Effect of Polymer Architecture on the Ionic Conductivity. Densely

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Effect of Polymer Architecture on the Ionic Conductivity. Densely Grafted Poly(ethylene oxide) Brushes Doped with LiTf George Zardalidis,*,† Achilleas Pipertzis,† Grigoris Mountrichas,‡ Stergios Pispas,‡ Markus Mezger,§ and George Floudas*,† †

Department of Physics, University of Ioannina, P.O. Box 1186, 451 10 Ioannina, Greece Theoretical and Physical Chemistry Institute, National Hellenic Research Foundation, 116 35 Athens, Greece § Institute of Physics and Max Planck Institute for Polymer Research, Johannes Gutenberg University Mainz, 55128 Mainz, Germany ‡

S Supporting Information *

ABSTRACT: Densely grafted poly(ethylene oxide) (PEO) brushes on a poly(hydroxylstyrene) (PHOS) backbone (PHOS-g-PEO) as well as block copolymers with polystyrene (PS) (PS-b-(PHOS-g-PEO)) are designed as model systems for Li ion transport. This macromolecular design suppresses the propensity of PEO chains for complex crystal formation with LiTf as well as for crystallization. Li ion conductivities similar or even exceeding those in the archetypal electrolyte poly(ethylene oxide)/lithium triflate (PEO/LiCF3SO3 (LiTf)) are obtained for a range of temperatures and LiTf compositions. At the same time, PHOS-g-PEO and PS-b-(PHOS-g-PEO) show improved mechanical stability. Typically, at 333 K, the ionic conductivity is ∼6 × 10−5 S/cm and the modulus at ∼2 × 106 Pa for a [EO]:[Li+] = 8:1 composition. In the endeavor for suitable solid polymer electrolytes macromolecular architecture seems to play a decisive role.

1. INTRODUCTION Undoubtedly there has been a significant improvement in Li ion energy storage systems over the past three decades. Advancement of material engineering and processing has contributed to an almost tripling of the stored energy density.1 However, in comparison to another major technological field, the integrated circuit, integration density has increased by 3 orders of magnitude in the same time span; thus, the former number seems quite small. This fact raises a question with respect to the disproportional advancement in the two technologies that go hand-by-hand almost in every mobile electronic device or system. The answer can be found in the complexity of structures and behaviors found in the electrolytes as well as in the electrodes employed for Li-ion storage. Concerning the electrolytes, they must fulfill a number of requirements in order to produce true solid polymer electrolytes (SPEs).2−7 In addition, despite recent progress, the lack of a full understanding of the mechanisms responsible for ionic conduction prevented the technology to mature and precluded the mass production of batteries based on SPEs. Starting from the requirement of a polymer that is effective in dissolving metallic salts, poly(ethylene oxide) (PEO) turned out to be one of the best candidates. PEO is a flexible polymer with high solvating capability for a number of lithium salts and at the same time is biodegradable8 and environmentally friendly. However, PEO is lacking the necessary mechanical stability and high elastic modulus necessary to prevent dendritic growth at the electrodes.3 Hence, PEO needs to be “combined” with some stiffer macromolecules that could provide the © XXXX American Chemical Society

required elastic modulus, applicable to SPEs. This fact has initiated a pursue for effective scaffolds for PEO electrolytes based on more stiff macromolecules. In this respect, polymer electrolytes based on block copolymers offer particular advantages due to the simultaneous existence of soft and hard nanophases.9−11 The soft phase is utilized for ion transport while the hard phase provides the necessary mechanical stability. Based on these properties, major research has been conducted on poly(styrene-b-ethylene oxide) (PS-bPEO) copolymers doped with lithium salts.12−26 However, despite the molecular simplicity of block copolymers, there are several microscopic and macroscopic parameters that can affect the mechanism of ion transport. On a microscopic level, nanodomain spacing and interfacial thickness have been considered.18,22,25,27−29 By comparing different copolymers with variable interaction parameters but in the presence of the same conductive phase (PEO) and the same electrolyte, it was recently concluded that it is the nanodomain spacing that exerts a stronger influence on ion transport.29 On a macroscopic level, the details of nanodomain morphology with respect to lamellar orientation, grain size, grain boundaries, and electrode wetting effects adds even more to the complexity.3,17,20,24,30−33 An ideal SPE should be composed of rather short PEO chains of high segmental mobility and at the same time of suppressed tendency for PEO crystallization as well as for Received: February 6, 2016 Revised: March 19, 2016

A

DOI: 10.1021/acs.macromol.6b00290 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Table 1. Molecular Characteristics of the Copolymers

a

sample

Mw,grafta (g/mol) × 106

Mw,backboneb (g/mol) × 104

Nbranches PEO

Mw,PEO branchc × 104

Mw/Mn,graftb

PEOc (wt %)

PHOS82-g-PEO322 PS256-b-(PHOS179-g-PEO320)

1.17 2.57

0.99 4.83

82 179

1.42 1.41

1.18 1.14

99 98

By light scattering in THF. bBy SEC in THF. cBy 1H NMR (and the known characteristics of the backbone).

crystalline complex formation. It is known that the latter two structures, which are always present in SPEs composed of linear PEO, play a dominant role in ion transport. For example, the crystalline complex melts over a range of temperatures and in this process effectively provides additional ions for transport. However, high PEO crystallinity and high propensity for complex crystal formation, in general, impend ion transport.34 Therefore, attempts have been made to suppress complex crystal formation and PEO crystallization by chain architecture.35−37 Zhang et al.36 reported on electrolytes of grafted PEO brush copolymers using poly(ethylene glycol) methyl ether methacrylate as the macromonomer, enriched with LiCF3SO3. They obtained amorphous and homogeneous materials with high ionic conductivity (σ ∼ 10−3 S/cm at [EO]:[Li+] = 10:1) that did not flow at room temperature constituting mechanically soft elastic materials. More recently, triblock copolymers of poly[(norbornene-g-styrene)-b-(norbornene-g-ethylene oxide)-b-(norbornene-g-styrene)] with a brush architecture were found to be promising materials as SPEs. When blended with lithium bis(trifluoromethane)sulfonimide (LiTFSI) they exhibit Li conductivity of 10−3 S/cm at 378 K and a modulus of ∼104 Pa.37 Exploring the most suitable polymer architecture for Li ion transport, we designed short PEO chains densely grafted to a backbone composed of PS. This particular architecture provides several advantages. Densely grafted PEO chains provide the usual ion-transport phase composed of several chain ends that lower the glass temperature and increase the PEO and Li ion mobility. In addition, the PS backbone provides the required rigidity. A possible advantage from such an architecture stems from a reduced propensity for PEO crystallization and complex crystal formation. The latter can be anticipated by the restrictions provided by the dense grafting of PEO chains. In the present study we explore the dc−conductivity in relation to the local and global structure as well as the pertinent viscoelastic properties of two systems composed of grafted poly(hydroxystyrene) (PHOS) with poly(ethylene oxide) (PEO) and of its block copolymer with polystyrene (PS). We effectively demonstrate, for the first time, that densely grafted PEO chains to a PHOS backbone can produce Li ion conductivities that are similar to or even higher than in the respective PEO/LiTf electrolytes. At the same time PS-b(PHOS-g-PEO)/LiTf electrolytes show improved mechanical properties.

Figure 1. Structure of grafted poly(hydroxystyrene) (PHOS) with poly(ethylene oxide) (PEO) (left) and of its block copolymer with polystyrene (PS) (right). sealed ampules. The phosphazene base, t-BuP4 (solution in n-hexane), was subjected to drying overnight by evaporation of hexane under vacuum and redissolution in purified benzene. The synthesis procedure of PHOS82, via conventional anionic polymerization of ptert-butoxystyrene (t-BOS) and postpolymerization hydrolysis, has been described before.2 The synthesis of the PS256-b-PHOS179 diblock was realized similarly to the homopolymer by sequential polymerization of PS and t-BOS in THF, followed by acidic hydrolysis.39 The molecular weight polydispersity (Mw/Mn) of the homopolymer and the block copolymer was 1.09 and 1.19, respectively. Before each graft copolymer synthesis, the appropriate amount of the homopolymer or the block copolymer sample was purified by azeotropic distillation with toluene, left under high vacuum for 24 h, and finally dissolved in dry THF by the aid of cryo-distillation. Graft copolymers were synthesized via sequential metal-free anionic polymerization of EO in THF in the presence of t-BuP4 using PHOS as the backbone.40 Purified solvent (dry THF, ∼30 mL/70 mL for the block copolymer) was distilled into the polymerization apparatus. The apparatus was degassed and flamesealed. A predetermined amount of t-BuP4 solution (0.9 × 10−5 mol/ mL in benzene, 4.63 mL/8.93 mL for the block copolymer) was first added. Subsequently, temperature was decreased to 203 K, and the THF solution of PHOS (0.05 g in ∼10 mL of THF/0.2157 g for the block copolymer) was added via a break-seal, upon which the mixture became somewhat turbid because of the poor solubility of the consequently generated polyanion. This solution was stirred for 10 min at 203 K before EO (5.9 g/11.5 g for the block copolymer) was introduced by cryo-distillation. The reaction mixture was stirred at 203 K for 2 h and then slowly heated to 318 K. The turbidity disappeared gradually at this temperature, indicating the growth of PEO side chains on the backbone. The reaction mixture was stirred at this temperature for 96 h to ensure complete consumption of EO. Finally, the polymerization was terminated by degassed methanol (∼1 mL) and few drops of concentrated hydrochloric acid (HCl). After polymerization, the solvent was evaporated in a rotary evaporator. The crude product was dried in vacuum overnight and weighed in order to determine the conversion of the monomers, which approached 100% in both cases. The crude product was purified by at least two fractionations, using chloroform as the solvent and n-hexane as the precipitant, in order to remove reaction byproducts, especially the linear polymers. 2.2. Preparation of Copolymer/LiTF Composites. Initially, a polymeric solution in THF (5% w/v) was prepared, typically by heating at 333 K in order to ensure complete dissolution. A solution of the LiTf salt in THF (5% w/v) is also prepared. Subsequently, mixing

2. EXPERIMENTAL SECTION 2.1. Materials. The molecular characteristics of the copolymers are summarized in Table 1. Figure 1 gives the chemical structures of the grafted poly(hydroxystyrene) (PHOS) with poly(ethylene oxide) (PEO) (left) and of its block copolymer with polystyrene (PS) (PS-b(PHOS-g-PEO)). A PEO homopolymer of Mw = 1.4 × 104 g/mol and Mw/Mn = 1.06 was used for comparison. All reagents were purchased from Aldrich and purified by wellestablished procedures for anionic polymerization high-vacuum techniques.38 Tetrahydrofuran (THF) was initially dried over CaH2 and then over Na/K alloy. Ethylene oxide (EO) was stirred sequentially with CaH2 and n-butyllithium and then distilled into B

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Macromolecules of the two solutions took place in the appropriate ratio, followed by evaporation of the solvent. It is important that the evaporation is realized in slightly increased temperature (∼303 K) in order to avoid crystallization of PEO. 2.3. Characterization Methods. Differential Scanning Calorimetry. The thermal properties of the electrolytes were studied with a Q2000 (TA Instruments) differential scanning calorimeter (DSC). Cooling and heating cycles were performed at a rate of 10 K/min and in a temperature range between 173 and 453 K. The instrument was calibrated for best performance on the specific temperature range and heating/cooling rate. The calibration sequence included a baseline calibration for the determination of the time constants and capacitances of the sample and reference sensor using a sapphire standard, an enthalpy and temperature calibration for the correction of thermal resistance using indium as standard (ΔH = 28.71 J/g, Tm = 428.8 K), and a heat capacity calibration with sapphire standard. Polarizing Optical Microscopy. A Zeiss Axioskop 40 equipped with a video camera and a fast frame grabber was used to follow the superstructure formation in the graft and block-graft systems. A Linkham temperature control unit (THMS600), equipped with TMS94 temperature programmer, was employed for the temperature-dependent studies. Images were recorded following slow cooling (1 K/min) from the melt state. In a second experiment the kinetics of superstructure formation were investigated by performing T-jumps from high temperatures (T = 473 K) to different final crystallization temperatures where the growth of the crystalline complex was followed. Subsequently, the system was heated with 1 K/min and the apparent melting temperature of the superstructure was recorded. X-ray Scattering. Small-angle X-ray scattering (SAXS) measurements were made using Cu Kα radiation (RigakuMicroMax 007 X-ray generator, Osmic Confocal Max-Flux curved multilayer optics). Oriented fibers of 1.0 mm diameter were prepared by a mini-extruder. The patterns were anisotropic with strong peaks in the equatorial direction revealing ordering of chains along the extrusion direction. 2D scattering patterns were recorded on a Mar345 image plate at a sample-to-detector distance of 1.8 m. Wide-angle X-ray scattering (WAXS) measurements were made with a sample-to-detector distance of 35 cm using also a Mar345 detector. Intensity distributions as a function of the modulus of the scattering vector, q = (4π/λ) sin(2θ/2), where 2θ is the scattering angle, were obtained by radial averaging. In both SAXS and WAXS experiments, temperature-dependent measurements were made in the range from 303 to 433 K in 10 K steps on heating and subsequent cooling. Dielectric Spectroscopy (DS). Sample capacitors were prepared in a glovebox under a controlled nitrogen atmosphere. Samples were heated up to the melting temperature of the complex and hot pressed between two mirror polished brass electrodes to a thickness of 100 μm maintained with Teflon spacers. Dielectric spectroscopy measurements were made with a Novocontrol BDS system composed of a frequency response analyzer (Solartron Schlumberger FRA 1260) and a broadband dielectric converter. The experiments were performed at atmospheric pressure in the temperature range from 193 to 473 K and for frequencies in the range from 0.01 to 106 Hz. The conductivity has been studied using the analysis of the complex conductivity function through σ* = σ′ + iσ″, which is related to the complex dielectric permittivity with σ* = iωε0ε*.41,42 The dc conductivity has been analyzed with respect to the random free energy barrier model by J. Dyre according to which the charge carriers are hopping in a spatially randomly varying energy barriers.43 According to the model, the onset of dc conductivity is determined by crossing the highest energy barrier. The model provides an analytical expression for the complex dielectric function as ε*(ω) = ε∞ + σ0τe/[ε0 ln(1 + iωτe)], where ε∞ is the value of ε′ in the limit of high frequencies and σ0, τe are the dc conductivity and characteristic time of ion motion. It predicts a universal shape for the conductivity contribution based on these two parameters. The model predictions were tested earlier against the experimental data for the PEO/LiTf electrolyte.34 The predictions could only partially fit the experimental data and within a limited frequency range. We attribute these deviations to the nature of polymer electrolyte. The present system cannot be considered as disordered, and in addition, possible

formation of the complex is strongly temperature-dependent. Alternatively, the dc conductivity can be obtained by the plateau in the real part, σ′, without invoking any model. Rheology. A TA Instruments AR-G2 with a magnetic bearing that allows for nanotorque control was used for recording the viscoelastic properties of the polymer electrolytes. Measurements were made with the environmental test chamber (ETC) as a function of temperature. Samples were prepared on the lower rheometer plate (8 mm), the upper plate was brought into contact, and the gap thickness was adjusted. The linear and nonlinear viscoelastic regions were determined by the strain amplitude dependence of the complex shear modulus |G*| at ω = 10 rad/s. Measurements involved isothermal frequency scans within the range 10−1 < ω < 102 rad/s at selected temperatures and isochronal temperature ramps with ω = 10 rad/s between 303 and 423 K.

3. RESULTS AND DISCUSSION 3.1. Thermal Properties. The thermodynamic properties of the grafted (PHOS-g-PEO) and block-grafted (PS-b-(PHOSg-PEO)) systems are discussed with respect to the DSC traces of Figure 2. Evidently their qualitative characteristics are similar

Figure 2. DSC traces of (a) PHOS-g-PEO/LiTf, (b) PS-b-(PHOS-gPEO)/LiTf, and (c) PEO/LiTf electrolytes. All measurements correspond to the second heating run (rate 10 K/min). Vertical lines indicate the glass temperature of PEO.

to those of the respective PEO homopolymer electrolytes, studied earlier under the same conditions (Figure 2c).34 Neat PHOS-g-PEO (Figure 2a) exhibits the thermal characteristics of PEO with a glass temperature at 218 K followed by the melting of crystalline PEO at 331 K. The glass temperature of the PS backbone is not evident, presumably because of its small fraction. The electrolytes with low salt concentration, i.e., [EO]:[Li+] = 12:1, shows a glass temperature at 229 K followed by the cold crystallization of PEO and the simultaneous formation of the crystalline complex. Subsequently, PEO crystals melt at 322 K. At even higher temperatures there is a broad endotherm with a peak around 379 K that signifies the melting of the crystalline complex. The enthalpy of melting of the crystalline complex is 6 times lower than in the corresponding PEO homopolymer electrolyte at the same composition, revealing a suppressed complex formation. In the electrolyte with [EO]:[Li+] = 8:1 composition, there exists again a glass temperature at 236 K (i.e., some 18 K higher than in PEO/LiTf), followed by the melting of PEO crystals at 324 K and by a broad and weak melting peak with maximum at 402 K due to the melting of the crystalline complex. The enthalpy of melting is now 2 times lower than that in the corresponding PEO/LiTf. At [EO]:[Li+] = 6:1, a glass temperature is evident at 235 K. The melting of C

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Macromolecules PEO crystals now take place at 318 K while the complex melts at 407 K with nearly 4 times lower enthalpic content as compared to the PEO/LiTf electrolyte. For the grafted electrolytes with the higher salt concentrations, 4:1 and 3:1, the melting of the crystalline complex appears at 425 and 429 K, respectively, while the enthalpy of melting is also about 3 times lower than the corresponding PEO/LiTf electrolytes. In all cases, there is reduced crystallinity both for PEO crystals as well as for the complex. The degree of crystallinity of PEO is calculated using Xc = ΔH/(wPEOΔH0PEO), where wPEO is the weight fraction of PEO (ΔH0PEO = 196 J/g). The thermodynamic data for the PHOS-g-PEO/LiTf grafted electrolytes are summarized in Table 2.

Figure 3. Melting temperatures of complex (filled symbols) and of PEO crystals (empty symbols) in PEO/LiTf (red symbols), PHOS-gPEO/LiTf (green symbols), and PS-b-(PHOS-g-PEO)/LiTf electrolytes. The liquidus (melting temperatures of crystalline complex) and solidus lines (melting temperatures of PEO crystals) are also depicted.

Table 2. Melting Temperature of PEO Crystals and of the Crystalline Complex, Corresponding Change in Enthalpy and PEO Degree of Crystallinity for the PHOS-g-PEO/LiTf Grafted Electrolytes PHOS-gPEO/LiTf

Tg (K)

T′mPEO (K)

ΔHCrPEO (J/g)

XcPEO

neat 12:1 8:1 6:1 4:1 3:1

218 229 236 235 231

331 322 324 318

154.2 44.0 39.0 20.0

0.79 0.29 0.29 0.16

T′mcomplex (K)

ΔHcomplex (J/g)

379 402 407 425 429

4.5 24.0 21.0 35.0 47.0

regimes can be discussed. At temperatures below 326 K (i.e., solidus), the electrolytes consist of a crystalline PEO phase (except at the stoichiometric composition), a crystalline complex of the type (PEO)3LiTf,44 and an amorphous phase. At temperatures above 326 K but below the melting of the crystalline complex (i.e., liquidus) both electrolytes are composed of an amorphous phase and of a crystalline complex that starts to dissolve gradually to the amorphous phase and disappear at the liquidus line. At temperatures above the liquidus line, the systems are not totally amorphous as the anions still have some level of coordination (see below with respect to the viscoelastic properties). In comparing the PEO/ LiTf with the grafted and block-grafted electrolytes, we find that the PEO melting temperatures is invariably around 326 K; however, the melting temperatures of the complex are reduced. As a final note, the phase diagram of Figure 3 conceals the fact that the PEO and complex degrees of crystallinity are significantly reduced in the PHOS-g-PEO/LiTf and PS-b(PHOS-g-PEO)/LiTf electrolytes. In order to better understand the complex crystal formation in the PHOS-g-PEO/LiTf and PS-b-(PHOS-g-PEO)/LiTf electrolytes, we used POM to follow the self-assembly at the level of the spherulitic superstructure. Some representative POM images are shown in Figure 4. In general, at [EO]:[Li+] = 4:1 salt concentration the superstructure is nearly spherulitic whereas at the more dilute, [EO]:[Li+] = 12:1 concentration there are deviations from a spherulitic shape suggesting variable growth rates at different directions. We followed the kinetics of superstructure formation, and the growth rates were obtained under isothermal conditions for different crystallization temperatures. Figure 4 depicts the apparent melting temperatures and spherulitic growth rates as a function of the crystallization temperature for [EO]:[Li+] = 12:1 and 4:1 salt concentrations. The equilibrium melting temperatures (obtained by extrapolation) are at 430 and 452 K for the PHOS-gPEO/LiTf with the low and high electrolyte concentration, respectively. In the PS-b-(PHOS-g-PEO)/LiTf electrolytes these temperatures are at 435 and 459 K, respectively. In any case, the equilibrium melting temperature is higher for the higher salt concentrations, as anticipated from DSC. The growth rates at the high salt concentration ([EO]:[Li+] = 4:1) are significantly higher and show a strong dependence on the crystallization temperature as anticipated from the homopolymer electrolytes PEO/LiTf.34 However, the growth

The thermodynamic properties of the PS-b-(PHOS-g-PEO)/ LiTf block-grafted electrolytes (Figure 2b) bear similarities to the PHOS-g-PEO/LiTf grafted electrolytes and to the corresponding PEO/LiTf. In all cases there is a reduction of PEO and of complex crystallinity. For the [EO]:[Li+] = 12:1 composition, the melting enthalpy of the complex is almost 8 times lower than that of the corresponding PEO/LiTf electrolyte. For the remaining compositions the enthalpy of melting of the complex is about 3 times lower than the corresponding homopolymer electrolytes. The glass temperatures and melting of PEO and complex crystals as well as the change in enthalpy are summarized in Table 3. Evidently, this specific molecular architecture, of densely grafted PEO chains on a PS backbone, significantly reduces the ability for PEO crystals and complex crystals to form. The phase diagrams of the PHOS-g-PEO/LiTf and PS-b(PHOS-g-PEO)/LiTf electrolytes are shown in Figure 3 for compositions in the range [EO]:[Li+] = 50:1 to 3:1 in comparison to the PEO/LiTf system. Different temperature Table 3. Melting Temperature of PEO Crystals and of the Crystalline Complex, Corresponding Change in Enthalpy, and PEO Degree of Crystallinity for the PS-b-(PHOS-gPEO)/LiTf Block-Grafted Electrolytes PS-b(PHOS-gPEO)/ LiTf neat 12:1 8:1 6:1 4:1 3:1

T′ m PEO Tg (K)

(K)

ΔHCrPEO (J/g)

XcPEO

216 232/245 238 241 238

330 326 326 324 317

151.0 36.0 37.3 32.0 1.0

0.78 0.24 0.28 0.26 0.01

T′mcomplex (K)

ΔHcomplex (J/g)

363 405 408 435 288

3.8 18.6 26.3 36.6 39.7 D

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Figure 4. (left) Apparent melting temperatures (top) and spherulitic growth rates (bottom) due to complex formation as a function of the final crystallization temperature. Solid symbols correspond to the grafted electrolytes and empty symbols correspond to the blockgrafted electrolytes with [EO]:[Li+] concentrations of 12:1 (squares) and 4:1 (circles). Arrows indicate the (extrapolated) equilibrium melting temperatures. (right) POM images, obtained under isochronal conditions at 404 K for PHOS-g-PEO/LiTf 4:1 (top) and at 365 K for PHOS-g-PEO/LiTf 12:1 polymer electrolytes (bottom). The scale bars indicate 100 μm.

Figure 5. WAXS patterns of PEO/LiTf (black line), PHOS-g-PEO/ LiTf (red line), and PS-b-(PHOS-g-PEO)/LiTf (blue line) polymer electrolytes with [EO]:[Li+] = 4:1, at 278 K. Dashed lines give the position and (hkl) indices of the Bragg reflections from the crystalline PEO3LiCF3SO3 complex.44 Gray regions indicate the positions of the most intense reflections of crystalline PEO in its monoclinic unit cell.

Figure 6 shows SAXS patterns for the PS-b-(PHOS-g-PEO) electrolytes as a function of the modulus of scattering vector, q.

rates at the low salt concentration systems ([EO]:[Li+] = 12:1) have nearly no dependence on the final crystallization temperature. In this particular case, we also measured different growth rates for different superstructures formed at the same temperature (red symbols connected by a red vertical line). We anticipate that these facts originate from the highly reduced crystallinities for the complex especially at the [EO]:[Li+] = 12:1 salt concentration (Tables 2 and 3). As we will discuss below, with respect to the more local structure, the reduced propensity for complex crystal formation plays an important role on ion transport. 3.2. Structural Characteristics. At the length scale of the unit cell, the structure was examined with WAXS. The main reflections of PHOS-g-PEO/LiTf and PS-b-(PHOS-g-PEO) are compared with those of the PEO/LiTf homopolymer electrolyte, at [EO]:[Li+] = 4:1 salt concentration in Figure 5. For this salt concentration and at 278 K, the complex crystal is the only ordered structure, and the observed reflections of both PHOSg-PEO/LiTf and PS-b-(PHOS-g-PEO)/LiTf coincide with those of PEO/LiTf. The unit cell of the complex crystal is unaffected by the specific molecular architectures and adopts the known crystal structure of PEO3LiCF3SO3.44 In the latter, PEO adopts a helical configuration parallel to the crystal b-axis. Each Li ion is coordinated by five oxygensthree from the PEO oxygens and one oxygen from each of the two adjacent CF3SO3− groups. Overall, the crystalline complex in LiTf as well as in PHOS-g-PEO/LiTf and PS-b-(PHOS-g-PEO)/LiTf has a monoclinic unit cell with lattice parameter a = 1.6768 nm, b = 0.8613 nm, c = 1.007 nm, and β = 121.02°. In addition, at this salt concentration there is a near complete suppression of PEO crystallinity as indicated by the weak intensity of the (120) and (032) reflections associated with the monoclinic unit cell of PEO. To determine the environment in which Li+ transport takes place the nanodomain structure was investigated with SAXS.

Figure 6. SAXS patterns of neat PS-b-(PHOS-g-PEO) (top) and its electrolytes with LiTf, at two temperatures below and above the PEO melting temperature: 313 K (left) and 343 K (right). Arrows and vertical lines indicate the peak positions from crystalline PEO (blue) and crystalline complex (green), respectively.

At 313 K, neat PS-b-(PHOS-g-PEO) exhibits a broad peak centered at q ∼ 0.33 nm−1. At this temperature PEO is crystalline and the peak reflects PEO crystals having average correlation distances of d1 ∼ 19 nm. At temperatures higher than the melting point of PEO crystals this feature is absent, suggesting correlation hole scattering.45 For the [EO]:[Li+] = 4:1 concentration one peak is observed at 0.46 nm−1. At this high salt concentration, the sample primarily consists of complex crystals having an average correlation distance of d2 ∼ 15 nm. For the [EO]:[Li+] = 12:1, 8:1, and 6:1 compositions the structure is more complex. At 313 K, the SAXS patterns reveal broad peaks. At these concentrations and according to E

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Macromolecules results from DSC, POM, and WAXS, both complex crystals and PEO crystals exist. The superposition of the corresponding broad SAXS features explains the complex diffraction pattern observed for the intermediate salt concentrations. In particular, for the 12:1 concentration at 323 K, the pattern suggests two independent SAXS peaks reflecting correlation distances between PEO crystals (d1 ∼ 19 nm) and between complex crystals (d2 ∼ 15 nm) (Figure S1, Supporting Information). At 343 K, PEO crystals melt leaving only the peak due to the complex. At salt concentrations [EO]:[Li+] = 6:1 and 4:1, the peak is at 0.35 and 0.46 nm−1, respectively. This is again in agreement with the DSC and POM results. In Figure 7 the structure of the complex and PEO crystals on the local (unit cell) and intermediate (lamellar) length scale is

Figure 8. (Top) dc conductivity of the PHOS-g-PEO/LiTf electrolytes (empty symbols) and (middle) of PS-b-(PHOS-g-PEO) blockgrafted electrolytes (empty crossed symbols) in comparison to PEO/ LiTf electrolytes (dotted lines). (Bottom) normalized conductivity defined as (σgraft or block‑graft)/( f PEO//LiTf*σPEO/LiTf). Red squares, green circles, and blue triangles correspond to the [EO]:[Li+] = 12:1, 8:1, and 4:1 salt concentrations, respectively.

Figure 7. Highly schematic representation of the PS-b-(PHOS-gPEO)/LiTf structure with the crystalline complex (left) and crystalline PEO domains (right). Dashed lines indicate the region of crystalline PEO. Arrows indicate characteristic correlation distances between PEO crystals (d1) and complex crystals (d2).

of the crystalline complex. Here a high fraction of Li ions are bound in the crystalline complex, therefore having a very low mobility. Moreover, the temperature dependence of all conductivities is neither Arrhenius nor Vogel−Fulcher− Tammann. The peculiar temperature dependence reflects the underlying structural changes (melting of PEO crystals and continuous melting of complex). These features of ion conductivity on lithium content were anticipated based on the behavior of the homopolymer electrolytes investigated earlier.34 These facts imply that the mechanism of ion transport is unaffected by the PHOS backbone. Ion transport takes place by segmental motion assisted by consecutive sub-Khun length jumps.34 Interestingly, for temperatures below the melting temperature of PEO crystals, the measured dc conductivities are significantly higher than those of the PEO/LiTf. This is a consequence of the reduced PEO degree of crystallinity in the present systems (Tables 2 and 3). The dense grafting architecture significantly enhances Li ion transport by reducing the degree of PEO and complex crystallinity. This is also reflected in the normalized conductivity values of Figure 8 (bottom). The normalized conductivity of PHOS-g-PEO/LiTf with [EO]:[Li+] = 12:1 is unity, while for the [EO]:[Li+] = 8:1 composition is 1.2 times higher than the PEO homopolymer case. To our knowledge, higher conductivities from the archetypal electrolyte PEO/LiTf have not been reported without the addition of oligomers or ionic liquids. The PS-b-(PHOS-g-PEO)/LiTf system with the blockgrafted architecture show similar qualitative characteristics with respect to the ionic conductivity (Figure 8, middle). The difference here from the previous case is the presence of a short PS block. The behavior is more informative when examining the normalized conductivities. For [EO]:[Li+] = 12:1, it shows 1.2 times higher conductivity than the corresponding

schematically summarized. Results obtained from SAXS/WAXS indicate the existence of an amorphous PEO matrix with chains stretching away from the PHOS backbone. Embedded in this matrix are regions of crystalline PEO and of crystalline complex as well as regions of amorphous PS mixed with the PHOS-gPEO chains. Because of the small volume fraction, the latter does not form a separate phase. In this picture ion transport can be facilitated by Li ions located near the PEO chain ends and away from PEO crystals. The role of ion mobility within the crystalline complex was discussed earlier.34 3.3. Ionic Conductivity. Knowledge of the structure at the different length scales allows investigating the effect of molecular architecture on the ionic conductivity. Measurements of dc conductivity were made by following two cooling/heating circles. Results from the heating/cooling circles for PHOS-gPEO/LiTf) and PS-b-(PHOS-g-PEO)/LiTf at [EO]:[Li+] = 8:1 are shown in Figure S2. They depict the usual hysteresis associated with PEO melting/crystallization. The results for PHOS-g-PEO/LiTf and PS-b-(PHOS-g-PEO)/LiTf are compared in Figure 8 for different temperatures and all salt concentrations. The normalized conductivity is also depicted defined as σ/φσ0.46,19 Here, σ is the conductivity of the electrolyte under study, φ is the volume fraction of the salt enriched phase, and σ0 is the conductivity of the homopolymer electrolyte (i.e., the PEO/LiTf). Starting with PHOS-g-PEO/ LiTf, conductivities with values similar to or even exceeding those of the homopolymer electrolytes at some salt concentrations are obtained. More specifically, conductivities are higher for the [EO]:[Li+] = 12:1, followed by the [EO]: [Li+] = 8:1 salt concentration. On the other hand, at [EO]: [Li+] = 4:1 salt concentration, conductivity is reduced because F

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8:1 and [EO]:[Li+] = 12:1) there is a two-step increase of the moduli at T ∼ 370 K and T ∼ 310 K signifying respectively the formation of the complex and of PEO crystals. The “plateau” modulus on complex formation is distinctly different in the three compositions: from a storage modulus of ∼108 Pa in the 4:1 composition to 4 × 106 Pa in 8:1 and to 3 × 104 Pa at the 12:1 composition. The low moduli of the complex crystal in the more dilute compositions are in agreement with the lower degree of crystallinity of the complex (Table 2) and the nearly temperature independent growth rates of the truncated superstructures found in POM (Figure 4). The temperature dependence of the moduli in the PS-b(PHOS-g-PEO) block-grafted electrolytes exhibits similarities to the graft electrolytes with two distinguishing characteristics (Figure 9). First, the storage and loss moduli at temperatures above the melting of complex are significantly higher than those of the PHOS-g-PEO system at the same salt concentration. Interestingly, even though the PS block, visualized as isolated PS molecules embedded in a PEO matrix (Figure 7), does not form a continuous phase, it affects the moduli. This can nicely be seen by comparing the moduli over heating/cooling cycles for the PHOS-g-PEO and PS-b-(PHOS-g-PEO) at the same salt concentration ([EO]:[Li+] = 8:1) (Figure S3). Second, G′ > G″ over a broad temperature range, except at temperatures in the vicinity of PEO crystallization. Among the different compositions, the PS-b-(PHOS-gPEO)/LiTf with the stoichiometric composition [EO]:[Li+] = 4:1 bears the highest moduli. This is not surprising given the higher crystallinity of the complex (Table 3) and high growth rates of the superstructures seen in POM (Figure 4). The temperature-dependent storage and loss moduli at this composition are shown on a cooling/heating cycle in Figure 10. The observed hysteresis is consistent with the DSC and POM results. Furthermore, the frequency-dependent moduli at 403 K reveal a purely elastic response of the complex crystals (G′(ω) ∼ ω0, G″(ω) ∼ ω0, and G′ > G″), in contrast to 453 K, where the dependence is viscoelastic. The latter point isat first sitesurprising, given the fact that at 453 K the crystalline complex is completely molten. We will return to this point below with respect to Figure S4 where the viscoelastic properties at high temperatures are compared. We now focus in the [EO]:[Li+] = 8:1 system because it combines favorable mechanical properties with high ionic conductivity. More specifically, for both PHOS-g-PEO/LiTf and PS-b-(PHOS-g-PEO) electrolytes, frequency sweeps at T = 338 K, i.e., at a temperature where the crystalline complex is the sole structure (Figure S4), display the typical response of a viscoelastic solid (G′(ω) ∼ G″(ω) ∼ ω1/4) due to the coordination and complexation of PEO with lithium ions. Interestingly, frequency sweeps at temperatures just above the melting of the complex, both for the grafted and block−grafted electrolytes, revealed a viscoelastic response (Figure S5). This situation here is similar to the PS-b-PEO/LiTf system29 and is discussed in terms of transient structures of the electrolyte with the PEO chains that exist at temperatures above the melting of the complex. Clearly, further work is needed to identify the characteristics of this structure. Nevertheless, both systems show enhanced Li ion conductivities and good mechanical properties at the [EO]:[Li+] = 8:1 concentration.

homopolymer, whereas the normalized conductivity of PS-b(PHOS-g-PEO)/LiTf [EO]:[Li+] = 8:1 is unity. The observed high ionic conductivities arise from the continuous conductive phase, the reduced propensity for complexation, and the suppression of PEO crystallization. All these properties originate from the specific macromolecular architecture. If the high ionic conductivity can be combined with elastic response and increased moduli, then this material can be favorable for applications. Therefore, their viscoelastic properties are examined next. 3.4. Viscoelastic Behavior. The linear viscoelastic properties of the grafted and block-grafted electrolytes are discussed with respect to Figures 9 and 10. Figure 9 compares the storage

Figure 9. Storage (filled symbols) and loss (open symbols) moduli during cooling with a rate of 2 K/min at a frequency of 10 rad/s for PHOS-g-PEO/LiTf (left) and PS-b-(PHOS-g-PEO) electrolytes (right) at three concentrations: [EO]:[Li+] = 12:1 (green), 8:1 (red), and 4:1 (blue) salt concentrations.

Figure 10. Storage (filled symbols) and loss (open symbols) moduli of the PS-b-(PHOS-g-PEO)/LiTf electrolyte during cooling (blue) and subsequent heating (red) with 5 K/min at a frequency of 10 rad/s. Inset: storage (filled symbols) and loss (open symbols) moduli versus frequency at 453 K (T < Tmcomplex) and at 403 K (T > Tmcomplex) for the PS-b-(PHOS-g-PEO)/LiTf electrolyte with [EO]:[Li+] = 4:1.

and loss moduli of the PHOS-g-PEO/LiTf and of the PS-b(PHOS-g-PEO) electrolytes at three salt concentrations. Measurements were made under isochronal conditions on cooling. For PHOS-g-PEO/LiTf electrolytes, at a [EO]:[Li+] = 4:1 salt concentration, a steep increase of the storage moduli is observed, from G′ ∼ 103 Pa to G′ ∼ 108 Pa, signifying the formation of the crystalline complex at the stoichiometric composition. For the more diluted electrolytes ([EO]:[Li+] =

4. CONCLUSION Macromolecular architecture is shown to be of key importance in the design of efficient solid polymer electrolytes. Densely G

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grafted PHOS-g-PEO as well as block copolymers with PS, PSb-(PHOS-g-PEO), were doped with LiCF3SO3, and their thermodynamic properties, structural characteristics, ionic conductivity, and viscoelastic properties were examined for a range of salt concentrations. This specific architecture with short and densely grafted PEO chains has a number of advantages. Thermodynamically speaking, the behavior is qualitatively similar to the PEO homopolymer electrolytes. However, a major difference being the significantly lower degree of crystallinity both for PEO crystals and complex crystals. This was shown to be a key point in the design and performance of the system. The suppressed PEO crystallization resulted in an increase by 1−2 orders of magnitude of the ionic conductivity below the melting temperature of PEO crystals. In addition, the conductivity in the vital zone above 333 K was similar to or even higher than that of the corresponding PEO homopolymer electrolytes. To our knowledge, higher conductivities from the archetypal electrolyte PEO/LiTf have not been reported without the addition of oligomers or ionic liquids. At the same time, both PHOS-g-PEO/LiTf and PS-b(PHOS-g-PEO)/LiTf had improved mechanical properties at all compositions investigated as compared to the archetypal PEO/LiTf. Along these lines the PS-b-(PHOS-g-PEO)/LiTf with of [EO]:[Li+] = 8:1 salt concentration combines high elasticity (G′ ∼ 2 × 106 Pa) with high ion conductivity (σ′ = 6 × 10−5 S/cm) at 343 K in an ideal way. Evidently, tailored macromolecular architectures in SPEs hold the key toward a simultaneous increase in ionic conductivity and elasticity. Further architectures that can completely suppress PEO crystallization and complex crystal formation need to be designed and examined. Work in this direction is currently in progress in our laboratory.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.6b00290. Small-angle X-ray scattering and viscoelastic properties of the PHOS-g-PEO as well as block copolymers with polystyrene PS-b-(PHOS-g-PEO) for certain [EO]:[Li+] ratios (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail gfl[email protected] (G.F.). *E-mail [email protected] (G.Z.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was cofinanced by the E.U.- European Social Fund and the Greek Ministry of Development - GSRT in the framework of the THALIS program, and the “Excellence in the Research Institutes” program. The current work was also supported by the Research unit on Dynamics and Thermodynamics of the UoI cofinanced by the European Union and the Greek state under NSRF 2007-2013 (Region of Epirus, call 18). H

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