowledged.


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 20 of 32

20

References 1.

Armand, M., Polymer solid electrolytes-an overview. Solid State Ionics 1983, 9, 745-754.

2.

MacCallum, J. R.; Vincent, C. A., Polymer electrolyte reviews. London, Elsevier, 1987

3.

Bruce, P.; Vincent, C., Polymer electrolytes. J.Chem. Soc., Faraday Trans. 1993, 89 (17), 3187-

3203. 4.

Fenton, D.; Parker, J.; Wright, P., Complexes of alkali metal ions with poly (ethylene oxide).

Polymer 1973, 14 (11), 589. 5.

Wright, P. V., Electrical conductivity in ionic complexes of poly (ethylene oxide). British Polym.

J. 1975, 7 (5), 319-327. 6.

Rietman, E.; Kaplan, M.; Cava, R., Lithium ion-poly (ethylene oxide) complexes. I. Effect of

anion on conductivity. Solid State Ionics 1985, 17 (1), 67-73. 7.

Meyer, W. H., Polymer electrolytes for lithium-ion batteries. Adv. Mater.1998, 10 (6), 439-448.

8.

Papke, B.; Ratner, M.; Shriver, D., Conformation and ion-transport models for the structure and

ionic conductivity in complexes of polyethers with alkali metal salts. J. Electrochem. Soc. 1982, 129 (8), 1694-1701. 9.

Ratner, M. A.; Shriver, D. F., Ion transport in solvent-free polymers. Chem. Rev. 1988, 88 (1),

109-124. 10.

Alarco, P. J.; Abu-Lebdeh, Y.; Abouimrane, A.; Armand, M., The plastic-crystalline phase of

succinonitrile as a universal matrix for solid-state ionic conductors. Nat. Mater. 2004, 3 (7), 476-81. 11.

Song, J.; Wang, Y.; Wan, C., Review of gel-type polymer electrolytes for lithium-ion batteries. J.

Power Sources 1999, 77 (2), 183-197. 12.

Michael, M.; Jacob, M.; Prabaharan, S.; Radhakrishna, S., Enhanced lithium ion transport in

PEO-based solid polymer electrolytes employing a novel class of plasticizers. Solid State Ionics 1997, 98 (3), 167-174.


Page 21 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

21 13.

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 Ionics 2002, 149 (1), 29-37. 14.

Turnbull, D.; Cohen, M. H., Free-volume model of the amorphous phase: glass transition. J.

Chem. Phys. 1961, 34 (1), 120-125. 15.

Miyamoto, T.; Shibayama, K., Free-volume model for ionic conductivity in polymers. J. Appl.

Phys. 1973, 44 (12), 5372-5376. 16.

Hall, P.; Davies, G.; McIntyre, J.; Ward, I.; Bannister, D.; Le Brocq, K., Ion conductivity in

polysiloxane comb polymers with ethylene glycol teeth. Polymer Commun. 1986, 27 (4), 98-100. 17.

Fish, D.; Khan, I. M.; Smid, J., Conductivity of solid complexes of lithium perchlorate with poly

{[ω-methoxyhexa (oxyethylene) ethoxy] methylsiloxane}. Die Makromolekulare Chemie, Rapid Commun. 1986, 7 (3), 115-120. 18.

Lee, A. S.; Lee, J. H.; Hong, S. M.; Lee, J.-C.; Hwang, S. S.; Koo, C. M., Ion conduction

behaviour in chemically crosslinked hybrid ionogels: effect of free-dangling oligoethyleneoxides. RSC Advances 2015, 5 (114), 94241-94247. 19.

Echeverri, M.; Hamad, C.; Kyu, T., Highly conductive, completely amorphous polymer

electrolyte membranes fabricated through photo-polymerization of poly(ethylene glycol diacrylate) in mixtures of solid plasticizer and lithium salt. Solid State Ionics 2014, 254, 92-100. 20.

He, R.; Echeverri, M.; Ward, D.; Zhu, Y.; Kyu, T., Highly conductive solvent-free polymer

electrolyte membrane for lithium-ion batteries: effect of prepolymer molecular weight. J. Membr. Sci. 2016, 498, 208-217. 21.

Fu, G.; Dempsey, J.; Izaki, K.; Adachi, K.; Tsukahara, Y.; Kyu, T., Highly conductive solid

polymer electrolyte membranes based on polyethylene glycol-bis-carbamate dimethacrylate networks. J. Power Sources 2017, 359, 441-449. 22.

Liao, C.; Sun, X.-G.; Dai, S., Crosslinked gel polymer electrolytes based on polyethylene glycol

methacrylate and ionic liquid for lithium ion battery applications. Electrochimi. Acta 2013, 87, 889-894.


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 22 of 32

22 23.

Paranjape, N.; Mandadapu, P. C.; Wu, G.; Lin, H. Q., Highly-branched cross-linked

poly(ethylene oxide) with enhanced ionic conductivity. Polymer 2017, 111, 1-8. 24.

Nielsen, L. E., Mechanical Properties of Polymers and Composites, second ed. CRC Press, 1993.

25.

Milhaupt, J. M.; Lodge, T. P.; Smith, S. D.; Hamersky, M. W., Composition and temperature

dependence of monomer friction in polystyrene/poly (methyl methacrylate) matrices. Macromolecules 2001, 34 (16), 5561-5570. 26.

Krumova, M.; Lopez, D.; Benavente, R.; Mijangos, C.; Perena, J., Effect of crosslinking on the

mechanical and thermal properties of poly (vinyl alcohol). Polymer 2000, 41 (26), 9265-9272. 27.

Sanders, J. F.; Ferry, J. D., Dynamic mechanical properties of cross-linked rubbers. VII. Butyl

rubber networks. Macromolecules 1974, 7 (5), 681-684. 28.

Yamazaki, H.; Takeda, M.; Kohno, Y.; Ando, H.; Urayama, K.; Takigawa, T., Dynamic

viscoelasticity of poly (butyl acrylate) elastomers containing dangling chains with controlled lengths. Macromolecules 2011, 44 (22), 8829-8834. 29.

Fan, L.-Z.; Maier, J., Composite effects in poly (ethylene oxide)–succinonitrile based all-solid

electrolytes. Electrochem. Commun. 2006, 8 (11), 1753-1756. 30.

Karatas, Y.; Banhatti, R. D.; Kaskhedikar, N.; Burjanadze, M.; Funke, K.; Wiemhöfer, H.-D.,

Synthesis and modeling of polysiloxane-based salt-in-polymer electrolytes with various additives. The J. Phys. Chem. B 2009, 113 (47), 15473-15484. 31.

Susan, M. A. B. H.; Kaneko, T.; Noda, A.; Watanabe, M., Ion gels prepared by in situ radical

polymerization of vinyl monomers in an ionic liquid and their characterization as polymer electrolytes. J. Am. Chem. Soc.2005, 127 (13), 4976-4983. 32.

Park, M.; Zhang, X.; Chung, M.; Less, G. B.; Sastry, A. M., A review of conduction phenomena

in Li-ion batteries. J. Power Sources 2010, 195 (24), 7904-7929. 33.

Seki, S.; Susan, M. A. B. H.; Kaneko, T.; Tokuda, H.; Noda, A.; Watanabe, M., Distinct

difference in ionic transport behavior in polymer electrolytes depending on the matrix polymers and incorporated salts. J. Phys. Chem. B 2005, 109 (9), 3886-3892.


Page 23 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

23 34.

Müller-Plathe, F.; van Gunsteren, W. F., Computer simulation of a polymer electrolyte: lithium

iodide in amorphous poly (ethylene oxide). J. Chem. Phys. 1995, 103 (11), 4745-4756. 35.

He, R.; Kyu, T., Effect of plasticization on ionic conductivity enhancement in relation to glass

transition temperature of crosslinked polymer electrolyte membranes. Macromolecules 2016, 49 (15), 5637-5648. 36.

He, R.; Peng, F.; Dunn, W. E.; Kyu, T., Chemical and electrochemical stability enhancement of

lithium bis-(oxalato)borate (LiBOB)-modified solid polymer electrolyte membrane in lithium ion halfcells. Electrochim. Acta 2017, 246, 123-134


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 24 of 32

24

Table of Contents Graphic


Page 25 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

Figure 1. (a) Schematic representation of UV-crosslinking of PEGMEA-co-PEGDA network, exhibiting PEGMEA side branching on the PEGDA network. The Infrared spectra of neat PEGDA oligomers, PEGMEA oligomers and the PEGMEA-co-PEGDA network with various compositions: (b) the enlarged region between 780 and 830 cm-1 associated with the C=C twisting of acrylate group and (c) the enlarged region between 1580 and 1660 cm-1 corresponding to the C=C stretching of acrylate group, showing the disappearance of the C=C double bonds upon photopolymerization. 369x132mm (96 x 96 DPI)


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 2. Dynamic mechanical analysis of crosslinked PEGMEA-co-PEGDA film showing the variation of (a) storage modulus and (b) tan δ as a function of copolymer composition 123x196mm (96 x 96 DPI)


Page 26 of 32

Page 27 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

Figure 3. The variations of (a) storage moduli and (b) tan δ versus temperature of the UV-crosslinked 20/40/40 PEGMEA-co-PEGDA/SCN/LiTFSI PEM for various PEGMEA/PEGDA weight ratios as measured by DMA. Note that the ratios of SCN and LiTFSI were both fixed at 40 wt% each. 124x199mm (96 x 96 DPI)


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 4. (a) Nyquist plots of PEGMEA/PEGDA/SCN/LiTFSI with various PEGMEA/PEGDA ratios measured at 22 oC; (b) Variation of ionic conductivity of the PEMs with reciprocal absolute temperature for various PEGMEA/PEGDA ratios. Note that 0/20 represents the neat PEGMEA matrix, whereas 20/0 represents the neat PEGDA network. The experimental data points of each composition were fitted by the solid curves calculated in the framework of the VTF equation. 313x146mm (96 x 96 DPI)


Page 28 of 32

Page 29 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

Figure 5. The variation of room temperature ionic conductivity of the PEM as a function PEGMEA/PEGDA ratios and Tg of the PEM as determined by the peak positions of tan δ. 232x166mm (96 x 96 DPI)


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 6. Cyclability of the PEM (PEGMEA/PEGDA/SCN/LiTFSI 10/10/40/40) in (a) Li4Ti5O12/PEM/Li half-cell (b) LiFePO4/PEM/Li half-cell at various current densities. Note that the denominator of the C stands for the time (in hour) required completing 1 charge/discharge cycle. The battery test was carried out at room temperature of 22 oC. 252x102mm (96 x 96 DPI)


Page 30 of 32

Page 31 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Langmuir

Figure 7. Galvanostatic charge-discharge measurement of the PEM (PEGMEA/PEGDA/SCN/LiTFSI 10/10/40/40) at the current rate of C/4. The battery test was undertaken at room temperature (22 oC). (a) Voltage as a function of the specific capacity of Li4Ti5O12/PEM/Li half-cell; (b) Variation of specific capacity as a function of cycles for Li4Ti5O12/PEM/Li half-cell; (a) Voltage as a function of the specific capacity of LiFePO4/PEM/Li half-cell; (b) Variation of specific capacity as a function of cycles for LiFePO4/PEM/Li half-cell. 232x185mm (96 x 96 DPI)


Langmuir

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Table of Contents Graphic 378x132mm (96 x 96 DPI)


Page 32 of 32

Effect of Side-Chain Branching on Enhancement of Ionic Conductivity

Recommend Documents