Anisotropic Ion Transport in a Poly(ethylene oxide)–LiClO4 Solid State

Jun 19, 2015 - Boyang Liu , Yunhui Gong , Kun Fu , Xiaogang Han , Yonggang Yao , Glenn Pastel , Chunpeng Yang , Hua Xie , Eric D. Wachsman , and ...
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Anisotropic Ion Transport in a Poly(ethylene oxide)−LiClO4 Solid State Electrolyte Templated by Graphene Oxide Shan Cheng, Derrick M. Smith, and Christopher Y. Li* Department of Materials Science and Engineering, Drexel University, Philadelphia, Pennsylvania 19104, United States S Supporting Information *

ABSTRACT: Solid polymer electrolytes (SPEs) have attracted intensive attention due to their potential applications in all-solid-state lithium batteries. Tailoring crystallization is crucial to the design of high performance poly(ethylene oxide) (PEO)−based SPEs. In this paper, we demonstrate that PEO crystal orientation in a PEO−lithium electrolyte results in anisotropic ionic conductivity along and through the crystalline lamellae. This conductivity anisotropy can be further enhanced by incorporating two-dimensional graphene oxide (GO) nanosheets, which retard PEO crystallization, template the crystal orientation, and guide the ion transport, leading to highly anisotropic and conductive nanocomposite polymer electrolytes.

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INTRODUCTION In both Li-metal and Li-ion battery systems, replacing the volatile, flammable liquid organic electrolyte with solid polymer electrolytes (SPEs) having adequate ionic conductivity, high cation mobility, sufficient mechanical properties, and a wide electrochemical window is critical to addressing the safety concerns of the batteries. Poly(ethylene oxide) (PEO) has been demonstrated as an excellent candidate for SPEs due to its high dielectric constant and strong lithium ion solvating ability.1−5 The ion conduction mechanism in PEO is considered as ion hopping along the polymer segments assisted by the ether oxygens. In most cases, it takes place in the amorphous region. Because linear PEO homopolymer tends to crystallize, the ionic conductivity of the SPEs at room temperature is far below the required value (>10−4 S/cm) for battery applications. To achieve considerable ionic conductivity while maintaining sufficient shear modulus to effectively prevent lithium dendrites formation during battery operation,6,7 the typical strategy is decoupling ion conduction and mechanical properties in the materials design using block copolymer (BCP),8−14 polymer blends, polyolefin porous membranes,14−16 holographic polymerized polymer electrolyte membranes,17−19 and nanoparticles.14,20−23 Our recent study has demonstrated that the two intertwined effects of PEO crystallization on ionic conductivity reduction, namely the structural/tortuosity effect and dynamic/tethered chain effect, can be decoupled and quantified using a model PEO single crystal SPE with controlled size, structure, crystallinity and crystal orientation. It further suggested that semicrystalline PEO-based SPEs offer an alternative solution to achieving both high mechanical integrity and moderate ionic conductivity as long as the PEO crystal orientation can be controlled so that the surfaces of PEO lamellae are parallel to the intended ion diffusion direction to minimize the tortuosity effect.24 While the polymer single crystal system can be used as © XXXX American Chemical Society

a model for fundamental study and to fabricate functional materials,25−28 it is not yet feasible for large scale SPE fabrication. Nanoparticles are capable of templating polymer crystal growth.29,30 In this study, a nanocomposite SPE approach with incorporation of 2-dimensional (2D) nanofillers has been explored to tailor the crystallization behavior of PEO and the ion conduction of the resultant SPEs. 2D Graphene oxide (GO) nanosheets have been selected as the nanofillers in the present study due to their high aspect ratio, high Young’s modulus and ease of dispersion in common organic solvents. In addition, relatively large quantities of single or a-few-layer GO nanosheets can be easily synthesized through an oxidation−reduction reaction from natural abundant graphite.31−36 Recently a few randomly oriented GO/PEO nanocomposite SPEs have been reported with enhanced ion conductivity at relatively low GO concentrations. It was suggested that the percolated GO network provides a fast conducting pathway.37−40 The GO orientation effect and anisotropic ion transport had not yet been demonstrated in the above-mentioned systems. On the other hand, anisotropic ion conductivity was found previously in PEO/clay SPE systems,41−44 and the conductivity anisotropy can be controlled via magnetic field45 and layer-by-layer (LbL) assembly46 methods. In the present work, the structure, crystallization behavior and ionic conductivity of a GO/PEO nanocomposite SPE containing 10 wt % of GO have been systematically studied. GO nanosheets act as passive fillers, the presence of which slows down the PEO crystallization and imparts significant enhancement of PEO crystal orientation in the resultant SPE membrane as compared with solution cast PEO SPEs without Received: May 6, 2015 Revised: June 9, 2015

A

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Macromolecules the GO fillers. The ion conductivity anisotropy and enhanced in-plane conductivity can be correlated with PEO crystalline structure and chain mobility, GO orientation, and crystallinity of the nanocomposite.

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SPEs were conducted using a PerkinElmer DSC7. Samples of 5−6 mg were hermetically sealed in 40 μL aluminum pans in the glovebox before measurements. Samples were heated/cooled between −60 to 120 °C at 10 °C/min under N2. Simultaneous 2D wide-angle X-ray diffraction (WAXD) and small-angle X-ray scattering (SAXS) experiments were performed using a Rigaku S-MAX 3000 SAXS system. WAXD patterns were collected using an image plate, and the SAXS patterns were collected simultaneously with a 2D multiwire area detector. In-situ heating/cooling WAXD/SAXS experiments were carried out using a Linkam high temperature control stage. Ionic Conductivity Measurements of PEO SPE and GO/PEO Nanocomposite SPE. Room temperature ionic conductivities were measured in an argon-purged glovebox using custom-made cells connected to a Princeton Applied Research Parstat 2273 Potentiostat. In-plane conductivity σ∥ was measured using a four-parallel-electrode method with two outer working electrodes and two inner sensing electrodes lying on the same side of the sample. The ac impedance spectroscopy was collected from 0.1 Hz to 1 MHz at 500 mV. Ionic conductivity was calculated using the following equation:

EXPERIMENTAL SECTION

Materials. Poly(ethylene oxide) (PEO) (average Mn ∼ 300 kDa), N,N′-dimethylformamide (DMF) (anhydrous, 99.8%), lithium perchlorate (LiClO4) (battery grade, dry, 99.99% trace metals basis) were purchased from Sigma-Aldrich. Both preweighed PEO and LiClO4 powders were heated in a vacuum oven at 100 °C overnight right before use to remove moisture. DMF was dried using a 4 Å molecular sieve for 10 min before use. Graphite oxide was synthesized by modified Hummer’s method.47 Detailed synthesis procedure and characterization can be found in ref48. One or a-few-layer GO nanosheets were obtained by sonicating graphite oxide in DMF. Preparation of PEO SPE and GO/PEO Nanocomposite SPE. Both PEO SPEs and GO/PEO SPEs were fabricated using a solution casting method as illustrated in Figure 1. For the pure PEO SPEs, 500

σ =

L w×t×R

(1)

Here L is the distance between two inner sensing electrodes, 2.5 mm in this study, and w and t are the width and thickness of the sample, respectively. The bulk resistance of the electrolyte R is determined from the low frequency plateau in the Bode plot of real impedance z as a function of frequency (the result is indistinguishable with that obtained from the semicircle fit method in a Nyquist plot). Through-plane conductivity σ⊥ was measured using a two-electrode setup with the sample film sandwiched between two stainless steel electrodes. The ac impedance spectroscopy was collected from 0.1 Hz to 1 MHz at 10 mV. The conductivity was calculated using the following equation:

σ⊥ =

L R×A

(2)

where L is the sample thickness, the typical electrolyte membrane thickness range from 100 to 150 μm; A is the area of the electrode and is 0.367 cm2 for the setup used in this study. The bulk resistance R is read as the intersection of the semicircle fit with axis of real impedance part in the Nyquist plot (also consistent with the result obtained from the high frequency plateau in the real impedance plot). Temperaturedependent ionic conductivity was performed using a Mettler Toledo hot stage with a temperature accuracy of less than 0.1 °C.

Figure 1. Illustration of the fabrication and composition of the GO/ PEO nanocomposite electrolyte film. mg PEO was dissolved in 8 mL DMF. LiClO4 powder with [EO] unit to [Li+] molar ratio of 12 was added and the mixture was stirred under nitrogen for 24 h at 45 °C to obtain a homogeneous solution. The solution was then casted in a PTFE lined Petri dish and dried under vacuum at 45 °C for 5 days. The resulting dry film was then transferred into an argon-purged glovebox and stored at room temperature. The fabrication of GO/PEO nanocomposite electrolyte involves two steps. First, 66.75 mg GO was sonicated in 5 mL DMF using a Branson Ultrasonic bath for 90 min to yield a uniform GO dispersion; the solution was then mixed with PEO solution (500 mg in 5 mL DMF) and stirred at 45 °C under nitrogen for 10 h. LiClO4 ([EO]/[Li+] molar ratio 12) was added and stirred for another 5 h before casting into films. The GO content was controlled to be 10 wt % of the total weight of the nanocomposite electrolyte film. The 10 wt % GO/PEO nanocomposite film was also fabricated in a similar way as the nanocomposite SPEs without adding the lithium salt. Here PEO SPE is referred to as P(EO)12:LiClO4 SPE and the GO/PEO nanocomposite SPE as GO0.1-P(EO)12:LiClO4 SPE. Structural and Morphological Characterization. The morphology of GO nanosheets was characterized by transmission electron microscopy (TEM) using a JEOL JEM2100 microscope with an accelerating voltage of 200 kV. Samples were drop-casted onto carbon coated nickel grids and dried overnight. X-ray powder diffraction (XRD) of graphite powder and graphite oxide was conducted using a Siemens D500 diffractometer with a Cu Kα wavelength of 1.54 Å. Samples were scanned from 5° to 50° at a rate of 0.03°/s. The chemical composition of graphite oxide was analyzed using elemental analysis and Karl Fischer Coulometry at Robertson Microlit Laboratories, Inc. (Ledgewood, NJ). Differential scanning calorimetry (DSC) measurements of PEO SPEs and GO/PEO nanocomposite

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RESULTS AND DISCUSSION Effect of GO Nanosheets on the Structure of PEO SPEs. The structures of the synthesized graphitic material and its derivatives were first characterized using powder XRD, in which case the powder samples were pressed onto a glass substrate and the diffraction pattern was collected using a θ−θ scanning setup. The d-spacing of the (002) plane increases from 0.34 nm for graphite to 0.96 nm for graphite oxide based on the XRD results (Figure 2a), suggesting the incorporation of additional functional groups between the graphite galleries after oxidation. The C/O ratio calculated from elemental analysis in combination with Karl Fischer Coulometry is ∼1.3, which indicates high degree of oxidation and is consistent with the expansion of the interlayer distance. Graphite oxide can be dispersed easily in aqueous solution and commonly used polar organic solvents to yield exfoliated single or a-few-layer GO nanosheets. Figure 2b shows a typical TEM micrograph of exfoliated GO nanosheets with 200−500 nm lateral size after sonication in DMF solution. Azimuthal integrations of the 2D WAXD patterns of GO/ PEO nanocomposite with and without lithium salt are also shown in Figure 2a. Two major diffraction peaks at 2θ = 19.15° B

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to a d-spacing of 1.3 nm), indicating the expansion of the graphite oxide galleries due to the intercalation of the PEO chains. The d-spacing of the graphite oxide (002) plane further increases from 1.3 to 2.1 nm (peak position at 2θ = 4.2°) when lithium salt is introduced. Fixed concentration of GO nanosheets, i.e., 10% of the total weight of the PEO electrolyte, was used in this study to prepare nanocomposites and SPEs. The control PEO and GO hybrid SPEs are abbreviated as P(EO)12:LiClO4 and GO0.1-P(EO)12:LiClO4, where the subscript 12 denotes the [EO]/ [Li+] molar ratio, while the superscript 0.1 stands for the GO weight percentage in the SPE. It is anticipated that such a high GO content would impose strong confinement on the structure of the electrolyte membrane, resulting in highly anisotropic ion transport. Because of the 2D geometry and large aspect ratio, GO nanosheets can be readily aligned parallel to the film surface upon controlled evaporation during the film forming process. To characterize the orientation of the nanocomposite SPE films, 2D WAXD was performed from two directions. Here we define z-axis as the film normal, and x and y-axes are parallel to the film plane. An in-plane direction (Figure 3a) refers to the X-ray beam being parallel to the x and y planes, while a through-plane direction (Figure 3d) is for the X-ray beam being parallel to the film normal z. As shown in Figure 3b, the in-plane WAXD pattern of the GO0.1-P(EO)12:LiClO4 nanocomposite SPE reveals that GO nanosheets are strongly oriented parallel to the film surface (low angle meridional

Figure 2. (a) Wide angle X-ray diffraction of graphite powder, graphite oxide, GO/PEO nanocomposite and the corresponding SPE. (b) Transmission electron micrograph of exfoliated GO nanosheets.

and 23.3° are observed in both cases, corresponding to the (120) and (032) planes of the PEO monoclinic crystal structure. The diffraction peak of the GO (002) plane shifts to 2θ = 6.8° of GO/PEO nanocomposite film (corresponding

Figure 3. 2D WAXD of SPEs performed in two directions: in-plane (a) where the incident beam is parallel to film surface and through-plane (d) where incident X-ray beam is perpendicular to film surface; 2D WAXD patterns from the in-plane direction for GO0.1-P(EO)12:LiClO4 nanocomposite SPE (b) and P(EO)12:LiClO4 SPE (c) and from through-plane directions for GO0.1-P(EO)12:LiClO4 nanocomposite SPE (e) and P(EO)12:LiClO4 SPE (f); Azimuthal profiles of PEO (120) diffraction plane for GO0.1-P(EO)12:LiClO4 nanocomposite SPE (h) and P(EO)12:LiClO4 SPE (g). All measurements were performed at room temperature. C

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favor ion conduction. Note that pure PEO has a Tg of −64 °C. The Tg increase in these SPEs is typically attributed to the cross-linking effect of Li+; the higher the concentration of the Li+, the greater the Tg. In the present case, P(EO)12:LiClO4 shows a 13 degree higher Tg than GO0.1-P(EO)12:LiClO4, possibly because of the reduced crystallinity in the latter that results in a higher effective [EO/[Li+] molar ratio. Isothermal crystallization measurements were further performed to investigate the crystallization kinetics. The nanocomposite and PEO films without salt doping were isothermally crystallized at 30 °C for 10 min and have been used as a control study. A well-defined isothermal crystallization peak could not be obtained when the SPEs were crystallized at 30 °C because the crystallization kinetics of all SPE samples were significantly slower, so a lower isothermal crystallization temperature of 28 °C was selected for both SPE samples in this study. The isothermal crystallization curves are plotted in Figure 5. Here

reflection of GO (002) plane) with the Herman’s orientation factor f 002 calculated to be −0.34, while no (002) reflection of GO is observed from the through-plane WAXD pattern in Figure 3e. This suggests that all the GO nanosheets are relatively parallel to the film surface. PEO crystals also exhibit a noticeable preferred orientation, as indicated by a pair of equatorial PEO (120) arcs (azimuthal profile of PEO (120) plane is shown in Figure 3h); the Herman’s orientation factor of PEO (120) plane f120 can be calculated to be ∼0.64. An isotropic (120) ring is observed in the through-plane pattern, indicating the c-axis of the polymer chain is aligned perpendicular to the film surface. Detailed calculation of Herman’s factor for GO nanosheets and PEO crystal orientation can be found in our previous studies.24,49 As a control, a P(EO)12:LiClO4 SPE without GO was prepared via a similar solution casting method and the structural characterization is shown in Figure 3, parts c, f, and g. The PEO chains also exhibit moderate alignment, while the degree of PEO lamella orientation is considerably lower than that of its GO/ PEO nanocomposite SPE counterpart (Herman’s factor f120 calculated to be ∼0.35 compared with 0.64 for nanocomposite SPE). The 2D WAXD results of the two SPEs suggest that a highly anisotropic electrolyte structure can be obtained using a simple solution casting method. The well−oriented GO nanosheets serve as template to confine the chain alignment of PEO lamella. Crystallization Behaviors of the Nanocomposite SPE. The ion transport behaviors of the nanocomposite SPEs are complicated by PEO crystallization. Both nonisothermal and isothermal DSC scans were performed to better understand PEO crystallization and its correlation to ion transport behaviors of the SPEs. The nonisothermal scans (from first heating) of the as cast SPEs that had been stored at room temperature for 10 days are shown in Figure 4. Both SPEs

Figure 5. (a) Isothermal crystallization at 30 °C of GO/PEO nanocomposite and PEO films without lithium salt doping. (b) Isothermal crystallization at 28 °C of GO/PEO nanocomposite and PEO SPEs with LiClO4 doping.

t0.05 and t1/2 are defined as the time taken to reach 5% and 50% of the final crystallinity, respectively, and the values are listed in Table 1. Two different trends have been observed. First, for the systems without lithium salt doping, t1/2 decreases from 0.21 min for the pure PEO film to 0.12 min for the nanocomposite, suggesting the crystallization kinetics is faster in the nanocomposite case. This is because that the GO nanosheets act as the heterogeneous nuclei for PEO crystallization. Second, an opposite trend has been observed in the corresponding SPE systems. The nanocomposite SPE has a slower crystallization kinetics compared with PEO SPEs, and it is consistent with the reduced crystallinity shown in Figure 4. An Avrami equation50 is typically employed to describe the crystallization process and can be expressed as follows:

Figure 4. DSC first heating scans for as cast P(EO)12:LiClO4 and GO0.1-P(EO)12:LiClO4 SPEs that had been stored at room temperature for 10 days.

exhibit broad melting peaks (30−65 °C for P(EO)12:LiClO4 and 30−55 °C for GO0.1-P(EO)12:LiClO4) with small shoulders, indicating a broad distribution of PEO lamellar thickness as a result of the strong PEO−Li interactions during the crystallization process. GO0.1-P(EO)12:LiClO4 has a lower melting temperature (50 °C of the peak position compared with 56 °C for the as cast P(EO)12:LiClO4) and reduced crystallinity (22.9% compared with 34.9% for as cast P(EO)12:LiClO4), suggesting a suppression of crystallization in the presence of GO nanoplatelets. In addition, the glass transition temperature (Tg) of the nanocomposite SPE drops to −40 °C compared with −27 °C for the pure PEO SPE, which implies a higher room temperature chain mobility that would

1 − X(t ) = exp( −Kt n)

(3)

where X(t) is the relative crystallinity, calculated as the ratio of the heat of fusion at time t and the total heat of fusion of the whole crystallization process, and is plotted in Figure 6a as a D

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Macromolecules Table 1. Summary of crystallization parameters obtained from DSC isothermal crystallization 0.1

GO -PEO PEO GO0.1-P(EO)12-LiClO4 P(EO)12LiClO4

Tc (°C)

t0.05 (min)

t1/2 (min)

n

K (min−n)

30 30 28 28

0.07 0.14 0.19 0.12

0.12 0.21 0.84 0.49

2.75 2.54 1.67 1.74

502.3 177.5 0.86 2.45

note no salt salt doping

crystallization kinetics. As previously discussed, the formation of highly concentrated Li+ layers on GO surface could lead to a lower ionic cross-linking of the PEO chain, which further explains the observed 13 deg Tg drop in the nanocomposite SPE. The Avrami exponent n typically indicates the growth dimension of a polymer crystal; however, it should be noted that in addition to growth dimensionality, the exponent also depends on many other factors. For example, 3D growth of athermal nucleation leads to an exponent of 3 while similar 3D growth of thermal nucleation has an exponent of 4. Other factors that complicate the situation include volume change during crystallization, changing growth rates during crystal growth, changing of the nucleation mechanism due to asymmetric nucleation agent, etc. Therefore, caution should be taken when applying Avrami analysis in polymer crystallization. The Avrami exponent n was calculated to be 2.75 and 2.54 for GO/PEO nanocomposite and PEO film without salt doping, suggesting a nearly 3D spherulitic growth. On the other hand, the calculated Avrami exponent n values for both nanocomposite and PEO SPEs fell below 2 (1.67 for nanocomposite SPE and 1.74 for PEO SPE), which is possibly due to the strong interaction between PEO chain and Li+ during the crystallization process. As the PEO starts to crystallize, the lithium ions are expelled from densely packed crystals and complex with polymer chains in the remaining amorphous region. The effective Li+ to ether oxygen ratio gradually increases as the crystallinity increases, which would result in a stronger restriction of the mobility of the amorphous chains that need to diffuse to the growth front. This could possibly change the linear growth rate of the crystal, leading to the Avrami exponent deviation. Correlation of Structural Anisotropy and Ion Transport in GO/PEO Nanocomposite SPE. The conducting behavior of GO nanoplatelets was first evaluated by EIS measurements. CV scans were performed on a GO thin film and the results suggested that the material is electrically insulating. An ac scan was also performed on the same GO film using two parallel stainless steel electrodes. The real impedance value approached the detection limit of our instrument at low frequency, and the ac conductivity was calculated to be below 4 × 10−12 S/cm. The results suggest that GO nanoplatelets are neither electrical nor ionic conductive and therefore can be used as a passive filler in the nanocomposite electrolyte. Temperature dependence of σ∥ and σ⊥ for P(EO)12:LiClO4 and GO0.1-P(EO)12:LiClO4 have been investigated, and Figure 7 shows the first heating curves at 1 °C/min for as cast SPEs that had been stored at room temperature for 10 days. Each curve was repeated three times and the standard deviations are shown in the figure. On the basis of previous DSC data (Figure 4), the melting ranges for the SPEs are 30−65 °C for P(EO)12:LiClO4 and 30−55 °C for GO0.1-P(EO)12:LiClO4, respectively. Therefore, we shall use two temperature regions to discuss the conductivity behavior: high temperature region where T > 65 °C, and low temperature region near room

Figure 6. (a) Relative crystallinity X(t) as a function of isothermal crystallization time t. (b) Avrami plots of GO/PEO nanocomposite at 30 °C (green open square), pure PEO at 30 °C (blue open triangle), P(EO)12:LiClO4 SPE at 28 °C (red solid triangle) and GO0.1P(EO)12:LiClO4 SPE at 28 °C (black solid square).

function of t, n is the Avrami exponent, and K is the crystallization rate parameter. Values of n and K are determined using the initial linear part of the Avrami plot (Figure 6b) and are listed in Table 1. The kinetics parameter K is reduced significantly for all SPEs compared with the pristine nanocomposite and PEO without salt doping, which can be attributed to the strong interaction between the salt and the polymer. K of the as cast PEO SPE is nearly as 2 times greater than that of the nanocomposite SPE, which is consistent with the t1/2 result. The retardation of PEO crystallization in multiphase systems is not unusual. Similar phenomenon had been observed in a PEO nanocomposite containing lithium dodecyl sulfate (LDS) stabilized single wall carbon nanotubes (SWNTs).51−53 It has been proposed that the presence of Li+ and SWNTs have a synergistic effect on the disruption of PEO crystallization. Specifically, Li+ ions have an affinity to SWNTs and strongly interact with PEO, resulting in a localized region of an amorphous PEO−Li complex phase near SWNTs. The nucleation and growth of PEO are further disrupted by the SWNT network that imposes an energy barrier for chain diffusion during the crystallization process. In the present GO/ PEO nanocomposite SPE, Li+ ions may be enriched in the vicinity of GO surface due to the relatively strong interaction between Li+ and the polar functional groups on GO surface. This surface Li+ layer can further bind with PEO, leading to only an amorphous PEO/Li+ complex phase. This process will significantly reduce the number of available nucleation sites for PEO heterogeneous crystallization, leading to a decrease of the E

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Figure 7. Temperature dependent ionic conductivity for as cast SPEs that stored at room temperature for 10 days. All curves represent first heating scans from room temperature to 100 °C at 1 °C/min: blue solid triangle, P(EO)12:LiClO4 in-plane; black open square, GO0.1P(EO)12 :LiClO 4 in-plane; red solid diamond, P(EO)12 :LiClO4 through-plane; green open circle, GO0.1-P(EO)12:LiClO4 throughplane.

temperature. In the high temperature region, all the polymer crystals are molten and there is no crystallization effect on the observed ion conductivity, while in the low temperature region, crystallization plays a significant role. We begin our discussion with the high temperature region. Three features can be observed: (i) σ∥ and σ⊥ curves (blue and red) for P(EO)12:LiClO4 are overlapped; (ii) σ∥ and σ⊥ curves (black and green) for GO0.1-P(EO)12:LiClO4 show appreciable differences, with an anisotropic factor (σ∥/σ⊥) of 4 (Figure 8a,b); (iii) σ∥ curves for samples with and without GO are almost superimposed (blue and black), while σ⊥ for GO− containing samples (green) are noticeable lower than the control (red). Overlapping of σ∥ and σ⊥ curves for P(EO)12:LiClO4 at the high temperature region suggests that the SPE is in an isotropic state, and there is no residual chain orientation after PEO crystal melts. This can be supported by the 2D WAXD experiments: the diffraction patterns from both in-plane and through-plane directions exhibit isotropic amorphous halo in the wide angle range (Figure 8c, bottom panels). In GO0.1-P(EO)12:LiClO4, PEO chains are also molten with an isotropic conformation, suggested by the 2D WAXD patterns (Figure 8c, top panels). The in-plane diffraction pattern (Figure 8c, top left) shows strong GO diffractions in the meridian, suggesting the GO orientation retains after the melting of PEO crystals (azimuthal profiles of GO(002) diffraction plane measured at 100 °C are shown in Figure S1 in the Supporting Information). This GO orientation therefore accounts for the observed conductivity anisotropy factor of ∼4 of GO0.1-P(EO)12:LiClO4 in the high temperature region. It is of interest that the in-plane conductivity profile of GO0.1P(EO)12:LiClO4 overlaps with both in- and through-plane conductivities of P(EO)12:LiClO4, suggesting no significant change of ion conductivity along the surfaces of GO nanosheets at high temperatures. Gao et al. has reported a 10 times enhancement of conductivity above Tm of PEO for a random GO/PEO nanocomposite at 0.6 wt % GO loading, in which case the enhancement was attributed to the conducting ability of the layered GO/Li +/EO/Li +/GO composite.37 The enhancement has not yet been observed in our system at high temperature. One notable difference in our SPE is that despite the GO content being as high as 10%, the GO nanosheets are intercalated further with a more open structure

Figure 8. (a) Conductivity anisotropy as a function of temperature calculated from the first heating scan; (b) zoom-in plot of part a at temperatures above PEO Tm; (c) Corresponding 2D WAXD patterns of SPEs from both in-plane and through-plane directions at 100 °C.

(the d-spacing of GO (002) plane is 2.1 nm in our 10% aligned nanocomposite SPE compared with 1.4 nm in the reported 0.6% random nanocomposite SPE). Detailed study on the correlation between the GO gallery distance and the resultant conductivity is ongoing. Note that the conductivity enhancement due to incorporating nanoparticles is not ubiquitous and contradictory results have been found in other SPEs. Best et al. found no improvement, and even a decrease of conductivity in some fully amorphous polyether based SPE with the addition of TiO2 or Al2O3 nanoparticles.54 Johansson et al. reported no significant influence of the SiO2 fillers on the amorphous PEO− LiTFSI SPE.55 Xie et al. studied a PEO−LiTFSI SPE containing fumed silica nanoparticles, and conductivity of the composite SPE was found to decrease above Tm when compared with ultrapure PEO SPE.56 It is apparent that more detailed studies are needed to better understand the nanoparticle effect on SPE conductivity. Nevertheless, the conductivity anisotropy induced by 2D fillers has been clearly demonstrated. We can then move to the discussion on the low temperature conductivity profiles in Figure 7. Three features can also be observed. (i) For P(EO)12:LiClO4, σ∥ is noticeably greater that σ⊥ with an anisotropy factor of ∼10 at room temperature. (ii) F

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mobility. Specifically, the higher σ∥ of GO0.1-P(EO)12:LiClO4 compared with P(EO)12:LiClO4 is due to the enhanced ion conducting pathway combined with lower PEO crystallinity and high segmental mobility (lower Tg). On the other hand, the addition of GO imposes tortuous pathway for ion transport along the through plane direction in GO0.1-P(EO)12:LiClO4, which competes with crystallinity reduction and chain mobility enhancement, resulting in similar σ⊥ compared with PEO SPE. Table 2 shows the ionic conductivity normalized by both Tg and the volume fraction of conducting phase to account for the crystallinity and chain mobility difference between GOcontaining SPE and pure PEO SPE. The normalized σ∥ for GO0.1-P(EO)12:LiClO4 is 2.6 times higher than that of P(EO)12:LiClO4 while the normalized σ⊥ is 4.5 times lower. The enhanced in-plane ion transport and impeded throughplane ion transport can be attributed to geometric confinement only.

For GO0.1-P(EO)12:LiClO4, σ∥ is also greater that σ⊥; however, the anisotropy factor is much greater, reaching as high as 70 at room temperature. (iii) σ∥ for GO0.1-P(EO)12:LiClO4 (black) is nearly 1 order of magnitude higher than that of P(EO)12:LiClO4 (blue), while the σ⊥ values for both SPEs (green and red) are similar at room temperature. The low temperature conductivity anisotropy arises from the combined effect of crystallization and GO orientation as shown in Figure 9. Our previous work showed that in a well-aligned

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CONCLUSIONS In summary, a GO/PEO nanocomposite SPE has been prepared by solution casting a homogeneous mixture of PEO, GO and LiClO4. GO was highly aligned with the nanoplatelet surface parallel to the film surface during a slow solvent evaporation process, which further confined PEO crystallization, resulting in the polymer chain perpendicular to the film surface (PEO crystalline lamellae surfaces parallel to the film surfaces). The presence of GO and Li+ ions had a synergistic effect of confining PEO crystal orientation and retarding PEO crystallization. The ion transport is guided by GO nanoplatelets and PEO lamellae, leading to highly anisotropic ionic conductivity in both SPEs. In particular, conductivity anisotropy factors as high as ∼70 have been achieved in the nanocomposite SPE. This study demonstrated that PEO crystallization can be tuned and controlled using 2D templates, furthering our understanding of the complex interactions during ion transport at the fundamental level, which can help guide engineering new and improved PEO-incorporated batteries.

Figure 9. Illustration of the crystalline structure change of P(EO)12:LiClO4 (top) and GO0.1-P(EO)12:LiClO4 (bottom) SPEs from room temperature (left) to 100 °C (right).

polymer single crystal SPE, an anisotropy factor of ∼1000 could be achieved, which was attributed to the tortuous ion pathway imposed by the 2D lamellar crystals. Since the dissolved polymer was directly casted onto the substrate from the mixture solution without the controlled single crystal growth step in the present case of P(EO)12:LiClO4 SPE, the crystals are much less orientated compared with the single crystal model systems. Nevertheless, the 2D WAXD pattern in Figure 3 does show a moderate orientation of PEO crystals, with a f120 of ∼0.35. Such orientation therefore leads to an observed anisotropy factor of ∼10 at room temperature as shown in Figure 8a. For GO0.1P(EO)12:LiClO4, a much greater anisotropy factor was observed (70 at room temperature). This apparently is due to the enhanced lamellar orientation and GO orientation. As previously discussed, the f120 of GO0.1-P(EO)12:LiClO4 was calculated to be ∼0.64, much higher than that of P(EO)12:LiClO4. It is also worth mentioning that this enhanced orientation of PEO lamellar is due to the GO templating effect. Another noticeable feature of the ionic conductivities below PEO Tm is that σ∥ of GO0.1-P(EO)12:LiClO4 is ∼1 order of magnitude higher than that of P(EO)12:LiClO4 while the σ⊥ values of these two SPEs at room temperature are identical. The low temperature conductivity behaviors for GO-containing SPE and PEO SPE result from the competition among geometric confinement, PEO crystallinity and PEO chain

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ASSOCIATED CONTENT

* Supporting Information S

Azimuthal profiles of GO (002) diffraction plane measured at 100 °C. The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/ acs.macromol.5b00972.

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AUTHOR INFORMATION

Corresponding Author

*(C.Y.L.) E-mail: [email protected]. Telephone: 215-8952083. Fax: 215-895-6760. Notes

The authors declare no competing financial interests.

Table 2. Ionic Conductivity Normalized by Tg and the Volume Fraction of the Conducting Phase |σ| at (T − Tg) = 67 °C (S/cm) GO0.1-P(EO)12LiClO4 SPE P(EO)12LiClO4 SPE

Tg (°C)

T (°C)

in-plane

through-plane

−40 −27

27 40

6.46 × 10−07 2.47 × 10−07

9.22 × 10−09 4.15 × 10−08

Note: |σ| refers to ionic conductivity normalized by the GO volume fraction ϕGO and the crystallinity Xc, where |σ| = σ/(1 − ϕGO)/(1 − Xc). G

DOI: 10.1021/acs.macromol.5b00972 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules

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(33) Dikin, D. A.; Stankovich, S.; Zimney, E. J.; Piner, R. D.; Dommett, G. H. B.; Evmenenko, G.; Nguyen, S. T.; Ruoff, R. S. Nature 2007, 448, 457. (34) Stankovich, S.; Dikin, D. A.; Piner, R. D.; Kohlhaas, K. A.; Kleinhammes, A.; Jia, Y.; Wu, Y.; Nguyen, S. T.; Ruoff, R. S. Carbon 2007, 45, 1558. (35) Park, S.; An, J. H.; Jung, I. W.; Piner, R. D.; An, S. J.; Li, X. S.; Velamakanni, A.; Ruoff, R. S. Nano Lett. 2009, 9, 1593. (36) Zhu, Y. W.; Murali, S.; Cai, W. W.; Li, X. S.; Suk, J. W.; Potts, J. R.; Ruoff, R. S. Adv. Mater. 2010, 22, 3906. (37) Gao, S.; Zhong, J.; Xue, G.; Wang, B. J. Membr. Sci. 2014, 470, 316. (38) Yuan, M.; Erdman, J.; Tang, C.; Ardebili, H. RSC Adv. 2014, 4, 59637. (39) Cao, Y.-C.; Xu, C.; Wu, X.; Wang, X.; Xing, L.; Scott, K. J. Power Sources 2011, 196, 8377. (40) Shim, J.; Kim, D.-G.; Kim, H. J.; Lee, J. H.; Baik, J.-H.; Lee, J.-C. J. Mater. Chem. A 2014, 2, 13873. (41) Ruizhitzky, E.; Aranda, P. Adv. Mater. 1990, 2, 545. (42) Aranda, P.; Galvan, J. C.; Casal, B.; Ruizhitzky, E. Electrochim. Acta 1992, 37, 1573. (43) Hutchison, J. C.; Bissessur, R.; Shriver, D. F. Chem. Mater. 1996, 8, 1597. (44) Elmahdy, M. M.; Chrissopoulou, K.; Afratis, A.; Floudas, G.; Anastasiadis, S. H. Macromolecules 2006, 39, 5170. (45) Kitajima, S.; Matsuda, M.; Yamato, M.; Tominaga, Y. Polym. J. 2013, 45, 738. (46) Lutkenhaus, J. L.; Olivetti, E. A.; Verploegen, E. A.; Cord, B. M.; Sadoway, D. R.; Hammond, P. T. Langmuir 2007, 23, 8515. (47) Hummers, W. S.; Offerman, R. E. J. Am. Chem. Soc. 1958, 80, 1339. (48) Cheng, S.; Chen, X.; Hsuan, Y. G.; Li, C. Y. Macromolecules 2011, 45, 993. (49) Cheng, S.; Cairncross, R. A.; Hsuan, Y. G.; Li, C. Y. Polymer 2013, 54, 5016. (50) Avarami, M. J. Chem. Phys. 1939, 7, 1103. (51) Chatterjee, T.; Lorenzo, A. T.; Krishnamoorti, R. Polymer 2011, 52, 4938. (52) Edman, L.; Ferry, A.; Jacobsson, P. Macromolecules 1999, 32, 4130. (53) Chatterjee, T.; Yurekli, K.; Hadjiev, V. G.; Krishnamoorti, R. Adv. Funct. Mater. 2005, 15, 1832. (54) Best, A. S.; Adebahr, J.; Jacobsson, P.; MacFarlane, D. R.; Forsyth, M. Macromolecules 2001, 34, 4549. (55) Johansson, P.; Ratner, M. A.; Shriver, D. F. J. Phys. Chem. B 2001, 105, 9016. (56) Xie, J.; Duan, R. G.; Han, Y.; Kerr, J. B. Solid State Ionics 2004, 175, 755.

ACKNOWLEDGMENTS We are grateful for the support from the National Science Foundation through Grants DMR-1308958, CBET-1510092, and CMMI-1334067. The Rigaku S-MAX 3000 SAXS system was purchased through Grant NSF MRI-1040166.

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REFERENCES

(1) Armand, M. Solid State Ionics 1983, 9−10 (Part 2), 745. (2) Armand, M. B. Annu. Rev. Mater. Sci. 1986, 16, 245. (3) Armand, M. Adv. Mater. 1990, 2, 278. (4) Hickner, M. A. Mater. Today 2010, 13, 34. (5) Di Noto, V.; Lavina, S.; Giffin, G. A.; Negro, E.; Scrosati, B. Electrochim. Acta 2011, 57, 4. (6) Tarascon, J. M.; Armand, M. Nature 2001, 414, 359. (7) Quartarone, E.; Mustarelli, P. Chem. Soc. Rev. 2011, 40, 2525. (8) Gray, F. M.; MacCallum, J. R.; Vincent, C. A.; Giles, J. R. M. Macromolecules 1988, 21, 392. (9) Ruzette, A.-V. G.; Soo, P. P.; Sadoway, D. R.; Mayes, A. M. J. Electrochem. Soci. 2001, 148, A537. (10) Niitani, T.; Shimada, M.; Kawamura, K.; Kanamura, K. J. Power Sources 2005, 146, 386. (11) Singh, M.; Odusanya, O.; Wilmes, G. M.; Eitouni, H. B.; Gomez, E. D.; Patel, A. J.; Chen, V. L.; Park, M. J.; Fragouli, P.; Iatrou, H.; Hadjichristidis, N.; Cookson, D.; Balsara, N. P. Macromolecules 2007, 40, 4578. (12) Panday, A.; Mullin, S.; Gomez, E. D.; Wanakule, N.; Chen, V. L.; Hexemer, A.; Pople, J.; Balsara, N. P. Macromolecules 2009, 42, 4632. (13) Young, W. S.; Kuan, W. F.; Epps, T. H. J. Polym. Sci., Polym. Phys. 2014, 52, 1. (14) Cheng, S.; Smith, D.; Pan, Q. W.; Wang, S. J.; Li, C. Y. RSC Adv. 2015, 5, 48793. (15) Wang, Y.; Travas-Sejdic, J.; Steiner, R. Solid State Ionics 2002, 148, 443. (16) Zhang, S. S.; Xu, K.; Foster, D. L.; Ervin, M. H.; Jow, T. R. J. Power Sources 2004, 125, 114. (17) Smith, D. M.; Dong, B.; Marron, R. W.; Birnkrant, M. J.; Elabd, Y. A.; Natarajan, L. V.; Tondiglia, V. P.; Bunning, T. J.; Li, C. Y. Nano Lett. 2011, 12, 310. (18) Smith, D. M.; Li, C. Y.; Bunning, T. J. J. Polym. Sci., Polym. Phys. 2014, 52, 232. (19) Smith, D. M.; Cheng, S.; Wang, W. D.; Bunning, T. J.; Li, C. Y. J. Power Sources 2014, 271, 597. (20) Croce, F.; Appetecchi, G. B.; Persi, L.; Scrosati, B. Nature 1998, 394, 456. (21) Croce, F.; Curini, R.; Martinelli, A.; Persi, L.; Ronci, F.; Scrosati, B.; Caminiti, R. J. Phys. Chem. B 1999, 103, 10632. (22) Appetecchi, G. B.; Croce, F.; Persi, L.; Ronci, F.; Scrosati, B. Electrochim. Acta 2000, 45, 1481. (23) Croce, F.; Persi, L.; Scrosati, B.; Serraino-Fiory, F.; Plichta, E.; Hendrickson, M. A. Electrochim. Acta 2001, 46, 2457. (24) Cheng, S.; Smith, D. M.; Li, C. Y. Macromolecules 2014, 47, 3978. (25) Li, C. Y. J. Polym. Sci., Polym. Phys. 2009, 47, 2436. (26) Zhou, T.; Dong, B.; Qi, H.; Mei, S.; Li, C. Y. J. Polym. Sci., Polym. Phys. 2014, 52, 1620. (27) Wang, W.; Huang, Z.; Laird, E. D.; Wang, S.; Li, C. Y. Polymer 2015, 59, 1. (28) Wang, W.; Li, C. Y. ACS Macro Lett. 2014, 3, 175. (29) Laird, E. D.; Li, C. Y. Macromolecules 2013, 46, 2877. (30) Dillon, D. R.; Tenneti, K. K.; Li, C. Y.; Ko, F. K.; Sics, I.; Hsiao, B. S. Polymer 2006, 47, 1678. (31) Hummers, W. S.; Offerman, R. E. J. Am. Chem. Soc. 1958, 80, 1339. (32) Stankovich, S.; Dikin, D. A.; Dommett, G. H. B.; Kohlhaas, K. M.; Zimney, E. J.; Stach, E. A.; Piner, R. D.; Nguyen, S. T.; Ruoff, R. S. Nature 2006, 442, 282. H

DOI: 10.1021/acs.macromol.5b00972 Macromolecules XXXX, XXX, XXX−XXX