Role of Preparation Methods on the Structural and Dielectric

7/21/2019 Role of Preparation Methods on the Structural and Dielectric

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Electrochimica Acta 142 (2014) 359370

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Electrochimica Acta

journal homepage: www.elsevier .com/ locate /e lectacta

Role ofpreparation methods on the structural and dielectricproperties ofplasticized polymer blend electrolytes: Correlationbetween ionic conductivity and dielectric parameters

R.J. Sengwa, Priyanka Dhatarwal, Shobhna Choudhary

DielectricResearch Laboratory, Department of Physics, Jai Narain VyasUniversity, Jodhpur342005, India

a r t i c l e i n f o

Article history:

Received 18 May 2014Received in revised form 22 July 2014Accepted 23 July 2014Available online 8 August 2014

Keywords:

Polymer blendSolid polymer electrolyteDielectric relaxationIonic conductivity

a b s t r a c t

The polymer blend based electrolyte films consisted of poly(ethylene oxide) (PEO) and poly(methylmethacrylate) (PMMA) with lithium triflate (LiCF3SO3) as a dopant ionic salt and poly(ethylene glycol)(PEG) as plasticizer have been prepared by solution cast meltpressed and ultrasonic assisted followedby microwave irradiated solution cast meltpressed methods. The Xray diffraction study infers thatthe amorphous phase of(PEOPMMA)LiCF3SO3x wt% PEG electrolytes decreases with the increase ofPEG concentration. The complex dielectric function, ac electrical conductivity, electric modulus and theimpedance spectra ofthese electrolytes have been investigated over the frequency range from 20Hz to1MHz. Itis found that allthe dielectric/electricalfunctions oftheseelectrolytes vary anomalouslywith theincrease ofPEG concentration, and also with the change ofsamples preparation methods. The activationenergies have been determined from the temperature dependent values ofdc ionic conductivity, polymersegmental relaxation time and dielectric strength. Results reveal that besides the amorphicity, the ionicconductivity ofthese electrolytes is also governed by the relaxation time and the dielectric strength, andthe transport ofions is due to hopping mechanism which is coupled with segmental motion ofpolymerschain. Room temperature ionic conductivity values ofthe PEOPMMA blend based electrolytes are foundabout one to two orders ofmagnitude higher than that ofthe PEO and PMMA based electrolytes.

2014 Elsevier Ltd. All rights reserved.

1. Introduction

Solid polymer electrolytes (SPEs) have been recognized as themost suitable flexible type, leak proof andlightweight novel mate-rials for the fabrication of ion conducting devices particularly therechargeable batteries for portable electronic equipments [118].The SPEs are complexes of polymer with cations of the dopant saltand mostly have high amorphicity. But the low ionic conductivityat ambient temperature limits their several technological appli-cations. In these electrolytes the dynamics of polymer chains iscritical for the ions transportation. Therefore, the knowledge onhow the ions are coupled with the chain dynamics of the multi-phase complex structures of the SPEs and the modification in thesestructures with the preparation methods for the improvement ofionic conductivity are nowadaysprime issues [5,6,12,15,16,1921] .

So far, enormous work is in progress on the SPE materials withthechoicesofpolymersandthealkalimetalsaltsinordertoachieve

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

their high ionic conductivity ( 105 S cm1) at ambient temper-ature for the realistic device applications. Mostly, poly(ethyleneoxide) (PEO) matrix based complexes with various lithium saltshave been an intense field of research [517,1930]. The PEO haslow lattice energy, low glass transition temperature, high solvat-ing power for alkali metal salts and an ease to form flexible typefilm. These facts make it an impending polymer in preparation ofthe SPEs. But the high crystallinity of PEO is its main disadvantagebecause effective iontransport can be achieved utmostin an amor-phous medium of the host polymer. Besides the PEO, poly(methylmethacrylate) (PMMA) matrix, which has amorphicity more than96%, is also preferredin preparation of SPEmaterials [3,3138]. Thehighly active electron donating carbonyl (C= ) functional groupof PMMA can easily coordinate with the cations of alkali metalsalts, which results in PMMAsalt complexes. Although PMMA haslight weight withhightransparency, highstrength anddimensionalstability, but its brittle property under a loaded force limits theindustrial and technological suitability of the PMMA based SPEs.

In the advancement of SPE materials, continuous work is inprogress on the PEOPMMA blend matrix based electrolytes toovercome the drawbacks of the pristine polymers [3948]. The

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360 R.J. Sengwaet al./ ElectrochimicaActa 142 (2014) 359370

advantages of this blend are; the addition of PEO in PMMA matrixresults in the increase of PMMA flexibility and reduces its brittle-ness, whereas the PMMA environment increases the amorphousphaseofPEO.ThesepropertiesofthePEOPMMAblendmatrixhavebeen recognized its suitability as novel polymeric blend matrix oftailormade properties for the preparation of SPEs. Our survey ofliterature reveals that PEOPMMA matrix is used with lithium tri-flate [39], lithium perchlorate [4045] and silver nitrate [4648]as ion conducting salt with different plasticizers and inorganicnanofillers for the enhancement of the ionic conductivity of theirsolution cast electrolyte films. But the PEOPMMA matrix withlithium trifluoromethanesulfonate (lithium triflate, LiCF3SO3) asion conducting salt and poly(ethylene glycol) (PEG) as plasticizerhas not been studied for its possible applications in the advance-ment of ion conducting electrochromic devices. As compared tothe high dielectric constant plasticizers ethylene carbonate (EC)s = 85.1 and propylene carbonate (PC) s = 69.0, at room temper-ature, the PEG has s = 22.1, which is very low. So it is interestingto know how a low dielectric constant plasticizer affects the con-ductivity of the electrolytes. Further, PEGis a linear chain moleculeand itsrepeat unit is same as PEO backbone unit andtherefore,howthe PEOstructure varies with increase of PEG concentration, is alsointeresting to study in these electrolytes.

In this paper, we have first time prepared the PEOPMMA blendmatrix based electrolytes comprising LiCF3SO3 salt and varyingconcentrations of PEG by classical solution cast meltpressed andultrasonic assisted followed by microwave irradiated solution castmeltpressed methods. The aim of ultrasound treatment is, firstly,to enhance the PEOPMMA blend miscibility, and secondly to dis-sociate the salt clusters in the electrolyte solution, if present any.The microwave irradiation is given to the solution with the aimto modify the dipolar orientation and the PEG dynamics whichoccur at microwave frequencies in the solutions. This study hasbeen carried outto explore the effects of samplepreparationmeth-ods and plasticizer concentration on the dielectric properties, ionicconductivity and structural dynamics of the SPE films using dielec-tric relaxation spectroscopy. Further, an attempt has been made

to correlate the temperature dependent ionic conductivity to thepolymer chain dynamics (segmental relaxation) by fitting the datato the VogelTammanFulcher (VTF) equation of these polymerelectrolytes.

2. Experimental

2.1. Sample preparation

ThePEO(Mw = 6105 gmol1), PMMA (Mw =3.5105 gmol1),PEG (Mw =200gmol1) and LiCF3SO3 were obtained fromSigmaAldrich, USA. The anhydrous acetonitrile and tetrahydro-furan of spectroscopic grade were purchased from Loba Chemie,

India. The polymers and the salt were vacuum dried at 50

C for atleast 12h before using. For the preparation of varying PEG concen-trations (PEOPMMA)LiCF3SO3x wt% PEG electrolyte films, the50:50 w/w PEO:PMMA blend was used. The average molar ratio9:1 was set for the total number of the ethylene oxide units (EO)and thecarbonyl groups(C = O)of the polymers blendto the lithiumcations (Li+) of the salt. The PEG concentrationsx= 0 , 5 , 10 and 15with respect to the weight of PEOPMMA blend were used.

Two different processing methods were used for preparationof the electrolyte films. (i) Classical solution casting (SC) method:In this process, initially, the required amounts of PEO and PMMAwere dissolved in acetonitrile and tetrahydrofuran, respectively,in separate glass bottles. After that LiCF3SO3 was added in PEOsolution, and it was dissolved and mixed homogenously by mag-

netic 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 concentrationswere added into the PEOPMMA blend electrolyte solutions andmixed homogeneously by magnetic stirring. Each PEG concentra-tion homogenous solution was divided into two equal parts. Thefirst parts were cast onto Teflon petri dishes and by slow evap-oration of solvent at room temperature, the solution cast (SC)prepared (PEOPMMA)LiCF3SO3x wt% PEG films were achieved.(ii) Ultrasonic (US) assisted followed by microwave (MW) irradi-ated solution casting method: In this method, the second partsof each polymeric electrolyte solutions of varying PEG concentra-tions,which were prepared as mentionedabove, were sonicatedbyultrasonicator (250W power, 25kHz frequency) for 10min dura-tion with ONOFF step of 15 s. In this processing the stainlesssteel sonotrode was directly immersed into the electrolyte solu-tion forstrongdose of theultrasound. After that each of thesolutionwas irradiated by microwave electromagnetic energy in a domes-tic microwave oven (600W power, 2.45GHz frequency) for 2 minduration and 10 s irradiation step with intermediate cooling. Thesesolutions were cast onto petri dishes whichresultedin the USMWirradiated solution cast films after the solvents evaporation.

The surfaces of the solution cast films prepared by above men-tioned methodswerefound uneven. Therefore, thesmooth surfaces

of these films were achieved by meltpressed technique. In thistechnique, each electrolyte film was initially vacuum dried at 40 Cfor 24h. After that each film was melted by heating it up to 130 Cin circular stainless steel die having suitable spacer using polymerfilm making unit. This melted material was pressed under 2 tonsof pressure per unit area and cooled slowly up to room tempera-ture which made the film of smooth surfaces. The same steps wererepeated for each film.

2.2. Characterizations

The Xray diffraction (XRD) patterns of the SPE films and theirconstituents were recorded in reflection mode using a PANalyticalXpert Pro MPD diffractometer of Cu K radiation (1.5406A) oper-

ated at 45kV and 40mA with a scanned step size of 0.05/s. Thepowder sample of LiCF3SO3was tightly filled in the sample holder,whereas the PEO, PMMA, PEOPMMA blend andthe SPEfilms wereplacedon the topof sampleholder duringtheir XRDmeasurementsin the 2 range from 1030 at room temperature.

The dielectric relaxation spectroscopy (DRS) of the electrolytefilms was carried out using Agilent technologies 4284A precisionLCR meter along with Agilent 16451B solid dielectric test fix-ture in 1V electric field of linear frequency f range from 20Hzto 1 MHz at 30 C, and also with the temperature variation forthe 10 wt% PEG concentration films. Frequency dependent val-ues of capacitance Cp, resistance Rpand loss tangent (tan=/)of the SPE films mounted in the dielectric cell were measured inthe parallel circuit operation for the determination of their dielec-

tric/electrical spectra. Prior the sample measurement, the opencircuit calibration of the cell was performed to eliminate the effectof stray capacitance of the cell leads. The dielectric test fixture wasplaced in a microprocessor-controlled heating chamber to recordthe measurements at different temperatures. Initially, all the mea-surements of varying PEG concentrations electrolyte films weremade at fixed temperature, 30C. After that the same films of 10wt% PEG concentration were used for their temperature depend-encestudyfrom 3055 C.Itwas found thatthe re-measured valuesofCp, Rpand tan have good reproducibility after one week dura-tion, which confirms that there is no immediate aging effect inthe studied electrolyte films. The spectra of intensive quantities,namely complex dielectricfunction*() =j, realpart ofalter-nating current (ac) electrical conductivity =0

and electric

modulus M*

() =M

+jM

, and the extensive quantity i.e. complex


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R.J. Sengwaet al./ ElectrochimicaActa 142 (2014) 359370 361

10 15 20 25 30

15 wt%

10 wt%

0 wt%US-MW

0 wt%

5 wt%

15 wt%

10 wt%

Intensity(

a.u.)

2

(

o

)

(b)SC

10 15 20 25 30

10 15 20 25 30

PMMA

PEO

PEO-PMMA

LiCF3SO

3

Intensity(a.u.)

2

(

o

)

(a)

PMMA

Fig. 1. XRD patterns of (a) PEO, PMMA and PEOPMMA blend films, and LiCF3SO3 powder; (b) (PEOPMMA)LiCF3SO3x wt% PEG electrolyte films prepared by SC andUSMW methods.

impedanceZ*() =ZjZ of the electrolyte films were determinedusing the expressions described in detail elsewhere [3,9].

3. Results and Discussion

3.1. Structural analysis

The XRD patterns of PEO, PMMA, PEOPMMA blend, LiCF3SO3,and the SC and USMW prepared (PEOPMMA)LiCF3SO3x wt%

PEGelectrolytes over the angular range 1030 are shown in Fig. 1.The XRD pattern of a material provides the structural propertiesrelated to crystalline phases (peak positions), phase concentra-tion (peak heights), amorphous content (back-ground hump) andcrystallite size (peak widths). The PEO peaks position 2 and thecorresponding intensity values of these materials are determinedby Xpert pro software and these values are given in Table 1. TheXRD pattern of pure PEO (Fig. 1(a)) has sharp crystalline peaks at2=19.22 and 23.41, which are corresponding to crystal reflec-tionplanes 120 and concerted112,032, respectively [8,15,49]. InsetofFig.1(a) shows that PMMA has a broad and diffused peak around16, which confirms its predominantly amorphous phase [3]. ThePEOPMMA blend also has peaks at 19.19 and 23.33 which arecorresponding to PEO crystalline reflection planes, but the intensi-

ties of these peaks are low as compared to pristine PEO (Table 1).The relative low intensity peaks of PEOPMMA blend reveal thatthere is some miscibility of PEO and PMMA which increases theamorphous content, but some amount of PEO crystalline phasealso exists in the blend. The LiCF3SO3has intense crystalline peaksat 16.62, 19.87, 20.53, 22.77 and 24.71, which confirm its highcrystallinity. The angular positions and relative intensities of thesepeaks arefoundin consistent with the earlier reported XRDspectraof the LiCF3SO3[50].

The differences in XRD patterns of (PEOPMMA)LiCF3SO3xwt% PEG films prepared by SC and USMW processing methods(Fig. 1(b)) confirm some structural variations due to the samplepreparation methods. The crystalline peaks of LiCF3SO3 are notfound 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 PEOPMMA blend. Fur-ther, the absence of PEO peaks in the USMW prepared film anda very low intensity peak in the SC prepared film of the unplasti-cized (PEOPMMA)LiCF3SO3 electrolyte infer that the formationof iondipolar complexes has changed the polymer blend intoan amorphous material. Due to four to six coordination sites oflithium cations with the oxygen atoms of PEO [10,22], and thecarbonyl groups of PMMA [32] and also formation of some mis-cible phase in the PEOPMMA blend result in complete amorphous

phase of the (PEOPMMA)LiCF3SO3 electrolyte having the saltconcentration molar ratio [EO + (C= O)]:Li+ =9:1. With the addi-tion of PEG as plasticizer in these electrolytes both the peaksof PEO have appeared, which confirm the recrystallization of asmall amount of PEO. The enhancement of PEO peaks intensi-ties 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 mate-rial crystallinity [8,49]. In comparison to the SC prepared films thepeak intensities of USMW prepared films are low(Table 1), whichconfirm that the high intensity ultrasonication disturbs the recrys-tallization of PEO. The end hydroxyl groups of PEG form the strongintra- and inter-molecular hydrogen bonding [51]. It seems thatthe hydroxyl groups of PEGform the ion-dipolar coordination with

the lithium cations which release some of the ether oxygen atomsof PEO from the polymer-ion complexes. Due to this fact someamount of uncomplexed PEO recrystallizes. As the PEG concentra-tion increases in the electrolyte, more amountof PEOreleases fromthe complexes, which results in gradual increase of PEO crystallinephase. Further, the presence of same unit in the backbone of PEGand PEOmoleculesalso favours theenhancement of PEOcrystallinephase as the PEG concentration increases in the PEOPMMA blendbased electrolytes. This structural finding of the PEG plasticizedelectrolyte is very interesting, and so far this type of behaviour hasnotbeenobserved inothertypeof plasticizer addedPEO basedsolidpolymer electrolytes [27,30,39,41,46] . Further, authors repeatedthe XRD scans of some of the samples within one week duration,and found good reproducibility of the scans which confirmed thatthe samples did not suffer from immediate aging effect.


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Table 1

Values of Braggs angles 2and crystalline peak intensities Iof pristine PEO film, PEOPMMA blend film, and (PEOPMMA)LiCF3SO3x wt% PEG electrolyte films preparedby SC and USMW methods.

x wt% PEG 21() I1(counts) 22() I2 (counts)

Pristine PEO film0 19.22 12964 23.41 12569

PEOPMMA blend film0 19.26 4746 23.49 4800

SC prepared (PEOPMMA)LiCF3SO3x wt% PEG electrolytes

0 19.34 329 5 19.12 1324 23.23 53710 18.99 1326 23.10 54915 19.11 1661 23.23 852

USMW prepared (PEOPMMA)LiCF3SO3x wt% PEG electrolytes0 10 19.24 402 23.46 26115 18.88 1495 23.09 773

3.2. Dielectric and electrical behaviour

The complex permittivity (real part and dielectric loss )spectra of (PEOPMMA)LiCF3SO3xwt%PEG electrolytefilms pre-pared through SC and USMW processes are shown in Figs. 2(a)and 2(b), respectively. The values of these electrolytes are in

the order of several thousands at low frequencies which decreasenonlinearly with increasing frequency and approach the steadystate (high frequency limiting permittivity ) near 1 MHz, asshown by the enlarged view in the inset ofFig. 2. The large val-ues of these polymeric electrolytes at low frequencies are due tothe electrode polarization (EP) effect occurring as a result of theaccumulation of ions near the electrodes surfaces of dielectric testfixture [9,10,15]. These spectra have point of inflection around1 kHz, therefore the values at this frequency are considered asthe static permittivity s of these electrolytes. The sand val-ues of the polymeric electrolytes with PEG concentration are givenin Table 2. The values of dielectric strength =s of thesematerials are also recorded in Table 2.

As compared to the polymer matrix, the plasticizers have high

valueofstaticpermittivity.Therefore,bythealgebraicadditiverule,it is expected that the increase of plasticizer concentration in thepolymeric electrolyte must increase the dielectric strength besidethe modification in material physical properties such as flexibil-ity, viscosity, microstructure, etc. Further, it is believed that theincrease of dielectric strength of the polymer electrolyte with theaddition of plasticizer must support the ion conduction process.But the dielectric strength of a composite material is governed bythe dipolar ordering in its complex structure. The parallel dipolarordering increases the dielectric strength whereas it is reduced incase of anti-parallel dipolar ordering. The comparison ofFigs. 2(a)and 2(b) and also the values given in Table 2 reveals thatthe values of these electrolyte films change anomalously withthe increase of PEG concentration and these values are also influ-enced by the sample preparation methods. These results confirm

that there is significant alteration in the strength of iondipolarinteractions and the dipolar ordering in the complexes withincrease of PEG concentration. The non-monotonous behaviour ofwith PEGconcentrationalso reveals that the interaction of end-hydroxyl groups of PEG with lithium cations and functional groupsof polymersmakes themrandomly aligned in thecomplexes, which

is also favoured by the changes in their crystalline phase (Table 1).Figs. 2(c) and 2(d) show the temperature dependent and

spectra of theSC andUSMWprepared (PEOPMMA)LiCF3SO310wt% PEG films. It is found that both the and values increasewith increasing temperature, which is attributed to the increasein charge density as an additional contribution from the interfa-cial polarization. The temperature dependents, and valuesof the 10% PEG concentration electrolyte film prepared by SC andUSMW methods are recorded in Table 3. It has also been revealedthat the magnitude of values of these electrolyte films followstheir values.Further,eachofthe spectrumhassinglerelaxationpeak corresponding to the segmental motion (-relaxation) in themiscible polymer blend. Due to large difference in the backboneunits of PEO and PMMA chains there must be separate relaxation

peaks in the

spectra on the frequency scale, but the appearanceof single relaxation peak also suggests a cooperative chain segmen-tal dielectric relaxation process in these PEOPMMA blend basedelectrolytes. The ColeCole plots ( vs ) of the electrolyte filmsare shown in insets of the spectra ofFig. 2. The high frequencydata in these plots hasappearedas semicirculararcs correspondingto the bulk properties whereas the low frequency data forms thespikesdue to dominancecontributionof theEP effectin the lowfre-quency dielectric properties. The values of these electrolytes asdetermined from their estimated svalues are in consistence withthe other polymeric electrolytes which are determined by fittingthe complex permittivity data to the HavriliakNegami function[18].

The ac ionic conductivity and the loss tangent tan spectra ofSC and USMW processed prepared (PEOPMMA)LiCF3SO3xwt%

Table 2

Values of static permittivity s , high frequency limiting permittivity , dielectric strength , loss tangent relaxation time tan, dc ionic conductivitydc , and fractionalexponent n of (PEOPMMA)LiCF3SO3x wt% PEGelectrolytefilms prepared by SC and USMW methods.

x wt% PEG s tan(s)

dc 105

(Scm1)n

SC prepared (PEOPMMA)LiCF3SO3x wt% PEG electrolytes0 2820.9 14.8 2806.1 1.40 1.01 0.915 4463.9 16.7 4447.2 1.19 1.46 0.9210 3060.4 13.9 3046.5 1.21 1.09 0.8815 1365.1 18.8 1346.3 0.35 1.62 0.96

USMW prepared (PEOPMMA)LiCF3SO3x wt% PEG electrolytes0 3478.4 8.0 3470.4 0.34 2.02 0.9210 2255.3 12.9 2242.4 0.93 1.07 0.8915 2562.4 14.2 2548.2 1.03 1.19 0.90


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0

2000

4000

'

'

0

4000

8000

010

15

x wt% PEG

'

(b)

105

106

0

60

120

0 200040000

2000

4000

''

'

0

2000

4000

'

'

0

4000

8000

0; 5

10; 15

x wt%PEG

'

105

106

0

100

200

0 3000 60000

3000

6000

'

'

'

(a)

101

102

103

104

105

106

0

3000

6000

''

f (Hz)

0

5000

10000

30oC

35oC

45oC

55oC

'

(c)

105

106

0

100

200

0 2000 40000

2000

4000

'

'

'

101

102

103

104

105

106

0

2000

4000

''

f(Hz)

(d)

0

4000

8000

30 C

35 C

45 C

55 C

'

105

106

0

100

200

0 2000 40000

2000

4000

'

'

'

Fig. 2. Frequency dependentreal part andloss of thecomplex dielectric function of (a)SC prepared and (b)USMWprepared(PEOPMMA)LiCF3SO3x wt% PEG filmsat 30 C; and (c) SC prepared (d)USMW prepared (PEOPMMA)LiCF3SO310 wt% PEG films with temperature variation.

PEG electrolytes are depicted in Figs. 3(a) and 3(b), respectively.

The

values of these polymeric electrolytes increase nonlinearlywith increase of frequency on logarithm scale. It has been estab-lished that the ion transportation in polymeric electrolytes occurson different time scale under the influence of alternating current(ac) electric field [8,17,18,52,53]. The jumps of ions between dif-ferent ion sites through hopping mechanism are demonstratedby random barrier model which consequences the ac conductiv-ity dispersion [52]. The shorttime ion dynamics is characterizedby backandforth motion over the limited range in disorderedpolymer matrix, subdiffusive dynamics, which leads to disper-siveac conductivity at high frequencies. Whereas thelongtime iondynamicsis characterizedby random walks resultingin longrangeiontransport,diffusivedynamics,whichleadstotheplateaucorre-spondingto thedirect current (dc)ionic conductivity at frequencieslower than that of the dispersive ac conductivity region. The dc

conductivity of such electrolytes mostly obeys the Arrhenius tem-

perature dependent behaviour [7,17]. The ac conductivity of thepolymericelectrolytesalsoobeys timetemperature superposition,i.e. it is possible to scale data at different temperatures to one sin-gle master curve which is roughly same for all disorder materials[18,52,54]. Further, at high frequencies, the loglog plot of ac con-ductivityfollowsapparentpower law behaviour [55]. Thereforethedetailed studies of frequency dependent ac conductivity behaviourof SPEs have academic interest, besides the confirmation of theirsuitability for technological applications.

The spectra of the investigated electrolytes show threeregions on frequency scale; (i) the electrode polarization domi-nated lowfrequency regionbelow10 kHz, (ii) thedc plateau regionabove 30kHz up to nearly 300 kHz, and (iii) the power law disper-sive region above 300 kHz. The values of dc ionic conductivitydcof these electrolytes are determined by fitting the

spectra above


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0.0

1.5

3.0

4.5 05

10

15

tan

10-8

10-7

10-6

10-5

x wt% PEG

'(Scm

-1)

(a)

0

2

4

6 010

15

tan

10-7

10-6

10-5

'(Scm

-1)

x wt% PEG

(b)

101

102

103

104

105

106

0.0

2.5

5.0 30 oC

35oC

45oC

55oC

tan

f(Hz)

10-7

10-6

10-5

'(Scm

-1)

(d)

101

102

103

104

105

106

1.5

3.0

4.5

30oC

35oC

45oC

55oC

tan

f (Hz)

10-7

10-6

10-5

'(Scm

-1)

(c)

Fig. 3. Frequency dependentreal part of ac conductivityand loss tangent (tan) of (a) SC prepared and (b) USMW prepared (PEOPMMA)LiCF3SO3x wt% PEG filmsat30 C; and (c) SC prepared (d) USMW prepared (PEOPMMA)LiCF3SO310 wt% PEGfilms with temperature variation.

30kHz to the conventional Jonschers power law() =dc +An

[55], whereA is the preexponential factor and n is the fractional

exponent ranging between 0 and 1. These fits are shown by solidlines in the spectra. Further, it is observed that the increaseswith the increasing temperature of the electrolytes (Figs. 3(c) and3(d)). This increase has twoimplications, firstly, the mobility of theions increases due to increased polymer chain segmental motion,and secondly, the ion concentration increases. But the XRD spec-tra reveals that the total salt is in dissociated form and hence theincrease of temperature does not have any additional effect in dis-sociating the salt. This fact confirms that the increase of valuesof these electrolytes is due to increase of ions mobility only. Fur-ther, the presence of small salt clusters of nm size is also ruled outbecause the salt concentration is not very high. The evaluated val-ues ofdcand n of the electrolytes with the PEG concentration andthe 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 elec-trolyte also takes place through hopping mechanism as reported

for other polymeric electrolytes [10,15,17,18,52,56] . Further, it isfound that the n values of the electrolyte increase with increaseof temperature (Table 3). This agrees well with the idea that nphysically represents the strength of the ionion and iondipolarinteractions and as the temperature increases these interactionsare expected to decrease [25].

The tan spectra of the electrolytes shown in Fig. 3 have Debyetype relaxation peaks appearingin the dc plateau frequency regionof the spectra. TheDebye-type relaxationpeak suggeststhe tran-sient behaviour of the complex structures of these electrolytes.Therefore, these peaks represent the polymers segmental dynam-ics of the complexed miscible blend. The intensities of these peaksvary with the PEG concentration and also with the sample prepara-tion methods (Figs. 3(a) and 3(b)). But the tan peak intensity hasan increase with increase of temperature and also shifts towards


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101 102 103 104 105 106

0.00

0.01

0.02

0.03

M''

f(Hz)

0.00

0.01

0.02

30oC

35oC

45oC

55oC

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)

Fig. 4. Frequency dependent real partM and loss M of complex electric modulus of (a) SC prepared (PEOPMMA)LiCF3SO3x wt% PEG filmsat 30C; and (b) SC prepared(PEOPMMA)LiCF3SO310 wt% PEGfilm with temperature variation.

higher frequency side (Figs. 3(c) and 3(d)) which infers that thepolymer dynamics in the complexed structures of the electrolyteincreases.The relaxation timetan correspondingto thesepolymerblend segmental motion of thecomplexedstructures is determinedby the relation tan =1/2fp(tan), where fp(tan) is the frequencycorresponding to tan peak. The observed tanvalues of the elec-trolytes are given in Tables2 and 3 with PEGconcentration and thetemperature variation, respectively. Table 2 shows that the tan

values of SC prepared films are nearly same (except 15 wt% PEG).Butin case of USMW prepared films thetan valueof0wt%PEGisfound significantlydifferent as compared to the 10 and 15 wt% PEGconcentrations. Interestingly, it is also observed that thetanvalueof SC prepared 15 wt% PEG is same as that of USMW prepared 0wt% PEG electrolyte film and these values are relatively very low.It is observed that the (PEOPMMA)LiCF3SO310 wt% PEG elec-trolyte films prepared by both the methods have decrease oftanvalues with the increase of temperature (Table 3) confirming theincrease of polymer dynamics in their complexes.

Figs. 4(a) and 4(b) show the electric modulus (real part M andloss M) spectra of the SC prepared (PEOPMMA)LiCF3SO3x wt%PEGelectrolytefilms of differentPEG concentration and the 10 wt%PEG concentration film at varying temperatures, respectively. Sim-

ilar types ofM

and M

spectra are also observed for the USMWprepared films. These spectra are free from the contribution of EPeffect,andindependentofthenatureofelectrodematerial,theelec-trode/dielectric specimen contact, and the adsorbed impurities inthesample. TheM andM spectra of theseelectrolytes havedisper-sion above 100 kHz, whereas in the EP effect dominated frequencyregion their values are found close to zero because of the prod-uct M*()*()= 1. Mostly, the M spectra of the ion conductingelectrolytes exhibit a peak corresponding to the ionic conductivityrelaxation time [2,10,15,17,46,53,56] . But in these electrolytes theM spectrapeaksseemtoappearabovetheupperlimitoftheexper-imental frequency range. With an increase of temperature, theM

dispersion has shift towards higher frequency side which revealsthat the ionic relaxation is thermally activated with the hops ofcharge carriers.

The Nyquist impedance plots (Z vs Z) of the(PEOPMMA)LiCF3SO3x wt% PEG electrolyte films preparedthrough SC and USMW processes are depicted in Figs. 5(a) and5(b), respectively. These plots arewidely used forthe electrochem-ical characterization of the ion conducting electrolyte materials.The nature of charge carriers (electrons or ions), the frequencyrange over which the EP effect contributes in bulk properties, thebehaviour of electrodeelectrolyte contact, and the dc resistance

and the dc ionic conductivity of the electrolytes are generallyanalyzed from these plots [2,4,12,18,34,37,41] . The ion conductingpolymeric electrolyte materials commonly exhibit a spike in thelow frequency region and an arc in the high frequency region oftheirZ vsZ plots. The insets ofFigs. 5(a) and 5(b) showthe similarbehaviour of the studied polymeric electrolytes. The spike of anideal capacitive element should be parallel to the imaginary axis,but for the studied electrolyte it deviates from ideal behaviourwhich is due to irregularities at the electrode/electrolyte contact.The interfacial impedance for such materials can be describedby a constant phase element (CPE). In the parallel equivalentcircuit consisting of bulk resistance Rbin parallel with geometricalcapacitance Cg, the CPE acts in series as shown in inset ofFig. 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 thisintercept point separates the bulk and EP affected frequencyregion of the electrolyte material [24,9,17]. The dc values ofthe ion conducting electrolyte film can also be determined usingthe relation dc = tg/RbA, where tg is the thickness and A is thesurface area of the film. In the present study, the dcvalues of theinvestigated electrolyte films are determined by power law fit tothe spectra, which are found nearly same as those evaluated usingRb values. Figs. 5(c) and 5(d) show the temperature dependentZ vs Z plots of (PEOPMMA)LiCF3SO310 wt% PEG electrolytesprepared through SC and USMW methods, respectively. Theinsets show that these plots have shift towards low resistance sideon the real axis which confirms the decrease of their Rb valueswith increase of temperature.


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0 5 10 15 20 25

0

5

10

15

20

250

5

10

15

Z"(k

)

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

)

(b)

0.1 0.2 0.30.0

0.1

0.2

0 4 8 12 16

0

3

6

9

30oC

35oC

45oC

55oC

Z"(k

)

Z' (k )

(c)

0.1 0.2 0.3

0.1

0.2

0 5 10 15 20

0

4

8

12

16

30oC

35oC

45oC

55oC

Z"(k

)

Z' (k )

(d)

0.1 0.2 0.30.0

0.1

0.2

Rb

C gCPE

Fig. 5. Complex impedance plane plots (Z vs Z) of (a) SC prepared and (b) USMW prepared (PEOPMMA)LiCF3SO3x wt% P EG films at 30C; and (c) SC prepared (d)USMW prepared (PEOPMMA)LiCF3SO310 wt% PEGfilms with temperature variation.

3.3. Correlation between ionic conductivity and dielectric

parameters

Fig. 6 shows the variation of the room temperature (RT), tan and dc values with the PEG concentration of (PEOPMMA)LiCF3SO3x wt% PEG electrolyte films prepared bySC and USMW methods. Fig. 6 has been plotted in order toidentify the dependence of ionic conductivity on the dielectricparameters of these electrolytes. It has been established that forthe solid polymeric electrolytes, the dc value increases with theincrease of dielectric strength and also with the decrease of poly-mer chainsegmental motion relaxation time [7,10,15,17,18,46,56] .In such electrolytes the transportation of ions occurs due to seg-mentalmotionofcationscoordinatedpolymerchain[25,36,4345].It is found that the dc value of USMW prepared film is twotimes higher than that of the SC prepared film at 0 wt% PEG(Table 2). This confirms that the USMW processing is effective

for the enhancement of ionic conductivity of the unplasticized(PEOPMMA)LiCF3SO3 electrolyte. This increase in conductivityis also favoured by the high and lowtanvalues of the USMW

electrolyte film as compared to that of the SC prepared unplasti-cized electrolyte film (Fig. 6). Further, it is observed that the dcvalues of the plasticized electrolytes at a fixed PEG concentrationare also governed by their and tan values. The dc values ofSC prepared plasticized electrolytes are found higher than that ofwithout plasticizer added electrolytes, andthese values havea non-monotonous behaviour with the PEG concentration. But in case ofUSMW prepared electrolytes, the dc values of plasticized elec-trolytes are found lower than that of the unplasticized electrolytes.The contribution of PEG plasticizer is found insignificant in theenhancement of ionic conductivity of (PEOPMMA)LiCF3SO3xwt% PEG electrolytes prepared by USMW method, which isexpected because PEG increases the crystalline phase of polymericelectrolytes as revealed from their XRD patterns.


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0 5 10 150.5

1.0

1.5

2.0

x wt % PEG

0.5

1.0

1.5

dcx

10

5(

Scm

-1)

tan(

s)

103

2x103

3x103

4x103

5x103

SC

US-MW

Fig. 6. PEG concentration dependent dielectric strength , loss tangent relax-ationtime tananddc ionic conductivitydcof (PEOPMMA)LiCF3SO3x wt% PEGelectrolyte films prepared by SC andUSMW methods.

Here, it is worthy to compare the dc values of the (PEOPMMA)LiCF3SO3 with the PEOLiCF3SO3 andPMMALiCF3SO3 electrolytes at RT in order to explore theeffect of PEOPMMA blending. The reported dc values ofPEOLiCF3SO3 are 10

7108 S cm1 [7], 3.5107 S cm1 [22],

2109 S cm1[23], 3.5107 S cm1 [24], 107 S cm1 [25],3108 S cm1 [28] and 3.8107 S cm1 [29], which are in therange of107109 S cm1. The variations in these values aredue to the differences in salt concentration of the electrolytes,the sample preparation methods and the evaluation of the exper-imental data. These values have further increased by one to twoorders of magnitude with the addition of ethylene carbonate(EC) and propylene carbonate (PC) as plasticizers [24,27,30], andalso alumina (Al2O3), silica (SiO2) and organomodified montmo-rillonite (MMT) clay as inorganic nanofiller [2224]. Survey ofliterature reveals that thedcvalues of PMMALiCF3SO3at RT are2.29106 S cm1 [36] and 1.16106 S cm1 [37], which arealso increased by one to two orders of magnitude with theadditionof EC and PC plasticizers [35,37]. But the dc values of the inves-

tigated (PEOPMMA)LiCF3SO3x wt% PEG electrolytes at RT are105 S cm1, which are about two to three orders of magnitudehigher than that of the PEOLiCF3SO3electrolytes, and nearly oneorder of magnitude higher than the PMMALiCF3SO3electrolytes.This finding also suggests that the PEOPMMA blending resultsin increase of their ionic conductivity up to the same order ofthe magnitude as that of plasticized PEO and PMMA electrolytes,which is very interesting and confirms its suitability as a novelsolid polymer electrolyte for the lithium ion rechargeable batteriesand the other electrochromic devices. It has also been confirmedthat the PEG plasticizer has less effect on the increase of ionicconductivity of these electrolytes, which may be due to the factsthat PEO itself acts as plasticizer for the PMMA in the PEOPMMAblend and also there is a complete dissociation of the salt in the

blend matrix.

3.1 3.2 3.3

10-5

3x10-5

SC

US-MW

1000/T (K-1

)

10-1

100

dc

(Scm

-1)

tan(

s)

2x103

4x103

6x103

8x103

Fig. 7. Reciprocal temperature dependence of dielectric strength , loss tangentrelaxation time tan and dc ionic conductivity dc of (PEOPMMA)LiCF3SO310wt% PEGelectrolyte films prepared by SC andUSMW methods.

Fig. 7 shows the temperature dependent , tan and dcvalues of the (PEOPMMA)LiCF3SO310 wt% PEG films pre-pared by SC and USMW methods. On logarithmic scale theseparameters have almost linear variation with reciprocal of tem-perature which confirm their Arrhenius behaviour. These plotsinfer that the temperature dependent dc values are also gov-erned by their corresponding and tan values. The decreaseoftanvalues favors the increase ofdc values which is also sup-

portedby theincreaseof values withincreasing temperature ofthese materials. Mostly, the PEO based electrolytes have increaseof their dc values by one to two orders of magnitude when thetemperature increases and exceeds the PEO melting temperature[7,13,22,24,29]. Duetoincreaseoftemperaturetheviscosityofelec-trolyte decreases and finally the material changes into amorphousphase, which results in increase of free volume and favourable ionconducting paths in thermally activated dynamical medium. In thepresent study, the temperature dependent conductivity study hasbeen carried outonly in thetemperaturerangeof 3055 C inorderto determine the activation energies of these polymers blend elec-trolytes.

The conductivity activation energy Eand dielectric relaxationtime activation energy E of the (PEOPMMA)LiCF3SO310 wt%

PEG electrolytes have been determined by the Arrhenius relationsdc =0exp(E/kT) and tan = 0exp(E/kT), respectively. On thecompressed scale, thedcand tanversus 1000/T plots are linear.The slopes of these plots were used for the determination of activa-tion energies. The obtained values of activation energies are listedin Table 4. Because of the linear behaviour of versus 1000/Tplots, the dielectric strength activation energy E values are alsodetermined using similar type of Arrhenius relation, and these arealso given in Table 4. The Eand Evalues of the SC prepared elec-trolyte film are found slightly higher as compared to the USMWprepared filmof the same composition material.Further,it is foundthat for these electrolytes the Evalues are a little higher than therespective E values. The observed E values are found in consis-tent with the other polymeric electrolytes [7,17,18,45]. The low

activation energy values of these electrolytes suggest that there is


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Table 3

Temperature dependentvaluesof static permittivitys, highfrequency limitingpermittivity, dielectricstrength, losstangentrelaxationtime tan, dc ionicconductivitydc , and fractional exponent n of (PEOPMMA)LiCF3SO310 wt% PEG electrolyte films prepared by SC and USMW methods.

Temperature (C) s tan(s) dc 105 (Scm1) n

SC prepared (PEOPMMA)LiCF3SO310 wt% PEG electrolyte30 3060.4 13.9 3046.5 1.21 1.09 0.8835 3293.0 15.0 3278.0 1.16 1.19 0.8945 3663.9 15.6 3648.3 0.81 1.64 0.9355 4613.1 16.3 4596.8 0.40 2.66 0.98

USMW prepared (PEOPMMA)LiCF3SO310 wt% PEGelectrolyte30 2255.3 12.9 2242.4 0.93 1.07 0.8935 2429.0 13.0 2416.0 0.90 1.13 0.9045 2758.3 13.5 2744.8 0.68 1.57 0.9055 3391.8 15.0 3376.8 0.41 2.12 0.97

Table 4

Values of conductivity activationenergyE, relaxationtimeactivationenergyE anddielectric strengthactivation energyEof(PEOPMMA)LiCF3SO310 wt%PEG electrolytefilms prepared by SC and USMW methods.

Electrolyte preparation methods E(eV) E(eV) E(eV)

SC 0.27 0.33 0.12USMW 0.22 0.25 0.12

Table 5

TheVTF fitted parametersfrom thetemperature dependenceof the dc ionicconductivity and thedielectricrelaxation timeof (PEOPMMA)LiCF3SO310 wt% PEGelectrolytefilms prepared by SC and USMW methods.

Electrolytes preparation methods VTF fit VTF fit

0(Scm1) Ev(eV) T0(K) 0(s) Ev(eV) T0(K)

SC 3.00107 0.315 351.07 1.97106 0.155 339.22USMW 1.70107 0.857 379.50 1.40106 0.137 340.85

relatively fast hopping mechanism for the ions, which is becauseof the thermally activated transient coupling between the mobilecations and the dynamical chain segmental structures of the poly-mers. A littledifference inEandEvalues of these electrolytes alsoinfers that the ion transportation occurs through hopping mecha-nism,andtheionshavetoovercomethesamebarrierwhilerelaxingas well as while conducting, as reported for the PEO based elec-trolytes [17,53]. This fact suggests that in ion transportation both

3.1 3.2 3.3

10-5

2x10-5

1000 / T (K-1

)

5x10-7

10-6

1.5x10

SC

US-MW

VTF fit

VTF fit

dc

(Scm

-1

)

tan(

s)

Fig. 8. Temperature dependent plots of the relaxation time and the ionic conduc-tivity of (PEOPMMA)LiCF3SO310 wt% PEG electrolyte films prepared by SC andUSMW methods. Thecontinuouslines show thefit of experimentaldatato theVTF

equations.

the segmental motion of polymer chain and hopping motion con-tribute equally in coupled form for the (PEOPMMA)LiCF3SO310wt% PEG electrolyte.

To have better insight into the temperature depend-ence of tan and dc values, the experimental data of (PEOPMMA)LiCF3SO310 wt% PEG electrolytes pre-pared by SC and USMW methods have been fitted to theVogelTammanFulcher (VTF) equations tan =0exp(Ev/k(TT0))and dc =0T

1/2exp(Ev/k(TT0)), respectively. In the VTF rela-tions, 0 and 0 are the pre-exponential factors, Ev is thepseudo-activation energy andT0is the equilibrium glass transitiontemperature. Although, the measurements in the present studyare in narrow temperature variation range (3055C), but onthe enlarged scale, these tan and dc versus 1000/T plots fittedto the VTF equation have curved shape (Fig. 8) which confirmsthe presence of free volume in the investigated electrolyte films.The VTF fitted parameters 0, 0, Ev and T0 are listed in Table 5.The T0 values of the PEOPMMA (50/50 wt%) based electrolyteswere found in good agreement to those reported earlier for thePEOPMMA blend (20/80 wt%) based electrolytes [44]. Further, thereasonably good fit of temperature dependence tanand dcdata

demonstrates the coupling between the ionic conductivity andsegmental relaxation in the PEOPMMA blend based electrolytesas established in earlier studies [25,45].

4. Conclusions

The detailed dielectric dispersion and electrical propertiesof (PEOPMMA)LiCF3SO3x wt% PEG electrolyte films preparedthrough SC and USMW methods were reported. It was revealedthat the dielectric parameters of these solid polymeric electrolytefilms change significantly with the sample preparation methodsand the PEG concentration. The ionic conductivity of unplasticizedelectrolyte film prepared through USMW method is two times

high as compared to that of the SC method prepared electrolyte


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