role of preparation methods on the structural and dielectric properties of plasticized polymer blend...

Accepted Manuscript

Title: Role of preparation methods on the structural anddielectric properties of plasticized polymer blend electrolytes:Correlation between ionic conductivity and dielectricparameters

Author: R.J. Sengwa Priyanka Dhatarwal Shobhna Choudhary

PII: S0013-4686(14)01539-4DOI: http://dx.doi.org/doi:10.1016/j.electacta.2014.07.120Reference: EA 23158

To appear in: Electrochimica Acta

Received date: 18-5-2014Revised date: 22-7-2014Accepted date: 23-7-2014

Please cite this article as: R.J. Sengwa, P. Dhatarwal, S. Choudhary, Role ofpreparation methods on the structural and dielectric properties of plasticized polymerblend electrolytes: Correlation between ionic conductivity and dielectric parameters,Electrochimica Acta (2014), http://dx.doi.org/10.1016/j.electacta.2014.07.120

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http://dx.doi.org/doi:10.1016/j.electacta.2014.07.120

http://dx.doi.org/10.1016/j.electacta.2014.07.120

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Graphical Abstract

Li+ −33SOCF

O

CO

C

Ion-dipolarinteraction

O

O

O

OH

O O OO

HO

PMMA

PEO

PEG

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Research Highlights

• PEO–PMMA–LiCF3SO3–x wt% PEG electrolytes were prepared by different

methods.

• Dielectric/electrical properties vary non-monotonously with PEG concentration.

• The ionic conductivity and amorphous phase change with preparation methods.

• Ion conduction is through hopping mechanism coupled with polymer segmental

motion.

• Ionic conductivity is ~10–5 S cm–1 at room temperature.

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Role of preparation methods on the structural and dielectric

properties of plasticized polymer blend electrolytes:

Correlation between ionic conductivity and dielectric

parameters

R. J. Sengwa∗, Priyanka Dhatarwal, Shobhna Choudhary Dielectric Research Laboratory, Department of Physics, Jai Narain Vyas University, Jodhpur – 342 005, India

ABSTRACT

The polymer blend based electrolyte films consisted of poly(ethylene oxide) (PEO) and

poly(methyl methacrylate) (PMMA) with lithium triflate (LiCF3SO3) as a dopant ionic

salt and poly(ethylene glycol) (PEG) as plasticizer have been prepared by solution cast

melt–pressed and ultrasonic assisted followed by microwave irradiated solution cast

melt–pressed methods. The X–ray diffraction study infers that the amorphous phase of

(PEO–PMMA)–LiCF3SO3–x wt% PEG electrolytes decreases with the increase of PEG

concentration. The complex dielectric function, ac electrical conductivity, electric

modulus and the impedance spectra of these electrolytes have been investigated over the

frequency range from 20 Hz to 1 MHz. It is found that all the dielectric/electrical

functions of these electrolytes vary anomalously with the increase of PEG concentration,

and also with the change of samples preparation methods. The activation energies have

been determined from the temperature dependent values of dc ionic conductivity,

polymer segmental relaxation time and dielectric strength. Results reveal that besides the

amorphicity, the ionic conductivity of these electrolytes is also governed by the

relaxation time and the dielectric strength, and the transport of ions is due to hopping

mechanism which is coupled with segmental motion of polymers chain. Room

temperature ionic conductivity values of the PEO–PMMA blend based electrolytes are

found about one to two orders of magnitude higher than that of the PEO and PMMA

based electrolytes.

∗ Corresponding author (E-mail: [email protected]) Tel.: +91 291 2720857; fax: +91 291 2649465

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Keywords: Polymer blend, Solid polymer electrolyte, Dielectric relaxation, Ionic

conductivity

1. Introduction

Solid polymer electrolytes (SPEs) have been recognized as the most suitable

flexible type, leak proof and light weight novel materials for the fabrication of ion

conducting devices particularly the rechargeable batteries for portable electronic

equipments [1–18]. The SPEs are complexes of polymer with cations of the dopant salt

and mostly have high amorphicity. But the low ionic conductivity at ambient temperature

limits their several technological applications. In these electrolytes the dynamics of

polymer chains is critical for the ions transportation. Therefore, the knowledge on how

the ions are coupled with the chain dynamics of the multiphase complex structures of the

SPEs and the modification in these structures with the preparation methods for the

improvement of ionic conductivity are nowadays prime issues [5,6,12,15,16,19–21].

So far, enormous work is in progress on the SPE materials with the choices of

polymers and the alkali metal salts in order to achieve their high ionic conductivity (≥ 10–

5 S cm–1) at ambient temperature for the realistic device applications. Mostly,

poly(ethylene oxide) (PEO) matrix based complexes with various lithium salts have been

an intense field of research [5–17,19–30]. The PEO has low lattice energy, low glass

transition temperature, high solvating power for alkali metal salts and an ease to form

flexible type film. These facts make it an impending polymer in preparation of the SPEs.

But the high crystallinity of PEO is its main disadvantage because effective ion transport

can be achieved utmost in an amorphous medium of the host polymer. Besides the PEO,

poly(methyl methacrylate) (PMMA) matrix, which has amorphicity more than 96%, is

also preferred in preparation of SPE materials [3,31–38]. The highly active electron

donating carbonyl (C=Ö) functional group of PMMA can easily coordinate with the

cations of alkali metal salts, which results in PMMA–salt complexes. Although PMMA

has light weight with high transparency, high strength and dimensional stability, but its

brittle property under a loaded force limits the industrial and technological suitability of

the PMMA based SPEs.

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In the advancement of SPE materials, continuous work is in progress on the PEO–

PMMA blend matrix based electrolytes to overcome the drawbacks of the pristine

polymers [39–48]. The advantages of this blend are; the addition of PEO in PMMA

matrix results in the increase of PMMA flexibility and reduces its brittleness, whereas the

PMMA environment increases the amorphous phase of PEO. These properties of the

PEO–PMMA blend matrix have been recognized its suitability as novel polymeric blend

matrix of tailor–made properties for the preparation of SPEs. Our survey of literature

reveals that PEO–PMMA matrix is used with lithium triflate [39], lithium perchlorate

[40–45] and silver nitrate [46–48] as ion conducting salt with different plasticizers and

inorganic nanofillers for the enhancement of the ionic conductivity of their solution cast

electrolyte films. But the PEO–PMMA matrix with lithium trifluoromethanesulfonate

(lithium triflate, LiCF3SO3) as ion conducting salt and poly(ethylene glycol) (PEG) as

plasticizer has not been studied for its possible applications in the advancement of ion

conducting electrochromic devices. As compared to the high dielectric constant

plasticizers ethylene carbonate (EC) εs = 85.1 and propylene carbonate (PC) εs = 69.0, at

room temperature, the PEG has εs = 22.1, which is very low. So it is interesting to know

how a low dielectric constant plasticizer affects the conductivity of the electrolytes.

Further, PEG is a linear chain molecule and its repeat unit is same as PEO backbone unit

and therefore, how the PEO structure varies with increase of PEG concentration, is also

interesting to study in these electrolytes.

In this paper, we have first time prepared the PEO–PMMA blend matrix based

electrolytes comprising LiCF3SO3 salt and varying concentrations of PEG by classical

solution cast melt–pressed and ultrasonic assisted followed by microwave irradiated

solution cast melt–pressed methods. The aim of ultrasound treatment is, firstly, to

enhance the PEO–PMMA blend miscibility, and secondly to dissociate the salt clusters in

the electrolyte solution, if present any. The microwave irradiation is given to the solution

with the aim to modify the dipolar orientation and the PEG dynamics which occur at

microwave frequencies in the solutions. This study has been carried out to explore the

effects of sample preparation methods and plasticizer concentration on the dielectric

properties, ionic conductivity and structural dynamics of the SPE films using dielectric

relaxation spectroscopy. Further, an attempt has been made to correlate the temperature

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dependent ionic conductivity to the polymer chain dynamics (segmental relaxation) by

fitting the data to the Vogel–Tamman–Fulcher (VTF) equation of these polymer

electrolytes.

2. Experimental

2.1. Sample preparation

The PEO (Mw = 6×105 g mol–1), PMMA (Mw = 3.5×105 g mol–1), PEG (Mw = 200

g mol–1) and LiCF3SO3 were obtained from Sigma–Aldrich, USA. The anhydrous

acetonitrile and tetrahydrofuran of spectroscopic grade were purchased from Loba

Chemie, India. The polymers and the salt were vacuum dried at 50 °C for at least 12 h

before using. For the preparation of varying PEG concentrations (PEO–PMMA)–

LiCF3SO3–x wt% PEG electrolyte films, the 50:50 w/w PEO:PMMA blend was used.

The average molar ratio 9:1 was set for the total number of the ethylene oxide units (EO)

and the carbonyl groups (C=O) of the polymers blend to the lithium cations (Li+) of the

salt. The PEG concentrations x = 0, 5, 10 and 15 with respect to the weight of PEO–

PMMA blend were used.

Two different processing methods were used for preparation of the electrolyte

films. (i) ‘Classical’ solution casting (SC) method: In this process, initially, the required

amounts of PEO and PMMA were dissolved in acetonitrile and tetrahydrofuran,

respectively, in separate glass bottles. After that LiCF3SO3 was added in PEO solution,

and it was dissolved and mixed homogenously by magnetic stirrer. This PEO electrolyte

solution was mixed with PMMA solution which resulted in the polymer blend electrolyte

solution. Finally, the required amounts of PEG for different concentrations were added

into the PEO–PMMA blend electrolyte solutions and mixed homogeneously by magnetic

stirring. Each PEG concentration homogenous solution was divided into two equal parts.

The first parts were cast onto Teflon petri dishes and by slow evaporation of solvent at

room temperature, the solution cast (SC) prepared (PEO–PMMA)–LiCF3SO3–x wt%

PEG films were achieved. (ii) Ultrasonic (US) assisted followed by microwave (MW)

irradiated solution casting method: In this method, the second parts of each polymeric

electrolyte solutions of varying PEG concentrations, which were prepared as mentioned

above, were sonicated by ultrasonicator (250 W power, 25 kHz frequency) for 10 min

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duration with ON–OFF step of 15 s. In this processing the stainless steel sonotrode was

directly immersed into the electrolyte solution for strong dose of the ultrasound. After

that each of the solution was irradiated by microwave electromagnetic energy in a

domestic microwave oven (600 W power, 2.45 GHz frequency) for 2 min duration and 10

s irradiation step with intermediate cooling. These solutions were cast onto petri dishes

which resulted in the US–MW irradiated solution cast films after the solvents

evaporation.

The surfaces of the solution cast films prepared by above mentioned methods

were found uneven. Therefore, the smooth surfaces of these films were achieved by

melt–pressed technique. In this technique, each electrolyte film was initially vacuum

dried at 40 °C for 24 h. After that each film was melted by heating it up to 130 °C in

circular stainless steel die having suitable spacer using polymer film making unit. This

melted material was pressed under 2 tons of pressure per unit area and cooled slowly up

to room temperature which made the film of smooth surfaces. The same steps were

repeated for each film.

2.2. Characterizations

The X–ray diffraction (XRD) patterns of the SPE films and their constituents

were recorded in reflection mode using a PANalytical X′pert Pro MPD diffractometer of

Cu Kα radiation (1.5406 Å) operated at 45 kV and 40 mA with a scanned step size of

0.05°/s. The powder sample of LiCF3SO3 was tightly filled in the sample holder, whereas

the PEO, PMMA, PEO–PMMA blend and the SPE films were placed on the top of

sample holder during their XRD measurements in the 2θ range from 10–30° at room

temperature.

The dielectric relaxation spectroscopy (DRS) of the electrolyte films was carried

out using Agilent technologies 4284A precision LCR meter along with Agilent 16451B

solid dielectric test fixture in 1 V electric field of linear frequency f range from 20 Hz to 1

MHz at 30 °C, and also with the temperature variation for the 10 wt% PEG concentration

films. Frequency dependent values of capacitance Cp, resistance Rp and loss tangent (tanδ

= ε″/ε′ ) of the SPE films mounted in the dielectric cell were measured in the parallel

circuit operation for the determination of their dielectric/electrical spectra. Prior the

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sample measurement, the open circuit calibration of the cell was performed to eliminate

the effect of stray capacitance of the cell leads. The dielectric test fixture was placed in a

microprocessor-controlled heating chamber to record the measurements at different

temperatures. Initially, all the measurements of varying PEG concentrations electrolyte

films were made at fixed temperature, 30 °C. After that the same films of 10 wt% PEG

concentration were used for their temperature dependence study from 30–55 °C. It was

found that the re-measured values of Cp, Rp and tanδ have good reproducibility after one

week duration, which confirms that there is no immediate aging effect in the studied

electrolyte films. The spectra of intensive quantities, namely complex dielectric function

ε*(ω) = ε′ – jε″, real part of alternating current (ac) electrical conductivity σ′ = ωε0ε″ and

electric modulus M*(ω) = M′ + jM″, and the extensive quantity i.e. complex impedance

Z*(ω) = Z′ – jZ″ of the electrolyte films were determined using the expressions described

in detail elsewhere [3,9].

3. Results and Discussion

3.1. Structural analysis

The XRD patterns of PEO, PMMA, PEO–PMMA blend, LiCF3SO3, and the SC

and US–MW prepared (PEO–PMMA)–LiCF3SO3–x wt% PEG electrolytes over the

angular range 10–30° are shown in Fig. 1. The XRD pattern of a material provides the

structural properties related to crystalline phases (peak positions), phase concentration

(peak heights), amorphous content (back-ground hump) and crystallite size (peak widths).

The PEO peaks position 2θ and the corresponding intensity values of these materials are

determined by X’pert pro® software and these values are given in Table 1. The XRD

pattern of pure PEO (Fig. 1(a)) has sharp crystalline peaks at 2θ = 19.22° and 23.41°,

which are corresponding to crystal reflection planes 120 and concerted 112,032,

respectively [8,15,49]. Inset of Fig. 1(a) shows that PMMA has a broad and diffused peak

around 16°, which confirms its predominantly amorphous phase [3]. The PEO–PMMA

blend also has peaks at 19.19° and 23.33° which are corresponding to PEO crystalline

reflection planes, but the intensities of these peaks are low as compared to pristine PEO

(Table 1). The relative low intensity peaks of PEO–PMMA blend reveal that there is

some miscibility of PEO and PMMA which increases the amorphous content, but some

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amount of PEO crystalline phase also exists in the blend. The LiCF3SO3 has intense

crystalline peaks at 16.62, 19.87, 20.53, 22.77 and 24.71°, which confirm its high

crystallinity. The angular positions and relative intensities of these peaks are found in

consistent with the earlier reported XRD spectra of the LiCF3SO3 [50].

The differences in XRD patterns of (PEO–PMMA)–LiCF3SO3–x wt% PEG films

prepared by SC and US–MW processing methods (Fig. 1(b)) confirm some structural

variations due to the sample preparation methods. The crystalline peaks of LiCF3SO3 are

not found in the XRD patterns of these electrolytes, which infer that the salt is completely

dissociated due to formation of the cations complexes with the functional groups of

PEO–PMMA blend. Further, the absence of PEO peaks in the US–MW prepared film and

a very low intensity peak in the SC prepared film of the unplasticized (PEO–PMMA)–

LiCF3SO3 electrolyte infer that the formation of ion–dipolar complexes has changed the

polymer blend into an amorphous material. Due to four to six coordination sites of

lithium cations with the oxygen atoms of PEO [10,22], and the carbonyl groups of

PMMA [32] and also formation of some miscible phase in the PEO–PMMA blend result

in complete amorphous phase of the (PEO–PMMA)–LiCF3SO3 electrolyte having the salt

concentration molar ratio [EO+(C=O)]:Li+=9:1. With the addition of PEG as plasticizer

in these electrolytes both the peaks of PEO have appeared, which confirm the

recrystallization of a small amount of PEO. The enhancement of PEO peaks intensities

with the increase of PEG concentration of the electrolytes (Table 1) is attributed to a

gradual increase of crystalline phase, because the peaks intensities are directly related

with the material crystallinity [8,49]. In comparison to the SC prepared films the peak

intensities of US–MW prepared films are low (Table 1), which confirm that the high

intensity ultrasonication disturbs the recrystallization of PEO. The end hydroxyl groups

of PEG form the strong intra- and inter-molecular hydrogen bonding [51]. It seems that

the hydroxyl groups of PEG form the ion-dipolar coordination with the lithium cations

which release some of the ether oxygen atoms of PEO from the polymer-ion complexes.

Due to this fact some amount of uncomplexed PEO recrystallizes. As the PEG

concentration increases in the electrolyte, more amount of PEO releases from the

complexes, which results in gradual increase of PEO crystalline phase. Further, the

presence of same unit in the backbone of PEG and PEO molecules also favours the

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enhancement of PEO crystalline phase as the PEG concentration increases in the PEO–

PMMA blend based electrolytes. This structural finding of the PEG plasticized

electrolyte is very interesting, and so far this type of behaviour has not been observed in

other type of plasticizer added PEO based solid polymer electrolytes [27,30,39,41,46].

Further, authors repeated the XRD scans of some of the samples within one week

duration, and found good reproducibility of the scans which confirmed that the samples

did not suffer from immediate aging effect.

3.2. Dielectric and electrical behaviour

The complex permittivity (real part ε′ and dielectric loss ε″) spectra of (PEO–

PMMA)–LiCF3SO3–x wt% PEG electrolyte films prepared through SC and US–MW

processes are shown in Figs. 2(a) and 2(b), respectively. The ε′ values of these

electrolytes are in the order of several thousands at low frequencies which decrease non–

linearly with increasing frequency and approach the steady state (high frequency limiting

permittivity ε∞) near 1 MHz, as shown by the enlarged view in the inset of Fig. 2. The

large ε′ values of these polymeric electrolytes at low frequencies are due to the electrode

polarization (EP) effect occurring as a result of the accumulation of ions near the

electrodes surfaces of dielectric test fixture [9,10,15]. These ε′ spectra have point of

inflection around 1 kHz, therefore the ε′ values at this frequency are considered as the

static permittivity εs of these electrolytes. The εs and ε∞ values of the polymeric

electrolytes with PEG concentration are given in Table 2. The values of dielectric

strength Δε = εs – ε∞ of these materials are also recorded in Table 2.

As compared to the polymer matrix, the plasticizers have high value of static

permittivity. Therefore, by the algebraic additive rule, it is expected that the increase of

plasticizer concentration in the polymeric electrolyte must increase the dielectric strength

beside the modification in material physical properties such as flexibility, viscosity,

microstructure, etc. Further, it is believed that the increase of dielectric strength of the

polymer electrolyte with the addition of plasticizer must support the ion conduction

process. But the dielectric strength of a composite material is governed by the dipolar

ordering in its complex structure. The parallel dipolar ordering increases the dielectric

strength whereas it is reduced in case of anti-parallel dipolar ordering. The comparison of

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Figs. 2(a) and 2(b) and also the Δε values given in Table 2 reveals that the ε′ values of

these electrolyte films change anomalously with the increase of PEG concentration and

these values are also influenced by the sample preparation methods. These results

confirm that there is significant alteration in the strength of ion–dipolar interactions and

the dipolar ordering in the complexes with increase of PEG concentration. The non-

monotonous behaviour of Δε with PEG concentration also reveals that the interaction of

end-hydroxyl groups of PEG with lithium cations and functional groups of polymers

makes them randomly aligned in the complexes, which is also favoured by the changes in

their crystalline phase (Table 1).

Figs. 2(c) and 2(d) show the temperature dependent ε′ and ε″ spectra of the SC

and US–MW prepared (PEO–PMMA)–LiCF3SO3–10 wt% PEG films. It is found that

both the ε′ and ε″ values increase with increasing temperature, which is attributed to the

increase in charge density as an additional contribution from the interfacial polarization.

The temperature dependent εs, ε∞ and Δε values of the 10% PEG concentration electrolyte

film prepared by SC and US–MW methods are recorded in Table 3. It has also been

revealed that the magnitude of ε″ values of these electrolyte films follows their ε′ values.

Further, each of the ε″ spectrum has single relaxation peak corresponding to the

segmental motion (α-relaxation) in the miscible polymer blend. Due to large difference in

the backbone units of PEO and PMMA chains there must be separate relaxation peaks in

the ε″ spectra on the frequency scale, but the appearance of single relaxation peak also

suggests a cooperative chain segmental dielectric relaxation process in these PEO–

PMMA blend based electrolytes. The Cole–Cole plots (ε″ vs ε′) of the electrolyte films

are shown in insets of the ε″ spectra of Fig. 2. The high frequency data in these plots has

appeared as semicircular arcs corresponding to the bulk properties whereas the low

frequency data forms the spikes due to dominance contribution of the EP effect in the low

frequency dielectric properties. The Δε values of these electrolytes as determined from

their estimated εs values are in consistence with the other polymeric electrolytes which

are determined by fitting the complex permittivity data to the Havriliak–Negami function

[18].

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The ac ionic conductivity σ′ and the loss tangent tanδ spectra of SC and US–MW

processed prepared (PEO–PMMA)–LiCF3SO3–x wt% PEG electrolytes are depicted in

Figs. 3(a) and 3(b), respectively. The σ′ values of these polymeric electrolytes increase

non–linearly with increase of frequency on logarithm scale. It has been established that

the ion transportation in polymeric electrolytes occurs on different time scale under the

influence of alternating current (ac) electric field [8,17,18,52,53]. The jumps of ions

between different ion sites through hopping mechanism are demonstrated by random

barrier model which consequences the ac conductivity dispersion [52]. The short–time

ion dynamics is characterized by back–and–forth motion over the limited range in

disordered polymer matrix, ‘subdiffusive’ dynamics, which leads to dispersive ac

conductivity at high frequencies. Whereas the long–time ion dynamics is characterized by

random walks resulting in long–range ion transport, ‘diffusive’ dynamics, which leads to

the plateau corresponding to the direct current (dc) ionic conductivity at frequencies

lower than that of the dispersive ac conductivity region. The dc conductivity of such

electrolytes mostly obeys the Arrhenius temperature dependent behaviour [7,17]. The ac

conductivity of the polymeric electrolytes also obeys time–temperature superposition, i.e.

it is possible to scale data at different temperatures to one single master curve which is

roughly same for all disorder materials [18,52,54]. Further, at high frequencies, the log–

log plot of ac conductivity follows apparent power law behaviour [55]. Therefore the

detailed studies of frequency dependent ac conductivity behaviour of SPEs have

academic interest, besides the confirmation of their suitability for technological

applications.

The σ′ spectra of the investigated electrolytes show three regions on frequency

scale; (i) the electrode polarization dominated low frequency region below 10 kHz, (ii)

the dc plateau region above 30 kHz up to nearly 300 kHz, and (iii) the power law

dispersive region above 300 kHz. The values of dc ionic conductivity σdc of these

electrolytes are determined by fitting the σ′ spectra above 30 kHz to the conventional

Jonscher’s power law σ′(ω) = σdc + Aωn [55], where A is the pre–exponential factor and n

is the fractional exponent ranging between 0 and 1. These fits are shown by solid lines in

the σ′ spectra. Further, it is observed that the σ′ increases with the increasing temperature

of the electrolytes (Figs. 3(c) and 3(d)). This increase has two implications, firstly, the

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mobility of the ions increases due to increased polymer chain segmental motion, and

secondly, the ion concentration increases. But the XRD spectra reveals that the total salt

is in dissociated form and hence the increase of temperature does not have any additional

effect in dissociating the salt. This fact confirms that the increase of σ′ values of these

electrolytes is due to increase of ions mobility only. Further, the presence of small salt

clusters of nm size is also ruled out because the salt concentration is not very high. The

evaluated values of σdc and n of the electrolytes with the PEG concentration and the

temperature variation are given in Tables 2 and 3, respectively. The n values of the

electrolytes are found in the range from 0.88 to 0.98, which suggests that the ions

transportation in these electrolyte also takes place through hopping mechanism as

reported for other polymeric electrolytes [10,15,17,18,52,56]. Further, it is found that the

n values of the electrolyte increase with increase of temperature (Table 3). This agrees

well with the idea that n physically represents the strength of the ion–ion and ion–dipolar

interactions and as the temperature increases these interactions are expected to decrease

[25].

The tanδ spectra of the electrolytes shown in Fig. 3 have Debye type relaxation

peaks appearing in the dc plateau frequency region of the σ′ spectra. The Debye-type

relaxation peak suggests the transient behaviour of the complex structures of these

electrolytes. Therefore, these peaks represent the polymers segmental dynamics of the

complexed miscible blend. The intensities of these peaks vary with the PEG

concentration and also with the sample preparation methods (Figs. 3(a) and 3(b)). But the

tanδ peak intensity has an increase with increase of temperature and also shifts towards

higher frequency side (Figs. 3(c) and 3(d)) which infers that the polymer dynamics in the

complexed structures of the electrolyte increases. The relaxation time τtanδ corresponding

to these polymer blend segmental motion of the complexed structures is determined by

the relation τtanδ = 1/2πfp(tanδ), where fp(tanδ) is the frequency corresponding to tanδ peak.

The observed τtanδ values of the electrolytes are given in Tables 2 and 3 with PEG

concentration and the temperature variation, respectively. Table 2 shows that the τtanδ

values of SC prepared films are nearly same (except 15 wt% PEG). But in case of US–

MW prepared films the τtanδ value of 0 wt% PEG is found significantly different as

compared to the 10 and 15 wt% PEG concentrations. Interestingly, it is also observed that

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the τtanδ value of SC prepared 15 wt% PEG is same as that of US–MW prepared 0 wt%

PEG electrolyte film and these values are relatively very low. It is observed that the

(PEO–PMMA)–LiCF3SO3–10 wt% PEG electrolyte films prepared by both the methods

have decrease of τtanδ values with the increase of temperature (Table 3) confirming the

increase of polymer dynamics in their complexes.

Figs. 4(a) and 4(b) show the electric modulus (real part M′ and loss M″ ) spectra

of the SC prepared (PEO–PMMA)–LiCF3SO3–x wt% PEG electrolyte films of different

PEG concentration and the 10 wt% PEG concentration film at varying temperatures,

respectively. Similar types of M′ and M″ spectra are also observed for the US–MW

prepared films. These spectra are free from the contribution of EP effect, and independent

of the nature of electrode material, the electrode/dielectric specimen contact, and the

adsorbed impurities in the sample. The M′ and M″ spectra of these electrolytes have

dispersion above 100 kHz, whereas in the EP effect dominated frequency region their

values are found close to zero because of the product M*(ω)·ε*(ω) = 1. Mostly, the M″

spectra of the ion conducting electrolytes exhibit a peak corresponding to the ionic

conductivity relaxation time [2,10,15,17,46,53,56]. But in these electrolytes the M″

spectra peaks seem to appear above the upper limit of the experimental frequency range.

With an increase of temperature, the M″ dispersion has shift towards higher frequency

side which reveals that the ionic relaxation is thermally activated with the hops of charge

carriers.

The Nyquist impedance plots (Z″ vs Z′ ) of the (PEO–PMMA)–LiCF3SO3–x wt%

PEG electrolyte films prepared through SC and US–MW processes are depicted in Figs.

5(a) and 5(b), respectively. These plots are widely used for the electrochemical

characterization of the ion conducting electrolyte materials. The nature of charge carriers

(electrons or ions), the frequency range over which the EP effect contributes in bulk

properties, the behaviour of electrode–electrolyte contact, and the dc resistance and the dc

ionic conductivity of the electrolytes are generally analyzed from these plots

[2,4,12,18,34,37,41]. The ion conducting polymeric electrolyte materials commonly

exhibit a spike in the low frequency region and an arc in the high frequency region of

their Z″ vs Z′ plots. The insets of Figs. 5(a) and 5(b) show the similar behaviour of the

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studied polymeric electrolytes. The spike of an ideal capacitive element should be

parallel to the imaginary axis, but for the studied electrolyte it deviates from ideal

behaviour which is due to irregularities at the electrode/electrolyte contact. The

interfacial impedance for such materials can be described by a constant phase element

(CPE). In the parallel equivalent circuit consisting of bulk resistance Rb in parallel with

geometrical capacitance Cg, the CPE acts in series as shown in inset of Fig. 5(b). The

common intercept of the arc and the spike line on the real axis gives the bulk resistance

Rb value of the electrolyte film [2,12,18]. Further, the frequency value corresponding to

this intercept point separates the bulk and EP affected frequency region of the electrolyte

material [2–4,9,17]. The σdc values of the ion conducting electrolyte film can also be

determined using the relation σdc = tg/RbA, where tg is the thickness and A is the surface

area of the film. In the present study, the σdc values of the investigated electrolyte films

are determined by power law fit to the spectra, which are found nearly same as those

evaluated using Rb values. Figs. 5(c) and 5(d) show the temperature dependent Z″ vs Z′

plots of (PEO–PMMA)–LiCF3SO3–10 wt% PEG electrolytes prepared through SC and

US–MW methods, respectively. The insets show that these plots have shift towards low

resistance side on the real axis which confirms the decrease of their Rb values with

increase of temperature.

3.3. Correlation between ionic conductivity and dielectric parameters

Fig. 6 shows the variation of the room temperature (RT) Δε, τtanδ and σdc values

with the PEG concentration of (PEO–PMMA)–LiCF3SO3–x wt% PEG electrolyte films

prepared by SC and US–MW methods. Fig. 6 has been plotted in order to identify the

dependence of ionic conductivity on the dielectric parameters of these electrolytes. It has

been established that for the solid polymeric electrolytes, the σdc value increases with the

increase of dielectric strength and also with the decrease of polymer chain segmental

motion relaxation time [7,10,15,17,18,46,56]. In such electrolytes the transportation of

ions occurs due to segmental motion of cations coordinated polymer chain [25,36,43–45].

It is found that the σdc value of US–MW prepared film is two times higher than that of the

SC prepared film at 0 wt% PEG (Table 2). This confirms that the US–MW processing is

effective for the enhancement of ionic conductivity of the unplasticized (PEO–PMMA)–

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LiCF3SO3 electrolyte. This increase in conductivity is also favoured by the high Δε and

low τtanδ values of the US–MW electrolyte film as compared to that of the SC prepared

unplasticized electrolyte film (Fig. 6). Further, it is observed that the σdc values of the

plasticized electrolytes at a fixed PEG concentration are also governed by their Δε and

τtanδ values. The σdc values of SC prepared plasticized electrolytes are found higher than

that of without plasticizer added electrolytes, and these values have a non-monotonous

behaviour with the PEG concentration. But in case of US–MW prepared electrolytes, the

σdc values of plasticized electrolytes are found lower than that of the unplasticized

electrolytes. The contribution of PEG plasticizer is found insignificant in the

enhancement of ionic conductivity of (PEO–PMMA)–LiCF3SO3–x wt% PEG electrolytes

prepared by US–MW method, which is expected because PEG increases the crystalline

phase of polymeric electrolytes as revealed from their XRD patterns.

Here, it is worthy to compare the σdc values of the (PEO–PMMA)–LiCF3SO3 with

the PEO–LiCF3SO3 and PMMA–LiCF3SO3 electrolytes at RT in order to explore the

effect of PEO–PMMA blending. The reported σdc values of PEO–LiCF3SO3 are ~10–7–

10–8 S cm–1 [7], 3.5×10–7 S cm–1 [22], 2×10–9 S cm–1[23], 3.5×10–7 S cm–1 [24], 10–7 S

cm–1 [25], 3×10–8 S cm–1 [28] and 3.8×10–7 S cm–1 [29], which are in the range of ~10–7–

10–9 S cm–1. The variations in these values are due to the differences in salt concentration

of the electrolytes, the sample preparation methods and the evaluation of the

experimental data. These values have further increased by one to two orders of magnitude

with the addition of ethylene carbonate (EC) and propylene carbonate (PC) as plasticizers

[24,27,30], and also alumina (Al2O3), silica (SiO2) and organomodified montmorillonite

(MMT) clay as inorganic nanofiller [22–24]. Survey of literature reveals that the σdc

values of PMMA–LiCF3SO3 at RT are ~2.29×10–6 S cm–1 [36] and 1.16×10–6 S cm–1

[37], which are also increased by one to two orders of magnitude with the addition of EC

and PC plasticizers [35,37]. But the σdc values of the investigated (PEO–PMMA)–

LiCF3SO3–x wt% PEG electrolytes at RT are ~10–5 S cm–1, which are about two to three

orders of magnitude higher than that of the PEO–LiCF3SO3 electrolytes, and nearly one

order of magnitude higher than the PMMA–LiCF3SO3 electrolytes. This finding also

suggests that the PEO–PMMA blending results in increase of their ionic conductivity up

to the same order of the magnitude as that of plasticized PEO and PMMA electrolytes,

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which is very interesting and confirms its suitability as a novel solid polymer electrolyte

for the lithium ion rechargeable batteries and the other electrochromic devices. It has also

been confirmed that the PEG plasticizer has less effect on the increase of ionic

conductivity of these electrolytes, which may be due to the facts that PEO itself acts as

plasticizer for the PMMA in the PEO–PMMA blend and also there is a complete

dissociation of the salt in the blend matrix.

Fig. 7 shows the temperature dependent Δε, τtanδ and σdc values of the (PEO–

PMMA)–LiCF3SO3–10 wt% PEG films prepared by SC and US–MW methods. On

logarithmic scale these parameters have almost linear variation with reciprocal of

temperature which confirm their Arrhenius behaviour. These plots infer that the

temperature dependent σdc values are also governed by their corresponding Δε and τtanδ

values. The decrease of τtanδ values favors the increase of σdc values which is also

supported by the increase of Δε values with increasing temperature of these materials.

Mostly, the PEO based electrolytes have increase of their σdc values by one to two orders

of magnitude when the temperature increases and exceeds the PEO melting temperature

[7,13,22,24,29]. Due to increase of temperature the viscosity of electrolyte decreases and

finally the material changes into amorphous phase, which results in increase of free

volume and favourable ion conducting paths in thermally activated dynamical medium.

In the present study, the temperature dependent conductivity study has been carried out

only in the temperature range of 30–55 °C in order to determine the activation energies of

these polymers blend electrolytes.

The conductivity activation energy Eσ and dielectric relaxation time activation

energy Eτ of the (PEO–PMMA)–LiCF3SO3–10 wt% PEG electrolytes have been

determined by the Arrhenius relations σdc = σ0 exp(–Eσ/kT) and τtanδ = τ0 exp(Eτ/kT),

respectively. On the compressed scale, the σdc and τtanδ versus 1000/T plots are linear.

The slopes of these plots were used for the determination of activation energies. The

obtained values of activation energies are listed in Table 4. Because of the linear

behaviour of Δε versus 1000/T plots, the dielectric strength activation energy Eε values

are also determined using similar type of Arrhenius relation, and these are also given in

Table 4. The Eσ and Eτ values of the SC prepared electrolyte film are found slightly

higher as compared to the US–MW prepared film of the same composition material.

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Further, it is found that for these electrolytes the Eτ values are a little higher than the

respective Eσ values. The observed Eσ values are found in consistent with the other

polymeric electrolytes [7,17,18,45]. The low activation energy values of these

electrolytes suggest that there is relatively fast hopping mechanism for the ions, which is

because of the thermally activated transient coupling between the mobile cations and the

dynamical chain segmental structures of the polymers. A little difference in Eσ and Eτ

values of these electrolytes also infers that the ion transportation occurs through hopping

mechanism, and the ions have to overcome the same barrier while relaxing as well as

while conducting, as reported for the PEO based electrolytes [17,53]. This fact suggests

that in ion transportation both the segmental motion of polymer chain and hopping

motion contribute equally in coupled form for the (PEO–PMMA)–LiCF3SO3–10 wt%

PEG electrolyte.

To have better insight into the temperature dependence of τtanδ and σdc values, the

experimental data of (PEO–PMMA)–LiCF3SO3–10 wt% PEG electrolytes prepared by

SC and US–MW methods have been fitted to the Vogel–Tamman–Fulcher (VTF)

equations τtanδ = τ0exp(Ev/k(T–T0)) and σdc = σ0T1/2exp(–Ev/k(T–T0)), respectively. In the

VTF relations, τ0 and σ0 are the pre-exponential factors, Ev is the pseudo-activation

energy and T0 is the equilibrium glass transition temperature. Although, the

measurements in the present study are in narrow temperature variation range (30–55 °C),

but on the enlarged scale, these τtanδ and σdc versus 1000/T plots fitted to the VTF

equation have curved shape which confirms the presence of free volume in the

investigated electrolyte films. The VTF fitted parameters τ0, σ0, Ev and T0 are listed in

Table 5. The T0 values of the PEO–PMMA (50/50 wt%) based electrolytes were found in

good agreement to those reported earlier for the PEO–PMMA blend (20/80 wt%) based

electrolytes [44]. Further, the reasonably good fit of temperature dependence τtanδ and σdc

data demonstrates the coupling between the ionic conductivity and segmental relaxation

in the PEO–PMMA blend based electrolytes as established in earlier studies [25,45].

4. Conclusions

The detailed dielectric dispersion and electrical properties of (PEO–PMMA)–

LiCF3SO3–x wt% PEG electrolyte films prepared through SC and US–MW methods were

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reported. It was revealed that the dielectric parameters of these solid polymeric

electrolyte films change significantly with the sample preparation methods and the PEG

concentration. The ionic conductivity of unplasticized electrolyte film prepared through

US–MW method is two times high as compared to that of the SC method prepared

electrolyte film. It is found that the ionic conductivity of these electrolytes has correlation

with the dielectric strength and the polymer chain segmental motion relaxation time. This

study confirms that the ions mobility is due to cations coordinated polymer chain

segmental motion and their transportations occur by hopping mechanism. The values of

ionic conductivity and relaxation time activation energies of these electrolytes are found

in the range 0.22–0.33 eV. The VTF behaviour of temperature dependence of the

relaxation time and conductivity confirms the coupling between the ions mobility and

polymer segmental dynamics in the polymer blend electrolytes. These PEO–PMMA

blend based electrolytes have about one to two orders of magnitude higher ionic

conductivity at room temperature as compared to the PEO and PMMA based electrolytes.

This significantly enhanced conductivity at room temperature confirms the suitability of

the PEO–PMMA blend based electrolytes for the lithium ion batteries and other

electrochromic devices. The unplasticized PEO–PMMA blend electrolytes structures are

found amorphous due to blend miscibility and its complexations with the lithium cations.

But the amorphous phase gradually reduces as the PEG concentration increases in the

electrolytes, which is owing to the same backbone units of PEG and PEO molecules.

Acknowledgements

Authors are grateful to the Department of Science and Technology (DST), New

Delhi for providing the experimental facilities through research projects Nos.

SR/S2/CMP-09/2002, SR/S2/CMP-0072/2010 and the DST–FIST program. One of the

authors SC is thankful to the DST, New Delhi for the award of SERB Fast Track Young

Scientist research project No. SR/FTP/PS-013/2012.

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Table 1 Values of Bragg’s angles 2θ and crystalline peak intensities I of pristine PEO film, PEO–PMMA blend film, and (PEO–PMMA)–LiCF3SO3–x wt% PEG electrolyte films prepared by SC and US–MW methods.

x wt% PEG

2θ1

(°) I1

(counts) 2θ2

(°) I2

(counts) Pristine PEO film

0 19.22 12964 23.41 12569 PEO–PMMA blend film

0 19.26 4746 23.49 4800 SC prepared (PEO–PMMA)–LiCF3SO3–x wt% PEG

0 19.34 329 – – 5 19.12 1324 23.23 537 10 18.99 1326 23.10 549 15 19.11 1661 23.23 852

US–MW prepared (PEO–PMMA)–LiCF3SO3–x wt% PEG 0 – – – – 10 19.24 402 23.46 261 15 18.88 1495 23.09 773

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Table 2 Values of static permittivity εs, high frequency limiting permittivity ε∞, dielectric strength Δε, loss tangent relaxation time τtanδ, dc ionic conductivity σdc, and fractional exponent n of (PEO–PMMA)–LiCF3SO3–x wt% PEG electrolyte films prepared by SC and US–MW methods.

Table 3 Temperature dependent values of static permittivity εs, high frequency limiting permittivity ε∞, dielectric strength Δε, loss tangent relaxation time τtanδ, dc ionic conductivity σdc, and fractional exponent n of (PEO–PMMA)–LiCF3SO3–10 wt% PEG electrolyte films prepared by SC and US–MW methods.

x wt% PEG εs ε∞ ∆ε τtanδ

(µs) σdc × 105

(S/cm) n

SC prepared (PEO–PMMA)–LiCF3SO3–x wt% PEG electrolytes 0 2820.9 14.8 2806.1 1.40 1.01 0.91 5 4463.9 16.7 4447.2 1.19 1.46 0.92 10 3060.4 13.9 3046.5 1.21 1.09 0.88 15 1365.1 18.8 1346.3 0.35 1.62 0.96

US–MW prepared (PEO–PMMA)–LiCF3SO3–x wt% PEG electrolytes 0 3478.4 8.0 3470.4 0.34 2.02 0.92 10 2255.3 12.9 2242.4 0.93 1.07 0.89 15 2562.4 14.2 2548.2 1.03 1.19 0.90

Temperature (°C) εs ε∞ ∆ε τtanδ

(µs) σdc × 105

(S/cm) n

SC prepared (PEO–PMMA)–LiCF3SO3–10 wt% PEG electrolytes 30 3060.4 13.9 3046.5 1.21 1.09 0.88 35 3293.0 15.0 3278.0 1.16 1.19 0.89 45 3663.9 15.6 3648.3 0.81 1.64 0.93 55 4613.1 16.3 4596.8 0.40 2.66 0.98

US–MW prepared (PEO–PMMA)–LiCF3SO3–10 wt% PEG electrolytes 30 2255.3 12.9 2242.4 0.93 1.07 0.89 35 2429.0 13.0 2416.0 0.90 1.13 0.90 45 2758.3 13.5 2744.8 0.68 1.57 0.90 55 3391.8 15.0 3376.8 0.41 2.12 0.97

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Table 4 Values of conductivity activation energy Eσ, relaxation time activation energy Eτ and dielectric strength activation energy Eε of (PEO–PMMA)–LiCF3SO3–10 wt% PEG electrolyte films prepared by SC and US–MW methods.

Table 5 The VTF fitted parameters from the temperature dependence of the conductivity and the dielectric relaxation time of (PEO–PMMA)–LiCF3SO3–10 wt% PEG electrolyte films prepared by SC and US–MW methods.

Electrolytes preparation methods

Eσ (eV)

Eτ (eV)

Eε (eV)

SC 0.27 0.33 0.12 US–MW 0.22 0.25 0.12

σ VTF fit τ VTF fit Electrolytes preparation methods σ0

(S cm–1) Ev

(eV) T0 (K)

τ0 (s)

Ev (eV)

T0 (K)

SC 3.00×10–7 0.315 351.07 1.97×10–6 0.155 339.22US–MW 1.70×10–7 0.857 379.50 1.40×10–6 0.137 340.85

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Fig. 1. XRD patterns of (a) PEO, PMMA, PEO–PMMA blend films, and LiCF3SO3 powder; (b) (PEO–PMMA)–LiCF3SO3–x wt% PEG electrolyte films prepared by SC and US–MW methods.

10 15 20 25 30

15 wt%

10 wt%

0 wt%US-MW

0 wt%

5 wt%

15 wt%

10 wt%

In

tens

ity (a

.u.)

2θ ( o )

(b) SC

10 15 20 25 30

10 15 20 25 30

PMMAPEO

PEO-PMMA

LiCF3SO3

Inte

nsity

(a.u

.)

2θ ( o )

(a)

PMMA

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0

2000

4000

ε''

0

4000

8000

0 10 15

x wt% PEG

ε'

(b)

105 1060

60

120

0 2000 40000

2000

4000

ε''

ε'

0

2000

4000

ε''

0

4000

8000

0; 5 10; 15

x wt% PEG

ε'

105 1060

100

200

0 3000 60000

3000

6000

ε''

ε'

(a)

101 102 103 104 105 106

0

3000

6000

ε''

f (Hz)

0

5000

10000

30 oC

35 oC

45 oC

55 oC

ε'

(c)

105 1060

100

200

0 2000 40000

2000

4000

ε''

ε'

101 102 103 104 105 106

0

2000

4000

ε''

f (Hz)

(d)

0

4000

8000

30 οC

35 οC

45 οC

55 οC

ε'

105 1060

100

200

0 2000 40000

2000

4000

ε''

ε'

Fig. 2. Frequency dependent real part ε′ and loss ε″ of the complex dielectric function of (a) SC prepared and (b) US–MW prepared (PEO–PMMA)–LiCF3SO3–x wt% PEG films at 30 °C; and (c) SC prepared (d) US–MW prepared (PEO–PMMA)–LiCF3SO3–10 wt% PEG films with temperature variation.

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Fig. 3. Frequency dependent real part of ac conductivity σ' and loss tangent (tanδ) of (a) SC prepared and (b) US–MW prepared (PEO–PMMA)–LiCF3SO3–x wt% PEG films at 30 °C; and (c) SC prepared (d) US–MW prepared (PEO–PMMA)–LiCF3SO3–10 wt% PEG films with temperature variation.

0.0

1.5

3.0

4.5 0 5 10 15

tan

δ

10-8

10-7

10-6

10-5

x wt% PEG

σ' (S

cm

-1)

(a)

0

2

4

6 0 10 15

tan

δ

10-7

10-6

10-5

σ' (S

cm

-1)

x wt% PEG

(b)

101 102 103 104 105 1060.0

2.5

5.0 30 oC

35 oC

45 oC

55 oC ta

n δ

f (Hz)

10-7

10-6

10-5

σ' (S

cm

-1)

(d)

101 102 103 104 105 106

1.5

3.0

4.5

30 oC

35 oC

45 oC

55 oC

tan

δ

f (Hz)

10-7

10-6

10-5

σ' (S

cm

-1)

(c)

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Fig. 4. Frequency dependent real part M′ and loss M″ of complex electric modulus of (a) SC prepared (PEO–PMMA)–LiCF3SO3–x wt% PEG films at 30 °C; and (b) SC prepared (PEO–PMMA)–LiCF3SO3–10 wt% PEG films with temperature variation.

101 102 103 104 105 106

0.00

0.01

0.02

0.03

M''

f (Hz)

0.00

0.01

0.02

30 oC

35 oC

45 oC

55 oC

M'

(b)

101 102 103 104 105 106

0.00

0.01

0.02

0.03

f (Hz)

0 5 10 15

x wt% PEG

M''

0.000

0.006

0.012

0.018

M'

(a)

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Fig. 5. Complex impedance plane plots (Z″ vs Z′) of (a) SC prepared and (b) US–MW prepared (PEO–PMMA)–LiCF3SO3–x wt% PEG films at 30 °C; and (c) SC prepared (d) US–MW prepared (PEO–PMMA)–LiCF3SO3–10 wt% PEG films with temperature variation.

0 5 10 15 20 25

0

5

10

15

20

25 0 5 10 15

Z" (k

Ω)

x wt% PEG

0.1 0.2 0.3

0.1

0.2

(a)

0 3 6 9 12 15

0

5

10

15x wt% PEG

0 10 15

Z" (k

Ω)

(b)

0.1 0.2 0.30.0

0.1

0.2

0 4 8 12 16

0

3

6

9

30 oC

35 oC

45 oC

55 oC

Z" (k

Ω)

Z' (kΩ)

(c)

0.1 0.2 0.3

0.1

0.2

0 5 10 15 20

0

4

8

12

16 30 oC

35 oC

45 oC

55 oC

Z"

(kΩ

)

Z' (kΩ)

(d)

0.1 0.2 0.30.0

0.1

0.2

Rb

Cg CPE

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Fig. 6. PEG concentration dependent dielectric strength Δε, loss tangent relaxation time τtanδ and dc ionic conductivity σdc of (PEO–PMMA)–LiCF3SO3–x wt% PEG electrolyte films prepared by SC and US–MW methods.

0 5 10 150.5

1.0

1.5

2.0

x wt % PEG

0.5

1.0

1.5

σ dcx

105 (S

cm

-1)

τ tanδ

(μs)

103

2x103

3x1034x1035x103

SC US-MW

Δε

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Fig. 7. Reciprocal temperature dependence of dielectric strength Δε, loss tangent relaxation time τtanδ and dc ionic conductivity σdc of (PEO–PMMA)–LiCF3SO3–10 wt% PEG electrolyte films prepared by SC and US–MW methods.

3.1 3.2 3.3

10-5

3x10-5

SC US-MW

1000/T (K-1)

10-1

100

σ dc (S

cm

-1)

τ tanδ

(μs)

2x103

4x103

6x1038x103

Δε

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Fig. 8. Temperature dependent plots of the relaxation time and the ionic conductivity of (PEO–PMMA)–LiCF3SO3–10 wt% PEG electrolyte films prepared by SC and US–MW methods. The continuous lines show the fit of experimental data to the VTF equations.

3.1 3.2 3.310-5

2x10-5

1000 / T (K-1)

5x10-7

10-6

1.5x10-6

SC US-MW VTF fit VTF fit

σ dc (S

cm

-1)

τ tanδ

(s)

role of preparation methods on the structural and dielectric properties of plasticized polymer blend...

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