Flexible Organic–Inorganic Hybrid Solid Electrolytes Formed via Thiol

Article pubs.acs.org/Macromolecules

Flexible Organic−Inorganic Hybrid Solid Electrolytes Formed via Thiol−Acrylate Photopolymerization Ji Hu,†,‡ Wanhui Wang,§ Haiyan Peng,† Mengke Guo,† Yuezhan Feng,† Zhigang Xue,*,† Yunsheng Ye,† and Xiaolin Xie*,† †

Key Laboratory for Material Chemistry of Energy Conversion and Storage, Ministry of Education, School of Chemistry and Chemical Engineering, Huazhong University of Science and Technology, Wuhan 430074, P. R. China ‡ School of Materials Science and Engineering and §School of Environmental Engineering and Chemistry, Luoyang Institute of Science and Technology, Luoyang 471023, P. R. China S Supporting Information *

ABSTRACT: Novel flexible organic−inorganic hybrid solid electrolytes with controlled network structures (cnHSEs) were formed via thiol−acrylate photopolymerization under UV irradiation, using a series of thiol-modified silica nanoparticles and poly(ethylene glycol) diacrylate. Because of the outstanding lithium ion mobility that is readily identifiable in the solid-state nuclear magnetic resonance spectra, the flexible cnHSEs demonstrate a relatively high ionic conductivity even at low temperature of −20 °C and up to the maximum of 7.3 × 10−4 S cm−1 at 30 °C. The ionic conductivity herein is higher than that for typical solid polymer electrolytes at the identical temperature. Additionally, the lithium ion transference number is also improved simultaneously up to 0.7 more due to the scavenger effect of silica nanoparticles embedded in the system. Moreover, the cnHSEs show a broad electrochemical stability window, and the fabricated cell comprising Li|cnHSEs|LiFePO4 exhibits a highly reversible electrochemical reaction and stable cycling performance.

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electrodes, resulting in short circuiting and cell failure.3 For decades considerable strategies have been proposed to overcome the underlying safety issues, such as using judiciously electrolyte additives,4−6 applying hybrid liquid electrolytes,7−9 preparing high modulus separators,10−12 and employing solid polymer electrolytes (SPEs), wherein the SPEs are believed by far the most promising media to afford safe solid-state batteries. To form SPEs, poly(ethylene oxide) (PEO) has been extensively explored due to its astonishing capability of forming useful complex with lithium salts as polymer electrolytes.13 Yet, linear PEO is typically in its semicrystalline state; the dominant crystalline domains hamper the mobility of lithium ions through the electrolyte.14 To address this problem, various strategies have been employed to suppress the PEO crystallinity for increasing the ionic conductivity,15−18 which primarily offer

INTRODUCTION Rechargeable lithium metal batteries (LMBs), wherein the metallic lithium acts as the anode material, have attracted wide attention due to their advantageous features. Lithium has the lowest negative electrochemical potential (−3.04 V vs standard hydrogen electrodes), low density (0.534 g cm−3), and extremely high theoretical specific capacity (3860 mAh g−1).1 When utilizing lithium metal as the lithium source in batteries, unlithiated materials can be also employed as the cathodes to fabricate novel lithium batteries with elevated capacity.2 All these outstanding characteristics make LMBs as promising alternatives to lithium-ion batteries (LIBs), as the LIBs with relatively low energy density are not able to meet the application standards in electrical vehicles, autonomous aircrafts, and so forth. Unfortunately, the uneven electrodeposition of lithium ion on the anode of lithium metal during the charge process usually results in the formation of dendrite structures. These dendrites grow irregularly and usually puncture the separator and ultimately connect the counter © XXXX American Chemical Society

Received: January 6, 2017 Revised: February 25, 2017

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DOI: 10.1021/acs.macromol.7b00035 Macromolecules XXXX, XXX, XXX−XXX

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Herein, we applied thiol decorated silica nanoparticle (SiO2SH) with varied grafting densities and poly(ethylene glycol) diacrylate (PEGDA, Mn = 200, 400, or 600) to fabricate organic−inorganic hybrid solid electrolytes with controlled network structure (cnHSEs) via one-pot thiol−acrylate photopolymerization, which combine the advantages of cross-linked polymer and composite materials. Both step-growth polymerization and chain-growth polymerization are considered to occur simultaneously and thus are able to afford controlled cross-linked network structures and improved mechanical performance. The composition of the cnHSEs was adjusted by changing the grafting density of SiO2-SH, molecular weight of PEGDA, and weight ratio between SiO2-SH and PEGDA, respectively. Because of the cross-linked structure in the system, the mechanical properties of the cnHSEs is improved significantly. The ionic conductivity and lithium ion transference number are enhanced simultaneously due to the better synergistic effects between the inorganic/organic segments in the hybrid system. Moreover, the prepared cell on the basis of cnHSEs as separators shows a highly reversible electrochemical reaction and stable cycling performance. The paradigm proposed in the present work paves the way for the design of SPEs with superior performance.

higher mobility of polymer chain and faster cation diffusion. Toward this end, building cross-linking framework as the electrolyte matrix is regarded to be promising for practical applications.19−21 The cross-linked polymer electrolytes exhibit fully amorphous features which promote ionic conductivity,22 and the mechanical strength of electrolytes is also elevated as the result of the tight network. Flexibility is the intrinsic feature of soft polymer materials, which can be improved owing to the ideal combination of plasticity and rigidity of the cross-linked system, simultaneously.23−25 Moreover, doping inorganic components in polymer matrices to fabricate hybrid solid electrolytes (HSEs) is an efficient route to promote the physical/chemical performance of electrolytes further. Weston and co-workers26 added inert fillers, such as α-alumina, to the polymer matrix, and the mechanical characteristics were thus improved. Fumed silica27 and titanium dioxide28 were employed as the additives to improve the electrochemical properties of polymer electrolytes, but the ionic conductivity decreased when incorporating high content of ceramic particles.29 In contrast to the original nature fillers, the additives modified by various chemical segments, such as ionic liquid,7,30,31 sulfonate,32 cellulose,33 etc., had been widely utilized to prepare HSEs for pursuing better practical performances. Very recently, Pan and co-workers reported an ingenious electrolyte composed of cross-linked polyhedral oligomeric silsesquioxane (POSS) and poly(ethylene glycol) (PEG) through thermal initiated amino−epoxy reaction. The corresponding cells are able to cycle approximately 2600 h at a current density of 0.3 mA cm−2, indicating an excellent lithium dendrite growth resistance.34 Nevertheless, POSS is typically expensive and always involves stringent synthesis with unsatisfying yield. On the contrary, it is more convenient and facile to prepare silicon-containing composites by using modified silica nanoparticles directly because the nanosilica can be readily modified by organic functional groups efficiently in a cost-effective way. For instance, vinyl-functionalized silica could play the role of cross-linker to fabricate network electrolyte with enhanced mechanical properties.35 Yet, the formation of desired solid electrolytes with nanosilica remians far beyond satisfying because of the noncontrolled cross-linked network formed via chain-growth polymerization. For the aim of forming cross-linked polymer electrolytes, photopolymerization has been employed as an effective way.36−38 Photopolymerization represents a valuable approach which exhibits unique spatial, temporal, and spectral control capabilities over the predesigned features.39−41 Furthermore, photopolymerization is rapid, reliable, easy processing, costeffective, and environmentally friendly. Thiol−ene photopolymerization, which is accomplished that the CC double bonds radically react with the thiol segments to form crosslinked systems, has become one of the most powerful and versatile synthetic methods in polymer chemistry and materials science due to the high chemoselectivity, rapid reaction kinetics, and mild reaction conditions.42−44 Though thiol− ene systems have been used for the preparation of UV-cured electrolytes and the polymer skeletons are adjusted in many cases,25,45−47 these electrolytes are neat polymer matrices without functional inorganic fillers embedded. Though uniform cross-linked polymer networks are able to form via a stepgrowth polymerization mechanism, these electrolytes show inferior mechanical performance due to the formed thioether group.

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EXPERIMENTAL SECTION

Materials. LUDOX SM colloidal silica (30 wt % suspension in H2O, Sigma-Aldrich), 3-mercaptopropyltrimethoxysilane (MPTMS, 95 wt %, Sigma-Aldrich), poly(ethylene glycol) diacrylate (PEGDA, Mn = 200, 400, or 600, Aladdin), and benzophenone (BP, SigmaAldrich) were used as received. Lithium perchlorate (LiClO4, Aladdin) was dried at 150 °C in vacuum oven for 24 h and stored in glovebox before use. Other chemicals, including alcohol, lithium iron phosphate (LFP), lithium cobalt oxides (LCO), acetylene black (AB), poly(vinylidene fluoride) (PVDF), N-methyl-2-pyrrolidone (NMP), and commercial liquid electrolyte (1 M LiPF6−ethylene carbonate/ ethylmethyl carbonate/dimethyl carbonate) were purchased locally. Deionized (DI) water was used in all experiments. Synthesis and Characterization of SiO2-SH. A series of thioldecorated silica nanoparticles (SiO2-SH) with varied grafting densities were prepared conveniently. In a typical process, 10 mL of MPTMS was dissolved in 40 mL of ethanol to form a transparent solution, and then the solution was dripped into 20 mL of LUDOX SM colloidal silica combined with 120 mL of deionized water using drastic mechanical stirring (700 rpm). The mixture was stirred under reflux at 70 °C for 24 h. The product was separated by centrifugation and washed with ethanol to remove the unreacted MPTMS and then dried at 60 °C under vacuum for 24 h. Variation in the proportion between LUDOX SM and MPTMS caused different grafting densities of SiO2SH. Transmission electron microscopy (TEM, Tecnai G2 20, FEI, Netherlands) was used to evaluate the particle morphology. The thiol content grafted on the surface of silica was estimated by thermogravimetric analysis (TGA) using a STA449F3 Jupiter thermogravimetric analyzer (Netzsch, Germany) in the range from ambient temperature to 800 °C at a heating rate of 10 °C min−1 under N2. All samples were dried at 60 °C under vacuum overnight before testing. Fabrication and Characterization of cnHSEs. The hybrid solid electrolytes with controlled network structures (cnHSEs) were prepared by taking advantage of thiol−acrylate photopolymerization while using BP as the photoinitiator.48 The composition of the cnHSEs was varied by changing the weight ratio and feature of the cross-linker (PEGDA) and SiO2-SH. As a particular fabrication procedure, the SiO2-SH (0.25 g) and PEGDA (2 g) were mixed with BP (0.1125 g, 5 wt % with respect to monomer concentration) and LiClO4 (0.2068 g, EO/Li+ = 16) without any other solvents. The mixture was stirred for 6 h in a vial in dark and then ultrasonicated in B

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Scheme 1. Schematic Diagram of the Preparation Procedure of SiO2-SH and One-Pot Fabrication Route of cnHSEs via Thiol− Acrylate Photopolymerization under UV Irradiation

aluminum mold before detection. All samples were examined by an Olympus microscope coupled to the Raman system. The laser beam with 200 μm of radius was focused on the surface of sample with a 50× objective, and the laser power irradiation over the samples was 0.5 mW. The region from 400 to 1800 cm−1 with a resolution of 0.4 cm−1 was used for all spectra. Solid-state nuclear magnetic resonance (ssNMR) was performed with a Bruker Avance III 500 MHz wide-bore spectrometer (Germany) with a 194.6 MHz of Larmor frequency for 7Li. The 7Li spectra with a single-pulse experiment using a π/12 pulse length of 2 μs were referenced to 1 M LiCl(aq) at 0 ppm and recorded using a 4 mm static NMR probe. Reported line widths corresponding to the full width at half-maximum intensity (fwhm) of the spectral peak were measured at room temperature. To study the lithium electrodeposition in the prepared polymer electrolytes, symmetric Li|Li coin cells (2032 type) containing cnHSEs were fabricated in a glovebox. The cells were evaluated by galvanostatic cycling experiments at 60 °C, and the charge (3 h) and discharge (3 h) cycling at specific current density was run to mimic the process of lithium stripping/plating. In addition, a control symmetric Li cell containing Celgard separator infused liquid electrolyte but without cnHSEs (Li|liquid electrolyte|Li) was also fabricated to compare lithium dendrite growth. The electrochemical stability of the polymer electrolytes was determined via linear sweep voltammograms (LSV) of the Li|cnHSE| SS cells at a 0.1 mV s−1 of scan rate over the 0−6 V range at 60 °C. The lithium ion transference number (tLi+) of all the prepared electrolytes was measured at 60 °C using the method proposed by Bruce49 and by Scrosati.50 The impedance of the lithium cell (Li| cnHSE|Li) was measured before and after the polarization with a dc voltage pulse, ΔV = 20 mV. The initial current is I0, and the current after decaying to a steady state is denoted as Iss. The equation tLi+ = Iss(ΔV − I0R0)/I0(ΔV − IssRss) was used to calculate the tLi+, where R0 and Rss are the interfacial resistances before and after polarization. Preparation and Electrochemical Characterization of cnHSEs-Based Cells. In the experiment of LMBs test, LFP was employed as the active cathode material. The active material powder, AB, and the PVDF binder were mixed with a weight ratio of 8:1:1, and the viscous slurry was cast on aluminum foil and dried in a vacuum oven at 110 °C for 12 h. The electrode film was cut into round with diameter of 16 mm for the coin cells. Furthermore, the cathode used LCO as active material was prepared following a similar procedure. The coin-cell full batteries (2032 type) were assembled using a Li anode and an LFP or LCO cathode with cnHSEs without a separator. The thickness of the cnHSEs and cathode was approximately 100 μm and 30, respectively. The cell assembly was carried out in a glovebox. The charge−discharge and cycling performances were evaluated and found under the potential window between 4.2 and 2.5 V.

water bath for 30 min with an amplitude of 50% to make SiO2-SH disperse in the mixture homogeneously. The solution was then degassed in a vacuum chamber for 1 h to eliminate the tiny amounts of air blended in the reaction system. The viscous slurry was cast on a Teflon mold with a doctor blade and then exposed to a 365 nm UV light. Flexible cnHSEs membranes with an average thickness of approximately 100 μm formed within 5 min; subsequently, the membranes were dried under a high vacuum condition at 60 °C for 8 h to minimize spurious effects by atmospheric moisture before further performance characterization. The cnHSEs with various weight ratios of SiO2-SH and PEGDA were fabricated following a similar procedure. The viscosity of the slurry mixtures consisted of raw materials for preparing cnHSEs was characterized at 30 °C using a rheometer (MCR 302, Anton-Paar, Austria) with two parallel plates with a diameter of 25 mm. A shear rate of 0−1000 s−1 was performed, and the gap between the two plates was set as 0.1 mm. The Fourier transformation infrared absorption spectra (FT-IR) were performed on an FT interferometer (Equinox 55, Bruker, Germany). The crystal structure of the cnHSEs was characterized via X-ray diffraction (XRD, X’Pert PRO, PANalytical B.V., Netherlands) data using Cu Kα radiation (40 mA/40 kV). Scanning electron microscopy (SEM) and energy dispersive X-ray (EDX) mapping were employed to detect the homogeneity of fillers in the chemical composition (Nova NanoSEM 450, FEI, Netherlands). Dynamic thermomechanical analysis (DMA) of the cnHSEs samples was determined with a dynamic mechanical thermal analyzer (Diamond DMA, PerkinElmer, USA) in a tensile mode at a frequency of 1 Hz and a heating rate of 5 °C min−1 in the temperature range of −90 to 150 °C. Differential scanning calorimetry (DSC) was performed on a differential scanning calorimeter (Q2000, TA, USA). The samples were weighed and sealed in aluminum pans under N2. The samples were first heated from room temperature to 170 °C and subsequently cooled to −90 °C at a rate of 10 °C min−1, and the DSC results were collected from the second heating−cooling cycle. Ionic conductivity was determined on an electrochemical test system (Autolab PGSTAT302N, Netherlands) by the two-electrode ac impedance method. Data were collected over a frequency range from 1 MHz to 100 Hz using a sinusoidal amplitude modulation of 10 mV at temperatures ranging from −20 to 110 °C at 10 °C intervals. A conductivity measurement cell containing two stainless steel electrodes was employed, and the samples were held at each temperature for more than 30 min to equilibrate prior to measurement. The ionic conductivity was calculated by using the equation σ = L/(SR), where L is the thickness of electrolyte film, S is the contact area between electrolyte and electrode, and R is the bulk resistance of polymer electrode. The Raman spectra were recorded at a LabRAM HR800 Raman spectrometer (Horiba Jobin-Yvon, France) using a synapse CCD detector with a laser of 532 nm wavelength. The cnHSEs membranes were placed onto the platform of the spectrometer for Raman characterization directly, and LiClO4 powder should be filled into an C

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Table 1. Composition, Glass Transition Temperature, Storage Modulus, Ionic Conductivity, and Lithium Ion Transference Number of cnHSEs cnHSEsa SiO2-SH(15)-8PEGDA400 SiO2-SH(21)-8PEGDA400 SiO2-SH(27)-8PEGDA400 SiO2-SH(27)-6PEGDA400 SiO2-SH(27)-8PEGDA200 SiO2-SH(27)-8PEGDA600

weight ratio (SiO2-SH:PEGDA)

Mnb/g mol−1

Tgc/°C

G′/MPa at 60 °C

400 400 400 400 200 600

− −d −39.7 −43.1 −8.8 −45.7

− −d 1.12 1.96 94.4 1.77

1:8 1:8 1:8 1:6 1:8 1:8

d

d

σ/S cm−1 at 30 °C 8.9 1.6 3.9 2.0 9.4 7.3

× × × × × ×

−5

10 10−4 10−4 10−4 10−5 10−4

tLi+ −d −d 0.70 0.74 −d 0.71

a

LiClO4 was used for all samples, EO/Li+ = 16. bMolecular weight of PEGDA. cObtained from DSC measurements. dNot detected for these samples.

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RESULTS AND DISCUSSION Structure and Morphological Characterization of SiO2-SH and Composition, Thermal Performance, and Mechanical Strength of cnHSEs. To obtain safer all-solidstate batteries, the hybrid solid electrolytes (HSEs) have stimulated much recent interest because of their improved dendrite resistance as well as the elimination of flammable organic solvents.51 In this work, the HSEs with controlled network structure (cnHSEs) are fabricated via one-pot thiol− acrylate radical photopolymerization with thiol-decorated silica nanoparticle (SiO2-SH) and PEGDA. By means of altering grafting density of SiO2-SH, molecular weight of PEGDA and weight ratio between inorganic filler and polymer, the network structure of the HSEs can be easily tailored. The fabrication processes of SiO2-SH and cnHSEs are elucidated in Scheme 1 intuitively. Six types of cnHSEs, abbreviated as SiO2-SH(x)-nPEGDAm, were prepared, and the compositions characteristics are summarized in Table 1. Here x denotes as the weight percentage of thiol component contained in SiO2-SH, which was calculated from the TGA results shown in Figure 1, m

viscosity of the samples increases with a rise in grafting density of SiO2-SH, implying that the diffusion of inorganic nanoparticles is more difficult in a higher grafting density. Notably, the reactive thiol group on the surface of the SiO2 particles allowed an in situ cross-linking reaction with the electrolyte precursor (PEGDA) containing methacrylate functional group by photopolymerization. The prepared SiO2-SH nanoparticle size is estimated to be ca. 20 nm by TEM (Figure S2), and no significant particles aggregation in cnHSEs is observed (Figure S3a). In addition, the spatial distributions of C, O, and Si atoms are obtained from EDX and corresponding SEM images (Figure S3). The results suggest that the Si of SiO2-SH and the C and O atoms in PEGDA are homogeneously distributed throughout the cnHSEs composite film. The uniform distribution of inorganic nanoparticle benefits the mechanical strength and high flexibility of the material.41,53 To understand the chemical structure change before and after the photopolymerization, FT-IR was conducted on the cnHSEs. As shown in Figure 2, the characteristic thiol group at 2570 cm−1 of SiO2-SH and the peak at 1635 cm−1 assigned to the CC band (methacrylate functional group) disappeared in the spectrum of SiO2-SH(27)-8PEGDA200, indicating that no detectable monomer is embedded in the matrix and the thiol− acrylate reaction is complete after UV irradiation. Macroscopically, as demonstrated in Figure S4, the entire cnHSEs membrane is translucent, which is able to be twisted, bent, and rolled up without cracking. In comparison with the PEGDA-based membrane without doping inorganic fillers (Figure S5), cnHSEs exhibit flexibility with good mechanical properties. The results suggest that the prepared cnHSEs have satisfying mechanical strength which can meet the requirements of application in lithium batteries. The crystalline structure of the cross-linked membrane composed of PEGDA600, LiClO4 (EO/Li+ = 16), and SiO2SH(27)-8PEGDA600 was studied by XRD. As shown in Figure 3, the peak at 20.8°, which corresponds to the characteristic diffraction peak of PEG, is high and sharp in the pattern of PEGDA600-based network without inorganic fillers, suggesting that the high degree of crystallinity of the material is obtained. On the contrary, in the case of SiO2-SH(27)-8PEGDA600, the reduction in intensity of PEG main peak along with broadened peak area of the diffraction peak indicates the decrease of crystallinity of PEG in the presence of nanoparticles,54 which results in an improvement of local relaxation and segmental motion of PEG chains by the increase in the amorphous content of these electrolytes, hence enhancing the ionic conductivity.14,55−57 Furthermore, Tg is known to be a crucial factor for the ionic conductivity of polymer electrolytes, which increases due to a strong ion-dipole interaction, whereas the ionic conductivity

Figure 1. TGA results for SiO2-SH with different grafting densities of the thiol functional group anchored on the surface of SiO 2 nanoparticles.

refers to the molecular weight of each PEGDA segment, and n is the weight ratio between SiO2-SH and PEGDA. Since the viscosity is a key factor to significantly impact the diffusion behavior of additives, that further influences the photopolymerization kinetics and segregation of SiO2-SH nanoparticles,52 the effect of grafting density of SiO2-SH on viscosity was studied on the slurry mixture consisting of the composition with fixed weight ratio between inorganic fillers and organic backbone, but changing the thiol content on the surface of silica nanoparticles (Figure S1, Supporting Information). The D

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Figure 4. DSC traces (second heating cycle) showing the Tg of four types of cnHSEs and the Tm of PEGDA600.

insignificant change. However, for the cnHSEs (SiO2-SH(27)8PEGDA200) with 1:8 weight ratio of SiO2-SH to PEGDA and short chain length of PEGDA200, an obvious increase in Tg can be observed in the DSC results. Though longer PEG chains can promote the molecular relaxation of the system, which leads to a lower Tg,60 the decline apparently slows down when the molecular weight of PEGDA achieves a high level, which may be attributed to the presence of PEGDA crystals. The broadening of the glass transition zone is suggestive of different chain dynamics resulting from varying network heterogeneity. On the other hand, in comparison with the DSC measurements of SiO2-SH(27)-8PEGDA400 and SiO2-SH(27)-6PEGDA400, Tg is inclined to be decreased by 3.4 °C when elevating the weight ratio of inorganic nanoparticles, suggesting that the increased amorphous nature and the increased segmental flexibility of the host polymer caused by the filler. Figure 5 shows the storage moduli for the cnHSEs with different weight ratios of SiO2-SH to PEGDA or various

Figure 2. FT-IR spectra of the pristine SiO2, SiO2-SH(27), PEGDA200, and SiO2-SH(27)-8PEGDA200. The spectra with selected wavenumber region between 2700 and 1500 cm−1 are illustrated concomitantly.

Figure 3. Comparison of XRD patterns for the cross-linked membrane composed of PEGDA600 blended with LiClO4 (EO/Li+ = 16) without inorganic additives and with SiO2-SH(27)-8PEGDA600. Figure 5. Curves of storage modulus vs temperature for the cnHSEs samples.

declines due to the lower ion mobility coupled to reduced chain mobility.58,59 The thermal behaviors, evaluated by DSC, of the cnHSEs samples together with the pure PEGDA600 sample are shown in Figure 4. For the pure PEGDA600, an endothermic peak assigned to the Tm is observed at 16.4 °C,58 indicating that the material is subjected to an innegligible liquid−solid phase change. However, in the DSC curves of cnHSEs, only glass transitions can be detected, suggesting that the PEG crystallization is dramatically suppressed in these cnHSEs samples, and most of the PEG chain ends are covalently linked to SiO2-SH.34 In addition, the reduction of Tg correlates inversely to the chain length of PEGDA, and the Tg of materials consisting of PEGDA400 or PEGDA600 shows a relatively

molecular weights of PEGDA, and the storage modulus at 60 °C is also shown in Table 1. As expected, the rubbery storage modulus of SiO2-SH(27)-8PEGDA200 is massively increased compared with other samples. The reduction in mobility of PEGDA200 with short chain length can result in the formation of network skeleton with high stiffness. On the other hand, with increasing the molecular weight of polymer segment or changing the weight percentage of inorganic filler, the storage modulus is almost immutable at the order of magnitude of 106 Pa. These results display similar trends as the DSC measureE

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polymer electrolyte due to the relaxing polymer chains.65 Hence, based on the Tg values shown in Table 1, the higher ionic conductivity is obtained in the cnHSEs composed of PEGDA with higher molecular weight. In addition, two VTF fitted lines of SiO2-SH(21)-8PEGDA400 (red line) and SiO2SH(27)-6PEGDA400 (pink line) cross at approximately 40 °C, and the ionic conduction of SiO2-SH(21)-8PEGDA400 is higher than SiO2-SH(27)-6PEGDA400 at a relatively high temperature, but showing an opposite tendency at low temperature. At a high temperature, the better ionic conductivity of cnHSEs with more PEGDA contents benefits from the polymer chain mobility.66−68 On the other hand, the chain mobility property is less differentiating at a low temperature, and thus the ionic conduction is dominated by the structure of inorganic fillers. In a word, the conductivity performance of cnHSEs is determined by a synergistic effect between grafting density of silica nanoparticle, weight ratio of SiO2-SH to PEGDA, and polymer chain length of PEGDA. Notably, the Arrhenius plots in Figure 6 show that SiO2SH(27)-8PEGDA600 is characterized by the highest ionic conductivity of 7.3 × 10−4 S cm−1 at 30 °C, which is comparable to and even better than the reported values in cross-linked solid polymer electrolytes and hybrid solid electrolytes.10,34,69−73 To study whether the electrical conductivity change discussed above is related to an enhanced mobility of the Li+ ions, the mobility of the Li+ cation in the cross-linked polymer electrolytes was probed using 7Li ssNMR.74−76 Figure 7 shows

ments, implying that the polymer chain length dominates the cross-linked structure performance. Ionic Conductivity and Lithium Ion Mobility Behavior of cnHSEs. Ionic conductivity is a crucial parameter for the application of polymer electrolytes in energy storage devices.18 Figure 6 presents the temperature dependence of ionic

Figure 6. Ionic conductivity of cnHSEs and neat PEGDA600; the solid lines in the figure are fits to VTF temperature dependence.

conductivities for cnHSEs and PEGDA600 doped with LiClO4 (EO/Li+ = 16). Notably, the plots of log σ versus T−1 exhibit a nonlinear relationship for all the electrolyte samples, which is able to be well described by Vogel− Tamman−Fulcher (VTF) empirical equation σ = A exp(−B/(T − T0)),7,61 where A is a pre-exponential factor equivalent to the ionic conductivity at high temperature, while B represents the effective activation energy for coupled ions and local segmental motion and T0 is a parameter correlated to the glass transition temperature, Tg, of the material and is usually about 30 K lower than Tg.62,63 Table S1 (Supporting Information) displays the related parameters A, B, and T0 calculated by a nonlinear leastsquares fitting regression on the experimental data. Because of the phase transition from liquid to solid, the ionic conductivity of PEGDA600 drops dramatically blow its melting temperature, which coincides with those obtained by other works.51,63,64 Very interestingly, the cnHSEs maintain a relatively high ionic conductivity at low temperature, implying that the segmental and chain motions of the tethered oligomers dominate the ionic conductivity mechanism. Significantly, the hybrid electrolytes arrive the same conductivity at high temperatures, implying that polymer chains in the system have similar dynamics. However, the distinction of ionic conductivity between each cnHSEs sample is obvious when lowering the temperature, resulting from the difference in the composition of cnHSEs. First of all, for the cnHSEs with PEGDA400 and 1:8 weight ratio of SiO2SH to PEGDA, the ionic conductivity increases from 8.9 × 10−5 (SiO2-SH(15)-8PEGDA400) to 1.6 × 10−4 (SiO2-SH(21)8PEGDA400) and 3.9 × 10−4 S cm−1 (SiO2-SH(27)8PEGDA400) at 30 °C, respectively, which can be attributed to the elevated amorphous region as the grafting density of silica nanoparticles increases. Nevertheless, for the network structure with the same inorganic filler (SiO2-SH(27)) and PEGDA400, the conductivity increases with increasing the weight percentage of PEGDA, denoting that more insulative additives are able to depress the ionic conductivity, and the segmental mobility becomes more dynamic with the increase of PEGDA contents. And more specifically, a lower Tg of the polymer is favorable to enhance the ionic conductivity of

Figure 7. 7Li ssNMR line widths for cnHSEs and cross-linked PEGDA600 separately.

the plots of the 7Li ssNMR full width half-maximum (fwhm) for each cnHSEs sample and cross-linked PEGDA600 (EO/Li+ = 16) at the room temperature, and the 7Li ssNMR spectra are shown in Figure S6. In general, the 7Li ssNMR peak of crosslinked PEGDA600 is border than others, implying the low mobility of lithium ions. In the case of SiO2-SH, the lower line width reflects an increase in the Li+ mobility, particularly for the SiO2-SH(27)-8PEGDA600 electrolyte system; the fwhm shows the lowest value of 208 Hz. This effect is expected because the homonuclear 7Li−7Li dipole−dipole interaction in the entire composite material is destroyed by doping inorganic fillers.77 The 7Li ssNMR results coincide with those obtained by ionic conductivity measurement, and therefore the lithium ion mobility is aptable to be promoted under the system using cnHSEs with relatively long chain length backbone. Since neutral ion pairs do not contribute to the conductivity, the ionic conduction performance of cnHSEs mainly contributes to the free charge carriers, not only lithium cations but also perchlorate anions in the total system. However, the movement of cations and anions in the dual-ion conductor F

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Macromolecules leads to a lower lithium ion transference number (tLi+) of 0.1− 0.3.78 The accumulation of ions closes to the surface of one electrode and depletion of ions on the other, resulting in a concentration polarization during the charging and discharging process.79 An unperturbed perchlorate anion has a tetrahedral structure, and the nine vibrational degrees of freedom are divided into four normal modes of vibration, which are Ramanactive; thus, the perturbation of the perchlorate anion can be indicated by Raman spectrocopy.80 Among all the four modes vibration, the fingerprint at approximately 933 cm−1 has been used to study the chemical association behavior of perchlorate anion with other segments in the hybrid mixture.81 As shown in Figure 8, compared with the Raman spectrum of neat LiClO4,

addition, the electrodeposition of lithium in the cnHSEs was studied by galvanostatic cycling experiments of Li|cnHSE|Li symmetric cells at 60 °C. The current density was fixed during the procedure, and the voltage became stable immediately when a stable Li/cnHSE interface was formed. However, after protracted lithium ion stripping/plating process, the cell will be short-circuited as the result of lithium dendrite grown on the rough surface of lithium foil. The typical galvanostatic cycling curves of the lithium cells fabricated from SiO2-SH(27)8PEGDA600 demonstrated in Figure 9a indicates that the cells

Figure 8. Raman spectra of LiClO4, cross-linked PEGDA600 (EO/Li+ = 16), and SiO2-SH(27)-6PEGDA400 individually. Solid line: experimental point; circle: fitting profile; short dashed line: fitted peak.

Figure 9. Galvanostatic cycling curves of the cell with SiO2-SH(27)8PEGDA600 as the separator at current densities of (a) 0.3 mA cm−2 and (b) 1.0 mA cm−2.

the internal standard peak of perchlorate anion is fitted to reveal three bands at around 929, 932, and 936 cm−1 when the salt blends with cross-linked PEGDA600 (EO/Li+ = 16), which is influenced by the coordination between perchlorate anions and polymer backbone.82 In other words, the phenomenon obtained from the fitting result implies that the free anions can be partly anchored on the skeleton of cross-linked PEGDA600. Moreover, with adding SiO2-SH in the system to fabricate cnHSEs, the characteristic peak of perchlorate anion almost completely disappeared, suggesting that the free anions in the composite are trapped by the ceramic fillers further, which scavenges the lithium ions away from the anions. The effects of inorganic fillers that can immobilize anions and lead to emancipate lithium cations from perchlorate anions have been reported in polymer electrolytes.83−85 Because of the fixed anions, the ionic mobility in cnHSEs is mainly derived from lithium cations dissociated from the anions, and thus the tLi+ of the prepared material can be improved (Table 1). Figure S7 shows the results of dc polarization and ac impedance measurements for SiO2-SH(27)-6PEGDA400 system. The obtained transport number of around 0.7 is noticeably close to the single-ion conductive polymers. Electrochemical Stability and Cell Performance of cnHSEs. The electrochemical stability windows of the cnHSEs were observed from LSV. The anodic current onset in the current−voltage curve is linked to electrochemically oxidized decomposition of the polymer electrolyte. The LSV result of SiO2-SH(27)-8PEGDA600 electrolyte shows that no decompositions in cnHSEs take place below 5.0 V vs Li+/Li (Figure S8), revealing that the cnHSEs with high anodic stability could be potentially applied to high voltage lithium batteries. In

show stable voltage profiles for more than 2100 h at a current density of 0.3 mA cm−2. Remarkably, even under harsh electrochemical condition at the current density of 1 mA cm−2, the cells still keep a stable interfacial property over 350 h (Figure 9b). In comparison, a control symmetric cell with commercial liquid electrolyte is already unstable from the first cycle of the test at the current density of 1 mA cm−2 (Figure S9), which has been reported by other works before.86 Compared with the standard cells obtained from the Celgard separator infused with liquid electrolyte, the cell containing the cnHSEs exhibits remarkably high cycling stability. Moreover, to the best of our knowledge, the cnHSEs prepared in our work have better cycling stability compared with previously reported lithium symmetric cells using solid-state polymer electrolytes.10,86−88 As opposed to the smooth surface of cnHSEs (Figure S3a), which becomes rough with some bulges and folds after cycling as the separator of the experimental cell at a fixed current density of 1 mA cm−2 for 350 h (Figure S10), the morphology change is caused by the growth of lithium clusters or lithium dendrites during the strip-plate test. To confirm the usefulness of the cnHSEs, SiO2-SH(27)8PEGDA600, which has the best ionic conductivity and moderate mechanical properties among all samples, was employed as the electrolyte separator to assemble the LFP/Li battery. Figure S11 shows the electrochemical impedance spectra of the fabricated cell at 30 °C, and the charge− discharge performance was also detected at the same temperature with 0.1 C rate current density. As shown in Figure S12, the LFP/Li battery delivers capacity about only 73 mAh g−1 at the test condition owing to the relatively low ionic G

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Figure 10. (a) Charge−discharge curves, (b) cycle performance of the Li|SiO2-SH(27)-8PEGDA600|LFP battery during galvanostatic cycling at 0.1 C rate, (c) discharge profiles, and (d) specific capacity of the LFP/Li cell at varied current density. The measurement was conducted at 60 °C.

conductivity and large polarization at 30 °C. Figure 10a illustrates the charge−discharge property of the Li|SiO2SH(27)-8PEGDA600|LFP cell cycled at 60 °C with a certain current density as 0.1 C. The initial discharge capacity is observed to be around 140 mAh g−1, and the battery retains a capacity of over 110 mAh g−1 up to 100 cycles, resulting in 79% of capacity retention. A distinct plateau at 3.49 V is shown in the charging curves, which is equivalent to the extraction of lithium ions from LFP at the voltage region for the Fe3+/Fe2+. Similarly, a corresponding voltage plateau at 3.36 V is illustrated in the discharging curves, resulting from lithium ions insertion into the Fe sites. The specific capacity during charging procedure is a little bit higher than that of discharge, and the difference becomes smaller with the charge−discharge process going on, indicating that the stabilization of interface is obtained through the reversible reaction. The comparably low cell resistance is also reflected by the low polarization of about 0.13 V (ΔV) between charge and discharge plateaus. Because of the good reversibility of the redox processed on the polymer battery system, the Coulombic efficiency increases to exceed 90% and even approaches 100% after the second cycle (Figure 10b). It is generally known that the durable rate capability of polymer electrolytes in batteries is a very important performance. Figure 10c depicts the discharge profiles of the Li|SiO2SH(27)-8PEGDA600|LFP cell with the voltage range of 2.5− 4.3 V and the current densities of charge/discharge varied from 0.1 to 0.5 C. Figure 10d shows the specific capacity with five cycles performed at each current density of LFP/Li battery. A reversible discharge capacity of approximate 120 mAh g−1 was obtained at 0.2 C, which was about 86% of the capacity at a rate of 0.1 C, and the battery achieved a reversible discharge capacity of 70 mAh g−1 at 0.5 C. Furthermore, a reversible capacity of 135 mAh g−1 was retained when the LFP/Li battery was charged and discharged at 0.1 C again, indicating that the

cnHSEs-based polymer battery possessed an excellent rate performance. In addition, the cnHSEs can be used in other cathodes with high working potential, such as LCO (Figure S13).

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CONCLUSIONS

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ASSOCIATED CONTENT

We developed a highly bendable hybrid solid electrolyte with controlled network structure (cnHSE) for use in lithium metal batteries. This approach was based on the one-pot cross-linked PEGDA and SiO2-SH, which was produced via a facile condensation route under UV irradiation. The successful fabrication of the SiO2-SH nanoparticles and cnHSEs was verified by analyzing the results of FT-IR, TGA, and DSC. The cnHSEs possess excellent ionic conductivity at room temperature and higher lithium ion transference number. Owing to these advantageous features, the cell fabricated using the new cnHSEs as the separator exhibited a highly reversible electrochemical reaction and stable cycling performance. An outstanding contribution of the present study is the provision of a new method to optimize physicochemical properties and electrochemical performance of solid-state electrolytes by exploiting facile UV-curable controlled network architecture as a polymer matrix.

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.7b00035. Figures S1−S13 and Table S1 (PDF) H

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AUTHOR INFORMATION

Corresponding Authors

*Tel +86 27 87540304; Fax +86 27 87543632; e-mail zgxue@ mail.hust.edu.cn (Z.X.). *Tel +86 27 87540304; Fax +86 27 87543632; e-mail xlxie@ mail.hust.edu.cn (X.X.). ORCID

Haiyan Peng: 0000-0002-0083-8589 Zhigang Xue: 0000-0003-2335-9537 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors are grateful to the National Natural Science Foundation of China (Grants 51622303, 51473056, and 51210004) for support of this work. We also specially thank Dr. Yao Yu, working at Wuhan High Magnetic Field Center, for his assistance with the 7Li ssNMR measurements.

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DOI: 10.1021/acs.macromol.7b00035 Macromolecules XXXX, XXX, XXX−XXX