NaBr complexed polymer blend electrolytes for electrochemical cell applications

Journal of Membrane Science 454 (2014) 200–211

Contents lists available at ScienceDirect

Journal of Membrane Science journal homepage: www.elsevier.com/locate/memsci

Investigations on PEO/PVP/NaBr complexed polymer blend electrolytes for electrochemical cell applications K. Kiran Kumar a,b,n, M. Ravi b, Y. Pavani b, S. Bhavani b, A.K. Sharma b, V.V.R. Narasimha Rao b a b

GITAM University, Department of Engineering Physics, GIT, Visakhapatnam 530045, AP, India Sri Venkateswara University, Department of Physics, Tirupati 517502, AP, India

art ic l e i nf o

a b s t r a c t

Article history: Received 21 June 2013 Received in revised form 2 November 2013 Accepted 9 December 2013 Available online 17 December 2013

Experimental investigations on a sodium ion conducting polymer blend electrolyte system based on polyethylene oxide (PEO) and polyvinyl pyrrolidone (PVP), complexed with NaBr salt are presented in this paper. The complexed polymer blend electrolytes were obtained in the form of dimensionally stable and free-standing films by solution cast technique. Physical characterization by X-ray diffraction (XRD), Fourier transform infrared spectroscopy (FTIR), scanning electron microscopy (SEM), differential scanning calorimetry (DSC) and thermogravimetric analysis (TGA) were performed to study the structural properties, their change along with ion–polymer interactions and thermal stability of the samples. The transport numbers of different mobile species were determined by means of d.c. polarization and combined a.c./d.c. techniques. Cationic transportation was found to be dominant in the films of high degree of amorphicity. Electrical properties were measured as a function of composition and temperature using complex impedance spectroscopy. The temperature dependent electrical conductivity showed Arrhenius type behavior. Activation energies were found to decrease with increasing concentration of the salt. Electrochemical cells were prepared using PEO/PVP/NaBr polymer blend electrolyte system and their discharge characteristics were studied. These studies suggested the practical realization of electrochemical cells using the present sodium ion conducting polymer blend electrolyte system. Crown Copyright & 2013 Published by Elsevier B.V. All rights reserved.

Keywords: XRD FTIR SEM DSC Activation energy

1. Introduction Development and characterization of new kinds of materials with uncommon properties have been drawing considerable attention in recent years. Among these new ‘strategic’ materials, conductive polymers, namely systems combining plastic nature with high electric transport have reached a prominent stage in a number of academic and industrial laboratories. Particularly, interest on solid polymer electrolytes (SPEs) has grown during the last decades, together with the search of new and more efficient, high performance devices for energy conversion and storage. The large interest in this field is due to the fact that these materials open the route for new and exciting applications such as realization of functionalized sensors, development of advanced rechargeable batteries, electrochemical cells, electro-chromic display devices, smart windows, fuel cells, photo electrochemical solar cells, dye-sensitized solar cells etc [1–5]. Solid polymer electrolytes have several advantages over conventional liquid electrolytes such as improved safety, appreciable theoretical

n Corresponding author at: GITAM University, Department of Engineering Physics, GIT, Visakhapatnam 530045, AP, India. Mobile: þ91 80085 28025. E-mail address: [email protected] (K.K. Kumar).

capacity, leakage free nature, low cost and easy fabrication into flexible geometries [6–8]. For most potential applications, it is desirable that the solid polymer electrolytes display a reasonable ionic conductivity, dimensional stability, processibility and flexibility under ambient conditions. These requirements have been achieved by adopting several approaches on various polymers using different methods of preparation. As a result, a variety of polymer electrolytes such as polymer–salt complexes, gel polymer electrolytes, composite polymer electrolytes and blend based polymer electrolytes have emerged. Polyethers, particularly polyethylene oxide (PEO) based electrolytes are the earliest and the most extensively studied systems that are usually referred to as solvent-free PEO/salt complexes. The high ionic transport in these polymer systems was first discussed by Wright [9]. Later Armand [10] demonstrated the feasibility of their use in electrochemical cells based on polyethylene oxide/Lisalt complexed solid polymer electrolyte films. Consequently, a large number of polymer electrolyte materials involving different kinds of transporting ions, namely, H þ , Li þ , Na þ , K þ , Ag þ , Mg2 þ etc. were reported. PEO is the most appropriate base material due to its ability to play host to various metal–salt systems for a wide range of concentrations. It also acts as a good binder for other phases and has excellent chemical stability. PEO has been complexed with a wide range of Li salts: LiX (X ¼I, Cl, Br, ClO4, SCN,

0376-7388/$ - see front matter Crown Copyright & 2013 Published by Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.memsci.2013.12.022

K.K. Kumar et al. / Journal of Membrane Science 454 (2014) 200–211

CF3SO3, BF4, ASF6, etc.) [11–14], of which PEO LiCF3SO3 [12,13] and PEO–LiClO4 [13,14] have been most widely studied. Polyethylene oxide, is a semicrystalline material with crystalline and amorphous elastomeric phases. Complex formation in PEOn–salt (n is the number of ether oxygen per mole of salt) is governed by competition between salvation energy and lattice energy of the polymer and the inorganic salt [15]. Low lattice energies of both the polymer and the complexing salt have been found to increase stabilities in the resulting polymer electrolyte. In PEO based Na–X complexes, PEO backbone was found as an open helix in which the oxygen atoms are directed inward [16]. In PEO salt complexes it has been revealed that the cation remains encapsulated within the helix of the PEO chain with the anions stacked outside the helix, thereby separating the cations from the X  counter anions [17]. This favors (i) the dissolution of salt in the PEO matrix following the salvation mechanism and (ii) the possibility of the anion to migrate within the polymer electrolyte which would deteriorate the performance of the battery leading to self-discharge as well as degradation of the electrode surface [18,19]. However, the cation transport in polymer electrolytes, a consequence of local relaxation as well as segmental motion of the polymer chains, is more favorable in the presence of high degree of amorphicity in the host polymer. The applications of PEO salt complexes in battery fabrication have been restricted by their low cationic transport numbers which is the consequence of the semicrystalline nature of PEO at ambient temperatures. Hence most of the research efforts have been directed towards the attainment of films containing large and stable amorphous phases in order to obtain good flexibility of the polymer chains which are responsible for the cation transport and to impede anion migration. Complexes of PEO with aromatic anions such as phenols and naphthols were prepared, but no ionic conductivities and transport numbers were reported [20,21]. Anions such as Cl  and I  have been complexed with azabased compounds, which, apart from acting as anion receptors, gave rubber like characteristics to the otherwise stiff polymer–salt complexes [22]. Attempts to impede the mobility of the anions were also made by synthesizing new structures with graft polymers, block copolymers and cross-linked polymer networks [23–25]. Comb polymers based on methyl vinyl ether/maleic anhydride copolymer containing oligo-oxyethylene, stable upto 140 1C, were studied by Ding et al. [26,27]. However, some of these polymer electrolytes showed poor ionic conductivities. One of the most promising alternate choices of enhancing the amorphous content in the polymer electrolyte system is polymer blending [28] and may result in highly dimensionally stable polymer electrolytes. Interest in these materials has arisen principally from their ability to modify mechanical properties and to tailor properties for desired applications. The inherent merits of using blend based polymer electrolytes are exemplified by several research groups [29–32]. It is because of these promising characteristics, that the present authors thought it worthwhile to investigate PEO based polymer blend electrolyte system to study its structural, electrical and electrochemical characteristics for battery applications. Polyvinyl pyrrolidone (PVP) has been chosen as a partner of PEO to prepare the polymer blend electrolyte system for electrochemical cell applications as it is a conjugated polymer with high amorphous content which can permit faster ionic mobility compared to other semicrystalline polymers. The presence of carbonyl group (CQO) lends marked Lewis base character to the side chains of PVP, leading to the formation of a variety of complexes with various inorganic salts. It also exhibits high Tg with good environmental stability. Furthermore, it provides outstanding thermal stability and mechanical strength to the blend material due to its ability to cross-link thermally. In the field of solid polymer electrochemical cells, so far, the main thrust of the work has been directed towards systems based on lithium, perhaps, mainly because of the larger number of available lithium insertion materials with well described properties. However lithium batteries are still facing some challenges and

201

safety limitations [33]. A few studies on Mg and Zn-battery systems, which are able to replace Li-batteries are reported [34,35]. Sodium also provides a compelling rationale for the study of sodium/ polymer battery systems as another alternative system, due to its lower cost, natural abundance, low toxicity and low atomic mass (23.0). Sodium is the most attractive as a negative electrode reactant on account of its high electrochemical reduction potential of -2.71 V. When coupled with an appropriate electropositive material, it is capable of giving a cell voltage of >2 V. The combination of low mass and high voltage leads to the possibility of employing sodium as electrode material in a battery to achieve high specific energy [36]. Sodium-ion batteries have been shown to be technically possible and have the potential advantage that sodium compounds are cheaper than lithium ones [37]. The chemistry of sodium is significantly different from lithium, particularly in the solubility of sodium salts in organic solvents. Furthermore, the mobility of smaller ions (Li þ and/or Mg2 þ ) is lower than that of cations with larger size (Na þ and/or Zn2 þ ) in polymer electrolytes [38]. The conductivity data obtained from them also supported the hypothesis that smaller cations are embedded or captured by the polymeric network and their mobility is lowered. The softness of sodium makes it easier to achieve and maintain contact with other components in electrochemical devices. Moreover, it does not form alloy with most of the metals, which may be used as current collectors in electrochemical devices. A few attempts have been made to develop electrolytes based on sodium complexed films [39–41]. However the studies available on PEO based blends complexed with sodium salts are limited, particularly for separate measurements of tcation and tanion with induced changes in crystallinity, which is an important parameter for controlling the transport properties of polymer electrolytes for battery applications. After reviewing a number of Li þ and Mg2 þ based electrolyte systems, we have used the experience gained to investigate sodium based PEO/ PVP blend system with the aim of making laboratory cells which can prove that the development of these systems is a realistic goal.

2. Experimental details 2.1. Materials Polyethylene oxide (PEO) of molecular weight 4  106 and polyvinyl pyrrolidone (PVP) of molecular weight 1.3  105 were obtained from Aldrich and used without further purification to prepare solid polymer blend electrolytes. The structure of polyethylene oxide (–CH2CH2O–), viewed perpendicular to the crystalline helix is shown in Scheme. 1. The helix conformation of PEO which is the basis of the structural unit in the crystalline phase has two turns in a fiber identity period of 1.93 nm. This polymer, in its

H

H

C

C

H

H

Hydrogen atom

O

Carbon atom

n

Oxygen atom

Scheme 1. Structural unit of PEO and model of PEO structure viewed perpendicular to the crystalline helix.

202

K.K. Kumar et al. / Journal of Membrane Science 454 (2014) 200–211

H

H

C

C

H

N

O n

Carbon Oxygen Nitrogen Hydrogen

Scheme 2. Structural unit of PVP and a meso dyad in the tt conformation along PVP chain.

pure form, is chemically and electrochemically stable since it contains only strong unstrained C–O, C–C and C–H bonds [42]. Polyvinyl pyrrolidone (PVP) is a non-ionic polymer having a hydrocarbon chain with strong polar side groups. A meso-dyad in the tt conformation along PVP chain is shown in Scheme 2. Sodium bromide (NaBr) (extra pure) salt (Sd. Fine-Chem Limited, India) was used as the dopant. Methanol (Merck, India) was used as the common solvent. Circular disks of aluminum electrodes (area¼1.168 cm2) were used as blocking electrodes in the electrochemical measurements. 2.2. Preparation of polymer electrolytes Polymer blend electrolyte films were prepared by using solution cast technique. In this process initially polyethylene oxide and polyvinyl pyrrolidone were added to methanol, in the required wt %, while stirring the solution magnetically at room temperature (  30 1C) for complete dissolution. Further, appropriate amounts (5, 10 and 15 wt%) of NaBr salt were dissolved in methanol and added to the PEO/PVP polymer solution under continuous stirring for 10–12 h. Finally, viscous solutions were poured into polypropylene dishes and the common solvent methanol was allowed to evaporate slowly at room temperature to obtain free-standing polymer electrolyte films at the bottom of dishes. The films were dried at  40 1C for 6 h to remove any traces of the residual solvent in the electrolytes. The harvested films were stored in highly evacuated desiccators to avoid any moisture absorption. Films of pure blend of PEO/PVP (70:30) and various compositions of NaBr doped PEO/PVP were obtained in weight percent ratios (67.5:27.5:5, 65:25:10, 62.5:22.5:15). Thickness of these films was determined by mechanical stylus method using German made ‘Perthometer’ and was found to be approximately 160 μm with an accuracy of about 75 μm. 2.3. Structural investigations In order to investigate the structure of the polymer blend electrolyte films, X-ray diffraction scans were taken using, PANalytical X'Pert PRO, X-ray diffractometer with CuKα radiation (λ ¼1.540 A1) and graphite monochromator at room temperature. The diffraction patterns were recorded at the Bragg's angle (2θ) in the range 10–801 with a scan rate of 0.11 per 6 s. FTIR absorption spectra were recorded using ‘Thermo Nicolet IR200’ Spectrometer in the wavenumber range of 400–4000 cm  1. FTIR experiments were performed in dynamic nitrogen atmosphere by averaging 16 scans per sample, keeping an optical resolution of 4 cm  1 for all

the spectra. The surface morphology of the polymer blend films was observed using CARL ZEISS EVO 25 scanning electron microscope. The samples were gold coated at 10 mA current under 10  2 Torr vacuum for 6 min using the sputter coater prior to imaging. Differential scanning calorimetry (DSC) thermograms of pure PEO/PVP and NaBr complexed PEO/PVP blend polymer electrolytes were recorded using Perkin Elmer PYRIS Diamond differential scanning calorimeter with a low temperature measuring head and liquid nitrogen as coolant. Samples of  5 70.1 mg weight were crimped in aluminum pans inside the glove box under argon atmosphere and transferred to the DSC cell for measurements. Samples in aluminum pans were stabilized by slow cooling to  70 1C and then heated to 300 1C with reference to the empty pan at a rate of 10 1C/min. The data was recorded during the heating scan. A Perkin Elmer TGA-7 (vertical furnace) system was used for the TGA experiments.

2.4. Electrical characterization The total ionic/electronic transport number of polymer blend electrolytes was measured by the well known Wagner's polarization method [43]. In this technique, polymer electrolyte is sandwiched between an electrode of the parent metal, Na, which acts as a reversible (non-blocking) electrode, and an ion-blocking electrode such as graphite (C) in the configuration Na|electrolyte| C. A d.c. step potential E, of 1.5 V, is applied to the asymmetric cell. The accuracy of the applied voltage was about 70.05 V and the resulting static current was monitored as a function of time using a Keithley (model 617) programmable electrometer. The transport numbers (tion, tele) were calculated from the polarization current versus time plots using the equations t ion ¼ 1  I f =I i

ð1Þ

t ele ¼ I f =I i

ð2Þ

where Ii is the initial current and If, the final current. Cationic/anionic transport numbers were evaluated using the combination of complex impedance and a dc polarization measurement technique suggested by Watanabe et al. [44]. In this technique, the polymer blend samples were sandwiched between two reversible Na electrodes in the cell geometry Na/electrolyte sample/Na. Sodium formed the non-blocking electrode for Na þ ions while it acted as a blocking electrode for the anions. The complex impedance was measured at 315 K using computer controlled phase sensitive multimeter (N4L PSM 1700) (impedance analyzer) in the frequency range 1 Hz to 1 MHz. The cell was then subjected to d.c. polarization measurement. A constant d.c. potential of 0.5 V was applied to the cell and the time dependence of the current was monitored by a high input impedance Keithley 617 electrometer, which gave the value of Is, the steady state current. This procedure was repeated for different constant potentials, 0.75 V, 1 V, 1.25 V and 1.5 V. Every time the steady state current Is was measured. The ionic conductivity of the complexed polymer blend electrolytes was determined by a.c. impedance measurements using a computer controlled meter (Newton 4th Ltd. UK) PSM 1700 (phase sensitive multimeter) in the frequency range 1 Hz to 1 MHz. The samples were vacuum dried at 313 K for 1 h and subsequently the measurements were recorded by sandwiching the electrolyte film between two electrodes which were assembled in a temperature controlled INDFURR furnace. The temperature dependence of ionic conductivity was measured by heating the samples from room temperature (303 K) to 348 K. Aluminum was selected as an electrode material due to its low work function and inertness with the components of the complexed polymer blend electrolyte films.

K.K. Kumar et al. / Journal of Membrane Science 454 (2014) 200–211

2.5. Fabrication of electrochemical cells Solid state electrochemical cells were fabricated by sandwiching the polymer blend film as electrolyte between anode and cathode pellets, with the configuration, Na|(PEO/PVP/NaBr)|(I2 þCþelectrolyte). This entire assembly was finally enclosed in the sample holder. The cells were prepared using the blend electrolytes complexed with different wt% of NaBr salt. Sodium metal was used as the anode. A proper mixture of iodine (I2), graphite (C) and the electrolyte material in weight ratio of 5:5:1 respectively was obtained by physical grinding. This was then pressed in the form of thin pellet/disc as the cathode at a pressure of 445 MPa after proper mixing of constituents. While iodine is an active cathodic material, graphite provides an adequate electronic conductivity and electrolyte material to favor electrode/electrolyte interfacial contact and helps in reducing electrode polarization [45].

3. Results and discussion 3.1. XRD analysis Qualitative identification of structure can be made by comparing the position of X-ray diffraction peaks of the specimen pattern with an index of standard patterns. XRD patterns of pure PEO, PVP and PEO/PVP blend films are shown in Fig. 1. The angular position of characteristic peaks of pure PEO [including their d-spacing given by the instrument and miller indices (PCPDF File nos. 49-2200 and 492201)] and amorphous peaks of pure PVP are given in the inset. PEO (Fig. 1(a)) exhibits a maximum intensity peak at 23.721, next maximum peak at 19.361 and a relatively less intense peak at 26.491 which are assigned to (112), (120) [46] and (222) planes [PCPDF File nos. 49-2200 and 49-2201] respectively. The d-spacings of the set of planes (120) and (112) of high intensity peaks at 4.65 A1 and 3.79 A1 have been reported earlier [47]. Several low intense peaks were also found around 12.61, 15.21, 22.51, 25.41, 29.71, 36.61 and 39.61 which are correlated with those of PCPDF File nos. 49-2200 and 49-2201. The sharp peaks are attributed to the crystalline phase of PEO, which originates from the ordering of polyether side chains due to the strong intermolecular interaction between PEO chains through the hydrogen bonding. The crystalline structure of PEO is a monoclinic unit cell [48]. The crystalline peaks of PEO were found to overlap on a low intensity broad peak of amorphous phase of PEO

203

between 181 and 451. PVP (Fig. 1(b)) shows two broad peaks at around 131 and 211, which are attributed to the amorphous nature of PVP [49–51]. The later is superimposed by the maximum intensity crystalline peak (23.721) of PEO, hence it appears as a shoulder of the sharp crystalline PEO peak in the XRD pattern of pure PEO/PVP blend (Fig. 1(c)). These observations confirm that the present polymer blend system possesses multiphase, having both crystalline and amorphous regions. The relative intensities of two crystalline peaks (19.361 and 23.721) of pure PEO/PVP blend are different from those of pure PEO which indicates that the PEO crystals in the blend are different from those formed in pure PEO polymer. Fig. 2 is a comparative study of XRD patterns of pure PEO/PVP (a), its complexes with various salt concentrations (b)–(d) and NaBr salt (e). The comparative study reveals that the intensity of all crystalline peaks of PEO decreases gradually, upon the addition of salt to the polymer blend, suggesting a decrease in the degree of crystallinity of the complexes. This could be due to the disruption of the semicrystalline structure of the film by salt. When NaBr dissolves in the polymer host, the interaction between PEO/PVP host matrix and NaBr leads to a decrease of the intermolecular interaction among the polymer chains which reduces the crystalline phase and hence increases the amorphous region. The relative intensities of sharp peaks of PEO at around 191 and 231 in the complexed PEO/PVP blend with 15 wt% NaBr are different from those of pure PEO/PVP blend which indicates that the dopant inhibits the orientation of PEO crystallites preferentially in certain directions [45]. The sharp peaks corresponding to NaBr salt (Fig. 2(e)) observed between 17.271 and 68.611 (PCPDF Card nos. 01-0901, 74-1182 and 72-1539) disappeared in the polymer blend complexes, which indicates the complete dissolution of salt in the polymer matrix [52].

3.2. FTIR analysis The IR spectra of pure PEO/PVP, complexes of PEO/PVP with different concentrations of NaBr and NaBr salt are shown in Fig. 3. The allowed modes of vibrations of various functional groups of PEO and PVP, correspond to the observed IR characteristic bands of pure polymer blend film which confirm the presence of PEO and PVP. These are given in Table 1 and the relevant references have been given in our earlier paper [53]. The co-existence of well resolved bands of ether oxygen groups (C–O–C) of PEO and carbonyl groups (CQO) of PVP indicates that PEO and PVP are miscible.

(e)

(b) (b) 1 →13°{Pos. [°2Th.]} , 2 → 21°{ Pos. [°2Th.]} 1

(a)

2

2 6

10

5

1

8 9 7

11

(a)

No. 1. 2. 4. 5. 6. 7. 8. 9. 10. 11.

Pos. [°2Th.] 15.64032 19.36874 23.7226 26.49313 27.1767 31.29217 32.85425 35.58744 36.51183 39.92792

d-spacing [Å] (h k l) 5.66129 (013) 4.5791 (120) 3.74763 (112) 3.36167 (222) 3.27864 (111) 2.85619 (142) 2.72387 (044/135) 2.52069 (313) 2.45896 (313) 2.2561 (118)

(d)

Intensity(arbitrary units)

Intensity(arbitrary units)

(c)

(c)

(b)

4 10

20

(a) 30

40

50

60

70

80

2θ Fig. 1. XRD patterns of pure (a) PEO, (b) PVP and (c) PEO/PVP blend films. [Inset: Angular positions and miller indices of crystalline peaks of pure PEO and angular positions of two broad amorphous peaks of pure PVP].

10

20

30

40

50

60

70

80

2θ Fig. 2. XRD patterns of PEO/PVP blend films containing NaBr ratios (a) 0 wt%; (b) 5 wt%; (c) 10 wt%; (d) 15 wt%; (e) NaBr salt.

204

K.K. Kumar et al. / Journal of Membrane Science 454 (2014) 200–211

- (CH2 - CH2 -O)n- + Na-Br →

- (CH2 - CH2

CH2 - CH2)nO

Na+- - - Br Transmittance %

Scheme 3. Formation of cross-link due to the interaction of cation with PEO ether oxygen atom.

3000

2500

2000

1500

1000

C

c

C

C

500

Wavenumber (Cm-1) Fig. 3. FTIR spectra of PEO/PVP blend films containing NaBr ratios (a) 0 wt%; (b) 5 wt%; (c) 10 wt%; (d) 15 wt% (e) NaBr salt.

Table 1 Different vibrational modes of various functional groups of PEO and PVP correspond to observed IR characteristic peaks of pure blend polymer [53]. Vibrational mode of functional groups

PEO wavenumber (cm  1)

PVP wavenumber (cm  1)

Symmetric C–H stretching Asymmetric C–H stretching of CH2 C–H stretching/C–C and C–O–C stretching CH2 bending second over tone C–H stretching/CH2 deformation combination C–H stretching/CQO stretching combination Asymmetric C–H stretching/C–H deformation combination C–O stretching C–N stretching Symmetric and asymmetric CQO stretching C–H stretching or C–O stretching second over tone combination C–N stretching CH2 scissoring mode C–H bending of CH2 CH2 wagging

3800–2700 2962–2695 2510 2357 2328 2240 2165

 2900, 2560 2510 2357 2328 2240 2165

2100 – – 1799

– 1964 1750–1550 –

– 1520–1400 1483 1360 (Crystalline) 1343 CH2 bending (Crystalline) C–H second over tone 1244 CH2 symmetric twisting 1236 Swinging vibration of C–H in CH2 group 1350 (Amorphous) 1282 Asymmetric CH2 twisting C–C stretching 1147 Symmetric and asymmetric C–O–C stretching 1100 (Amorphous) C–O stretching with some CH2 asymmetric 947 rocking 845 CH2 rocking in PVP and with some C–O stretching in PEO C–N bending –

-

C

a

3500

+

d

b

4000

N

1557, 655 – 1460 1433, 1283 – 1244 1241–1279 1375 1293 1166, 945 – – 845 465

C

C

C

C e

O

+

Na--Br

N ----- Br O ----- Na+ C C C

C

Scheme 4. Interaction of salt with PVP carbonyl group.

is primarily due to the CH2 rocking motion with a little C–O stretching motion mixed in, while the band at 947 cm  1 originates primarily due to the C–O stretching motion with some contribution from CH2 rocking motion. The presence of these bands, assigned to CH2 rocking vibrations, indicates gauche conformations of –CH2–CH2– groups in the crystalline portion of PEO which has all –O–C–C–O– torsion angles in gauche conformations. The absence of the characteristic IR band near 1320 cm  1, assigned to CH2 stretching vibration of ethylene groups, indicates the trans conformations. These conformations are responsible for helical conformation of PEO. In particular, NaX complexes of PEO based polymer electrolytes are believed to be in the crystalline state and have a helical configuration for the polymer [54]. The effect of dopant salt on the modes of vibrations was observed in terms of decrease in the intensity, broadening of the bands with salt content and shifting of the bands to lower wave numbers which result from the formation of cross-links (Scheme 3) between the cations and ether oxygen atoms in PEO and coordination between the Na þ cation and the oxygen atom of carbonyl groups of PVP (Scheme 4). In these interactions some of the PEO chains possibly wind around the Na þ cations and C–O single bond strength decreases in PEO. In PVP, the spectral perturbation of carbonyl band becomes extensive due to increase in the basicity of CQO group with the increased salt concentration, implying that the carbonyl group is able to act as a strong electron donor to interact with Na þ cation [55]. These interactions lead to the interruption of crystallization due to which the fraction of amorphous content increases. Table 2 shows the changes associated with the position of the bands with the variation of salt concentration. The intensity of the two bands at 947 and 845 cm  1, related to helical structure of PEO, decreases and their position is shifted to lower wavenumbers with increased salt concentration. It leads to the conclusion that the PEO helical conformation is distorted, or at least, stretched with the increase in salt concentration [54]. In addition to the above variations the small band at 1351 cm  1 which is attributed to amorphous content of PEO showed higher intensity when 15 wt% NaBr was added. This indicates the transformation of crystalline regions of PEO into amorphous regions due to the addition of NaBr [54]. The characteristic peaks of NaBr disappear completely in the spectra of complexed polymer blends. This may be due to the disruption of the initial order of the polymer blend by the salt [56]. 3.3. SEM analysis

The spectral region 1000–800 cm  1 consists of a mixture of CH2 rocking and C–O stretching vibrational modes. Spectral changes in this region reflect changes occurring in the local structure of the polymer backbone. The mode responsible for the band at 845 cm  1

Scanning electron microscopy is often used to assess the compatibility between various phases through the detection of phase separations and interfaces [57–60]. SEM micrographs of pure and complexed polymer blend electrolytes reveal different

K.K. Kumar et al. / Journal of Membrane Science 454 (2014) 200–211

205

Table 2 Shifting of IR bands of PEO/PVP/NaBr polymer blend electrolytes. Vibrational mode

C–H stretching/C–C and C–O–C stretching C–H stretching/CQO stretching combination C–O stretching Symmetric and asymmetric stretching of CQO C–O stretching with some CH2 asymmetric rocking CH2 rocking in PVP and with some C–O stretching in PEO C–N stretching

Spectral position (wave number, cm  1) Pure (PEO/PVP)

(PEO/PVP/NaBr) (67.5:27.5:5)

(PEO/PVP/NaBr) (65:25:10)

(PEO/PVP/NaBr) (62.5:22.5:15)

2510 2240 2100 1650 947 845 655

2497 2230 2094 1641 936 839 643

2486 2221 2087 1633 928 831 636

2478 2214 2076 1623 912 824 627

Fig. 4. Scanning electron microscope images of PEO/PVP blend films containing NaBr ratios (a) 0 wt%; (b) 5 wt%; (c) 10 wt%; (d) 15 wt%.

surface morphologies (Fig. 4). It can be seen from Fig. 4(a), that pure PEO/PVP shows several micro-cracks. After the addition of the salt to the pure blend, the surface texture becomes somewhat smoothened (Fig. 4(b–d)); a smooth appearance is generally associated with lowering of PEO/PVP crystallinity. This arises from random distribution and dissociation of salt which may introduce topological disorder in the electrolyte, which produces more amorphous phase in the system and makes the electrolyte more flexible [61–64]. X-ray diffraction and FTIR studies of the blend samples support these observations. As there is uniform distribution of one of the polymers in the other, PEO and PVP form a compatible blend. The formation of micro-voids and craters on the surface of films is due to rapid evaporation of the solvent from the polymer blend. 3.4. Thermal analysis With a view to understand the thermal history and stability of present polymer blend electrolyte samples, DSC measurements and TGA were carried out. The DSC thermograms of pure PEO/PVP and PEO/PVP/NaBr complexes of different compositions in the temperature range  70–300 1C (Fig. 5) show a strong endothermic effect superimposed on the heat flow shift in the region

between 43 1C and 75 1C which could be attributed to the melting of PEO crystallites [65,66]. Using the DSC thermograms, glass transition temperatures (Tg), percentage of relative crystallinity (%χC) and melting temperatures (Tm) of polymer blend electrolytes were measured and their variation with salt concentration is given in Table 3. DSC thermograms showed a glass transition temperature in the region 45–60 1C at around 54 1C which is often indicated as DSC Tg (onset)-value [67]. Although the literature confirms that the glass transition temperature of pure PEO lies around  57.6 1C [65] and that of PVP around 69 1C, present PEO/PVP blend samples showed single Tg that lies between those of the individual components, indicating that the present polymer blend is miscible. Furthermore, the glass transition temperature was found to shift towards lower temperatures, with the addition of salt of different concentrations. This effect is a result of reduction in cohesive forces of attraction among polymer chains due to penetration of the salt ions into the polymer matrix. This establishes polar attractive forces between the ions and the chain segments, thereby increasing the segmental mobility. Consequently the amorphous phase becomes more flexible. The percentage of relative crystallinity of complexed polymer blend electrolytes was calculated from their DSC curves based on

206

K.K. Kumar et al. / Journal of Membrane Science 454 (2014) 200–211

100

Heat flow (w/g)

Blend with 15 wt.% NaBr

b

90

a c

d

Blend with 10 wt.% NaBr 80

Blend with 5 wt.% NaBr

Weight % (%)

70

Onset Tg

Pure blend

60 50 40

Endo

30 20 10

Tm

0

30

60

90

120

Temperature (°C) Fig. 5. DSC thermograms of pure PEO/PVP and PEO/PVP/NaBr polymer blend films.

Table 3 Glass transition temperature, melting enthalpy, relative percentage of crystallinity and melting temperature of pure and complexed polymer blend electrolytes. S. no. Polymer electrolyte

Composition

1. 2. 3. 4.

(70:30) (67.5:27.5:5) (65:25:10) (62.5:22.5:15)

Pure (PEO/PVP) PEO/PVP/NaBr PEO/PVP/NaBr PEO/PVP/NaBr

Tg (onset) (1C)

ΔΗm

58.49 51.50 50.68 49.07

223 101.9 91.5 76.9

(J/g)

%χ c

Tm (1C)

100 45.7 41.0 34.5

65.25 62.29 61.77 61.31

the following equation, assuming that pure blend is 100% crystalline [64,68,69]. % Relative crystallinity ¼ ½ΔH m =ΔH om   100

100

200

300

400

500

600

700

800

Temperature (°C) Fig. 6. Weight loss curves of PEO/PVP blend films containing NaBr ratios (a) 0 wt%; (b) 5 wt%; (c) 10 wt%; (d) 15 wt%.

In semicrystalline polymers, Tm is the melting temperature of crystallites and Tg is the glass transition temperature of noncrystalline portions surrounding crystallites. In the present investigations, the value of Tg was observed to be more sensitive than that of Tm to the salt concentration. This is due to the fact that the concentration of NaBr salt in the crystalline and amorphous phases will be different due to the change in the relative amounts of the two phases present. However both Tg and Tm showed a decreasing trend as both are controlled principally by main-chain stiffness. The smaller crystallites, due to the presence of salt, are less stiff to give a low Tm and consequently the amorphous portions will be in or adjacent to less stiff environment with decreasing crystallite size, thus resulting in a low Tg.

ð3Þ

where ΔH om is the crystalline melting enthalpy i.e. the energy in the form of heat absorbed per unit weight of the polymeric sample obtained from crystalline endothermic melting curve of the pure blend (found to be 223 J/g in the present work) and ΔΗm is the melting enthalpy estimated from the crystalline melting peaks of complexed PEO/PVP blend films. The intensity of melting endotherm decreased and shifted slightly to lower temperatures with the increase in salt concentration in the blend samples. This is apparent with the decrease in the enthalpy of melting (ΔΗm), indicating a reduction in the degree of crystallinity of the polymer electrolyte films with increasing salt concentration. This causes an increase in the amorphous phase of all complexed blend electrolytes. All the polymer blend electrolytes exhibited a relatively sharp crystalline melting endotherm (dip) at around 65 1C. The melting temperature (Tm) is taken at the apex of the melting endotherm as shown in the figure, to be the point where the polymer gets completely melted. The melting temperature (Tm) of pure PEO/PVP shifted to lower temperatures when salt was added to the polymer system. The lowering of Tm upon salt addition is quite common [46] and has been related to the decrease in spherulite sizes [70] and their surface free energy. When salt concentration further increases Tm is lowered further, due to the suppression of crystallites, thereby increasing the amorphous content in the polymer matrix. As a result of more flexible amorphous environment getting trapped in or adjacent to the crystalline matrix, the suppressed crystalline portion of the PEO/PVP blend complexes melts probably at lower temperatures.

3.5. Thermogravimetric analysis Thermal stability of a polymer electrolyte is vital for a safe and endurable electrochemical cell. During the cell reactions, heat is known to get generated in the cell which can melt or degrade the polymer electrolyte within the cell and cause internal short circuits. Hence thermogravimetric analysis is very essential to study the thermal stability of polymeric systems under application conditions. Thermal stability is represented by the weight loss of the sample after heating over the temperature range 40–850 1C. Typical thermogravimetric analysis (TGA) traces of pure and NaBr complexed PEO/PVP polymer blend electrolytes are depicted in Fig. 6. The initial 1–4% weight loss between 55 1C and 85 1C is mainly due to the evaporation of moisture absorbed by the samples during the process of sample loading [71]. The weight loss curves showed degradation of samples above 350 1C in multisteps. This multistep trend proves that the present samples are blend polymers [72]. Furthermore, the thermal degradation of polymer samples involves an additional step when salt is added [61]. The first major degradation step and minor second step (poorly resolved) in the temperature range 350–650 1C are due to the degradation of polymer blend host and the other step in the range 650–850 1C may be due to the degradation of salt which is apparent in complexed polymer blend samples [61]. A polymer is thermally stable until the decomposition process starts. Two types of thermal decomposition processes are usually recognized in polymers, chain de-polymerization and random decomposition. Chain de-polymerization is the release of monomer units from a

K.K. Kumar et al. / Journal of Membrane Science 454 (2014) 200–211

chain end or at a weak link and is essentially the reverse process of polymerization. It is often called de-propagation or unzipping. Random degradation occurs by chain rupture at random points along the chain, giving a disperse mixture of fragments. Both processes cause sample mass losses at certain high temperatures. From Fig. 6 it was observed that the position of weight loss multistep curve shifts towards lower temperatures with increase in salt concentration, which means that the sample with higher salt concentration degrades at lower temperatures. This is due to the high flexibility acquired by the polymer samples with high salt concentration. Polymers with highly flexible chains need less energy and hence show low thermal resistance and decompose at lower temperatures. Thus the presence of salt contributes to a decrease in the thermal stability of the polymer system [73]. However, the maximum shift in the position of weight loss curve with the addition of salt is only about 75 1C on temperature axis. This indicates that the inclusion of salt does not greatly affect the thermal stability. The plateau region between 150 1C and 350 1C indicates that the samples are stable in this range. Thus it can be assumed that the present polymer blend electrolytes exhibited excellent thermal stability. 3.6. Transport number measurements

current (μ amps)

To evaluate the nature of species responsible for conductivity in the present electrolyte system, the transport numbers (which give a quantitative assessment of the extent of the ionic and electronic contributions to the total conductivity) were measured by applying a constant d.c. potential of 1.5 V across the cell in the configuration Na|(PEO/PVP/NaBr)|C at room temperature. Fig. 7 shows the polarization current versus time plots. As seen at the very beginning of polarization, the current (IT) rises up only to decay immediately and asymptotically approaches steady state after a long time of polarization before stabilizing at a much lower level. The mechanism by which current initially flows across the cell at the blocking electrode under the influence of an applied voltage can be visualized as follows: initially, after application of d. c. potential, the total current across the cell, due to the migration of ions and electrons (Ii or IT ¼Iion þIele) is proportional to the 2.5 2.0 1.5 1.0 0.5 0.0 2.0 1.5 1.0 0.5 0.0 2.0 1.5

Blend with 15 wt.% NaBr

Blend with 10 wt.% NaBr

Blend with 5 wt.% NaBr

1.0 0.5 0.0 0

100

200

300

400

500

600

Time (mins) Fig. 7. Polarization current versus time plots of PEO/PVP blend films containing NaBr ratios (a) 5 wt%; (b) 10 wt%; (c) 15 wt%.

207

applied field. At the blocking electrode, there is neither a source nor a sink for mobile ions. The migration of the ions under the influence of the electric field therefore leads to an enrichment of the mobile species in the region of the electrolyte adjacent to one electrode and depletion near the other electrode. The ionic motion is then opposed by a chemical potential gradient and when, after a short time, this has increased sufficiently to counterbalance the electric field, the migration stops. The cell is then said to be concentration polarized. Hence the current starts decreasing with time as the drift of ions is balanced by diffusion of ions due to their concentration gradient induced by the electrode which blocked the ions, but still active towards electrons and hence the cell gets polarized [74]. Thus it is apparent that during d.c. polarization, the interfacial resistance increases due to the formation of passivation layer of ions at the interface of blocking electrode [74]. As a result the ionic current is blocked, the polarization is exclusively carried by electrons and hence the final current is only due to electronic current (If ¼Iele). The ionic transport numbers of present electrolyte systems are in the range of 0.968–0.984 and are given in Table 4. This suggests that the charge transport in these polymer blend electrolyte systems is predominantly ionic with a negligible contribution (0.032–0.016) from electrons. The ion transport number increases with the increase of salt concentration. This is due to the enhancement of ionic concentration (both cationic and anionic) which results in high initial current. The ion transport number reached a high value (0.984) for 15 wt% of NaBr complexed electrolyte system, which may be sufficient to meet the requirements of solid state electrochemical cells [75–78]. The measurement of ionic transport numbers by d.c. polarization method does not give information about the individual contribution of Na þ and Br  ions to the total conductivity. Hence in the present work the contribution of cations/anions to the total ionic conductivity was measured using the combination of complex impedance technique and d.c. polarization measurement. Fig. 8(a) shows the typical impedance plot of 10 wt% NaBr complexed polymer blend electrolyte film at room temperature. Two semicircular arcs were observed in the impedance plots when cation-reversible electrodes were employed. The second arc is ascribed to a parallel total interfacial resistance (charge transfer resistance Re þ passive film resistance Rp) at the reversible electrode–electrolyte interface. At room temperature, the total interfacial resistance was found to be fairly constant over a period of time, and hence the contribution from Rp was neglected while evaluating the transport number. The intercept of the high frequency dispersion curve on the real axis gives the bulk resistance Rb of the material and the second semicircle at low frequency region corresponds to the electrode–electrolyte charge transfer resistance Re. Using the values of Rb and Re cationic transport number t Na þ was calculated using the following formula: t Na þ ¼

Rb ðV =I S Þ  Re

ð4Þ

where Is is the steady state current under a d.c. bias of V volts. In principle, the steady state current Is for a cell with ion-reversible

Table 4 Transport numbers, conductivity and activation energies of PEO/PVP/NaBr polymer blend electrolytes. [Uncertainty in conductivity values ¼ 7 0.5.] S. no.

1. 2. 3.

Polymer electrolyte

PEO/PVP/NaBr PEO/PVP/NaBr PEO/PVP/NaBr

Composition

(67.5:27.5:5) (65:25:10) (62.5:22.5:15)

t Numbers tion

tele

t Na þ

t Br 

0.968 0.976 0.984

0.032 0.024 0.016

0.475 0.518 0.579

0.493 0.458 0.405

Conductivity s (S/cm)

Activation energy (eV)

2.34  10  7 6.30  10  7 1.90  10  6

0.434 7 0.01 0.4077 0.01 0.3617 0.02

208

K.K. Kumar et al. / Journal of Membrane Science 454 (2014) 200–211

2000

Blend with 10wt% NaBr

1800 1600 1400

Z" KΩ

1200

Rb 198kΩ

1000

Rb+Re 1446kΩ

800 600 400 200 0 0

200

400

600

800

1000

1200

1400

1600

1800

2000

Z' KΩ 1.2

Blend with 10wt% NaBr

current (μ amps)

1.0

0.8

0.6

0.5

1.0

1.5

3.7. Electrical conductivity studies

Voltage (V) Fig. 8. (a) a.c. complex-impedance plot of Na/PEO/PVP:NaBr(65:25:10)/Na cell. (b) Residual steady state current IS as a function of applied voltage across Na/PEO: PVP:NaBr (65:25:10)/Na cell.

electrodes (sodium metal in the present case) is I s ¼ I s Na þ þ I e

ð5Þ

where I s Na þ is the current due to sodium ions where as Ie is the electronic current. From the total ionic transport number measurements, it was found that the electronic contribution to the total current is small. Hence the total steady state current is attributed to the sodium ion conduction, i:e:

I s  I s Na þ

ð6Þ

As Watanabe et al. [44] have pointed out that Is must be taken from the region in which the Is versus V plot is linear, for calculating the transport numbers, in the present investigations too the same method is followed (Fig. 8(b)). The values of t Na þ evaluated for the complexed polymer blend system PEO/PVP/NaBr (65:25:10) are presented in Table 4. Using the total ionic transport number tion and cationic transport number tcation anionic transport number tanion, was calculated as t anion ¼ ðt ion Þtotal  t cation

Table 4 shows the variation of tcation and tanion with the concentration of salt. It is apparent that the cationic transport number increases gradually over anionic transport number with increase of salt concentration. There seems to be a correlation between percentage crystallinity and t Na þ . One can find a trend in which higher crystallinity corresponds to lower cation transport number [79]. From FTIR studies it was observed that both the cations and the anions are complexed by the polymer blend. The cation is coordinated by the electron pairs of the polymer and resides within the polymer helix; this electronation of the cation is the reason for its solvation in the polymer. The anions similarly interact with the hydrogens or nitrogens of the polymer backbone to form hydrogen (proton) bond and may be considered as protonated. Such bonds are normally very weak so that their formation and breakage may take place without much involvement of energy. They are therefore, relatively free to move through the electrolyte [80]. They should be screened from the cations to some extent and become more potent nucleophiles as a result of cation complexation. Therefore, the current should be carried mainly by the anions while crystalline helical channels hardly allow cations to move at low salt concentrations. But when salt concentration increases, crystalline conformations get distorted into amorphous phases. Anions from the salt can play the role of a plasticizer. Since amorphous regions link the crystalline regions, the cation must be able to move out of the helix and pass through amorphous regions to account for the long range motion required for d.c. conductivity. This transport across the amorphous polymer may be achieved if the cations move through a series of bridging sites. Segmental polymer motions and cation motion may be closely coupled; as the amorphous polymer chain segments move, ion transport barriers must be changing on some time scale and motions of the chain segments may promote ion transport by providing a kick of some sort [81]. This is a reasonable microscopic picture of cation migration. Hence, as the amorphous portion progressively developed in the case of 10 wt% and 15 wt% of NaBr doped samples, the transport numbers of cations t Na þ were found to increase over anions.

ð7Þ

Typical impedance plots of PEO/PVP polymer blend electrolyte with 15 wt% NaBr salt at different temperatures, in the frequency range of 1 Hz to 1 MHz, below 75 1C, are shown in Fig. 9. The plots comprise a depressed semicircular arc in the high frequency region and a tilted straight line in the low frequency region. The presence of the depressed semicircle reveals the non Debye nature of the sample [82], as the potential well for each site, through which the ion transport takes place, is not equal. It is widely accepted that the semicircle is due to the bulk resistance of polymer electrolytes, whereas the inclined line is ascribed to the charge transfer resistance and capacitance of electric double layer formed at the electrode/electrolyte interface [83], which offers increasing impedance against ion transfer with decrease in frequency. The inclination of the line at an angle of less than 901 to the real axis is due to the roughness of the electrode/electrolyte interface. At higher temperatures, the lines lean slightly towards the imaginary axis, which is an indication of establishment of better contact with the electrode. A new type of circuit element, the ‘Constant Phase angle Element’ (CPE) is needed to account for semi-circle flattening and low frequency straight line tilting. It may be thought of conceptually as a leaky capacitor, the physical origin of which, for polymer electrolytes, is perhaps related to the presence of crystalline non-conducting regions interwoven with conducting amorphous material within the spherulites. Thus no combination of resistors and capacitors (or inductors) will produce either the

K.K. Kumar et al. / Journal of Membrane Science 454 (2014) 200–211

209

1.0x10

(PEO/PVP/NaBr)(62.5:22.5:15)

2x104

303K 313K 323K 333K 343K

8.0x10-6

303K 313K 323K 333K 343K

σ (S cm-1)

Z" Ohms

3x104

1x104

(PEO/PVP/NaBr)

6.0x10-6

4.0x10-6

2.0x10-6 0.0 0.0

2x104

1x104

3x104

0.0

Z' Ohms

0

Fig. 9. Cole–Cole plots of PEO/PVP/NaBr (62.5:22.5:15) polymer blend electrolyte at different temperatures.

depression or the tilting, but both of these are natural consequences of using constant phase elements, which are hybrids between a resistor and a capacitor with an impedance of the form Z cpe ¼ kð cos ½απ =2  j sin ½απ =2Þ=ωα ðthe cos term contributes to Z 0 and the sin term to Z″Þ:

ð8Þ

where 0 r α r1 and k is a constant. When α ¼ 0, Z is frequency independent and k is just the resistance and when α ¼ 1, Zcpe ¼k/ jω, the constant k represents the inverse of capacitance. For a simple circuit consisting of a resistor and a constant phase element in series, the impedance plot is represented by a line, inclined at an angle απ/2 to the Z' axis and contacting this axis when Z'¼ R. For a resistor and a CPE in parallel, a depressed semicircle is produced. This cuts the Z' axis at the origin and at Z' ¼R, it lies vertically above the far end of the diameter of the semicircle. Thus the intersection of the flattened semicircle and the tilted line on Z' axis is the bulk resistance, Rb, of the electrolyte. This is related to the ionic conductivity by s ¼ l/RbA where l is the thickness of the polymer blend electrolyte film and A is its area. It is evident from the Cole–Cole plots at different temperatures that the intercept of the semicircle (bulk resistance) on the real axis tends towards lower values with increasing temperature, i.e. conductivity is found to increase with the temperature. Fig. 10 shows the variation of electrical conductivity with the concentration of salt at different temperatures. From the figure it is observed that the ionic conductivity at all temperatures increases (i.e. bulk resistance decreases) with the increase of salt doping and at room temperature it attains a maximum value of 1.90  10  6 S cm  1 for 15 wt% salt content. The conductivity values of different complexes at room temperature are summarized in Table 4. The highest ionic conductivity at 15 wt% salt doping is attributed to the highest electrolyte uptake, as conductivity is related to the number of Na þ charge carriers, ni, and their mobility, mi, according to s ¼ Σiniqimi [84] (where qi is the charge of each charge carrier). The coordination interactions of the ether oxygen atoms of PEO or/and carbonyl oxygen atoms of PVP, with Na þ cations of NaBr salt resulting in an increase in number of dissociated mobile charge carriers and a reduction in the crystallinity of PEO/PVP mixture. This is responsible for the increase of ionic conductivity. These interactions have been monitored by FTIR analysis of the samples and a reduction in the crystallinity is observed in XRD patterns and DSC thermograms.

5

10

15

Salt wt.% Fig. 10. Variation of conductivity with salt concentration of (PEO/PVP/NaBr) polymer blend electrolytes at different temperatures

The temperature dependent conductivity of pure and complexed polymer blend electrolytes in the temperature range 303–343 K is shown in Fig. 11. It is found that the curves follow linear trend governed by the Arrhenius equation s ¼ so exp(  Ea/ kT), where s, so, Ea, k and T are the ionic conductivity, the preexponential factor, the activation energy, the Boltzmann constant and the absolute temperature respectively. The increase in the temperature leads to increase in ionic conductivity due to increase in the polymer chain flexibility producing more free volume, which leads to enhanced polymer segmental mobility. As a result, the polymer segments with associated ions constantly rearrange themselves. This re-arrangement changes the local position of the carrier ions. The solvation of the other segments dominates sometime after the change in the local position of the carrier ions, and the segmental motion again causes the ionic migration. The repetition of the association of the carrier ions to the polymer segments, the segmental motion with associated ions, and the dissociation from the polymer segments seems to cause the ionic transport in polymer electrolyte. Thus ionic migration does not occur by itself but with the segmental motion with associated carrier ions. The moving unit of polymer backbone with associated carrier ions is correlated with that involved in the relaxation process [85]. The activation energy calculated from Arrhenius plots was found to decrease with NaBr concentration. This suggests that addition of salt in the polymer blend matrix facilitates the migration of ions. The lowering of activation energy for Na þ ion transport upon addition of salt seems to be consistent with the increased amorphous nature of the polymer blend electrolytes.

3.8. Discharge characteristics of electrochemical cells: The discharge characteristics (variation of cell parameters as a function of time) of electrochemical cells with the configuration Na|electrolyte|(I2 þC þelectrolyte) for various compositions of the electrolyte at ambient temperatures for constant load 100 kΩ are shown in Fig. 11. For the above configuration the half cell reactions can be written as At the anode: Na!Na þ þ e 

ð9Þ

210

K.K. Kumar et al. / Journal of Membrane Science 454 (2014) 200–211

and at the cathode: 1 Br2 þ e  !Br  2

ð10Þ

The overall reaction can be written as 1 Na þ Br2 !NaBr 2

ð11Þ

-4

Arrhenius linear fit

log σ (S Cm-1)

-5

d

-6

c b

-7

-8

PEO/PVP/NaBr

2.8

2.9

4. Conclusions

a

3.0

3.1

3.2

3.3

1000/T (K-1) Fig. 11. Arrhenius plots of the conductivity of PEO/PVP blend films containing NaBr ratios (a) 0 wt%; (b) 5 wt%; (c) 10 wt%; (d) 15 wt%.

3 2 1

Voltage (v)

0 3 2 1 0 3 2 1 0 0

50

100

150

200

250

From the Fig. 12, it is found that during discharge, the cell voltage decreases initially and then remains constant for a particular duration (time of stable performance of the cell) after which voltage declines. The initial sharp decrease in voltage may be due to the activation polarization [86] and/or the formation of thin layer of sodium at electrode–electrolyte interface [87]. The cell parameters such as open-circuit voltage (OCV), short circuit current (SCC), plateau region (the region in which cell voltage remains constant), energy density etc for all the cells are listed in Table 5. From the table it is found that (PEO/PVP/NaBr) (62.5:22.5:15) cell showed better performance and stability in terms of all the parameters. This may be due to the higher ionic conductivity and greater amorphicity of the electrolyte system. The cell parameters of the present electrolyte systems are comparable with those of the earlier work reported on different polymer electrolyte systems [78,88–90]. Thus, based on these studies, solid state electrochemical cells with 15 wt% NaBr doped polymer blend film as electrolyte system may be promising in practical electrochemical cell applications. This clearly indicates the applicability of the present electrolyte systems as potential candidates for solid state electrochemical cells.

300

Time in hours Fig. 12. Discharge characteristics of PEO/PVP blend films containing NaBr ratios (a) 5 wt%; (b) 10 wt%; (c) 15 wt%.

Polymer blend electrolytes based on polyethylene oxide and polyvinyl pyrrollidone complexed with NaBr were prepared in the form of thin films using solution cast method. Sharp Bragg peaks along with broad peaks observed in XRD patterns revealed that the blend material is semicrystalline. The broadening and reduction in the intensity of the peaks with the addition of the salt to the blend polymer host confirmed the dissolution of salt in the polymer host and the increase in its amorphous nature. FTIR studies suggested the changes in the environment of the functional groups upon the addition of salt which confirms the coordination of cations and the interaction of anions with the polar groups of polymer blend host during the formation of complex. Furthermore co-existence of vibrational bands pertaining to the polar groups of PEO and PVP confirms the miscibility of the blend. Electron micrographs showed uniform distribution of one polymer in the other which confirms the compatibility of the blend components. Attainment of smooth surface morphology upon the addition of salt also suggests the enhancement of degree of amorphicity. This has been further confirmed by measuring the relative percentage of crystallinity using the DSC thermograms of the present polymer blend systems. A decrease in the percentage of relative crystallinity was observed as a function of NaBr salt concentration. The glass transition temperature Tg and melting temperature Tm also showed similar trend as the structure becomes more flexible at high salt concentrations. Single glass transition temperatures measured as onset value confirm that the present polymer blend is miscible. The multistep weight loss

Table 5 Cell parameters of Na/(PEO/PVP/NaBr)/(I2 þ Cþ electrolyte) polymer blend electrolytes. Cell parameters

Units

(PEO/PVP/NaBr) (67.5:27.5:5)

(PEO/PVP/NaBr) (65:25:10)

(PEO/PVP/NaBr) (62.5:22.5:15)

Open circuit voltage Short circuit current Area of the cell Weight of the cell Discharge time for plateau region Current density Power density Energy density Discharge capacity

(V) (mA) (cm2) (g) (h) (mΑ cm  2) (W kg  1) (W h kg  1) (μA h  1)

2.71 238 1.168 1.10 189 203.76 0.586 110.75 1.25

2.85 276.5 1.168 1.04 215.58 236.72 0.757 163.3 1.29

2.89 316 1.168 1.15 230.48 270.54 0.801 183 1.37

K.K. Kumar et al. / Journal of Membrane Science 454 (2014) 200–211

curves of TGA indicate the presence of blend in the host polymer. The present polymer blend system showed good thermal stability up to 350 1C. Transport number studies suggest that the charge transport in these polymer blend electrolytes is predominantly ionic and the high degree of amorphicity of samples favours cationic transportation. It was observed that increase in amorphous phase due to the rise of temperature and free ion concentration is responsible for conductivity improvement in these electrolytes. The activation energy was found to decrease gradually with increase of amorphous phase in polymer matrix. (PEO/ PVP/NaBr) (62.5:22.5:15) cell showed better performance and stability in terms of the parameters like OCV, SCC and discharge time for the plateau region etc. This is attributed to the presence of flexible matrix and high ionic conductivity. References [1] M.B. Armand, J.M. Chabagno, M.J. Duclot, in: P. Vashishta, J.N. Mundy, G. K. Shenoy (Eds.), Fast Ion Transport in Solids, North-Holland, New York, 1979. [2] M.B. Armand, Polymer electrolytes, Ann. Rev. Mater. Sci. 16 (1986) 245–261. [3] B. Scrosati, in: B. Scrosati (Ed.), Applications of Electro Active Polymers, Chapman and Hall, London, 1993. [4] M.A. Ratner, D.F. Shriver, Ion transport in solvent free polymers, Chem. Rev. 88 (1988) 109–124. [5] J.R. Owen, Super Ionic Solids and Solid Electrolytes—Recent Trends, Academic Press, New York, 1989. [6] N. Angulakshmi, Sabu Thomos, K.S. Nahm, A Manuel Stephan, R. Nimma Elizabeth, Ionics 17 (2011) 407–414. [7] A. Manuel Stephan, Eur. Polym. J. 42 (2006) 21–42. [8] A. Manuel Stephan, K.S. Nahm, Polymer 47 (2006) 5952–5964. [9] D.E. Fenton, J.M. Parker, P.V. Wright, Polymer 14 (1973) 589–590. [10] M.B. Armand, J.M. Chabagno, M. Duclot, in: St. Andrews (Ed.), Proceedings of 2nd International Meeting Solid Electrolytes, Scotland, 1978. [11] F.M. Gray, in: F.M. Gray (Ed.), Solid Polymer Electrolytes, VCH, New York, 1991, p. 4. [12] S.M. Jahurak, M.L. Kaplan, E.A. Reitman, D.W. Murray, R.J. Cava, Macromolecules 21 (3) (1988) 654–660. [13] A. Vallee, S. Besner, J. Prud’homme, Electrochim. Acta. 37 (1992) 1579–1583. [14] P.W.M. Jacobs, L.W. Lorimer, A. Russer, M. Wasiucionek, J. Power Sources 26 (1989) 503–510. [15] M. Armand, W. Gorecki, R. Andreani, in: B. Scrosati (Ed.), Proceedings of the Second International Meeting on Polymer Electrolytes, Elsevier, New York, 1990, pp. 91. [16] B.L. Papke, M.A. Ratner, D.F. Shriver, J. Phys. Chem. Solids 42 (1981) 493–500. [17] P. Lightfoot, M.A. Mehta, P.G. Bruce, Science 262 (1993) 883–885. [18] R.C. Agarwal, G.P. Pandey, J. Phys. D: Appl. Phys. 41 (2008) 18223001 41 (2008) 18. [19] Felix B. Dias, Lambertus Plomp, Jakobert B.J. Veldhuis, J. Power Sources 88 (2000) 169–191. [20] J.A. Siddiqui, P.V. Wright, Polym. Commun. 28 (1987) 89. [21] B. Mussarat, K. Conheeney, J.A. Siddiqui, P.V. Wright, Br. Polym. J. 20 (1988) 293. [22] X.Q. Yang, H.S. Lee, C. Xiang, J. McBreen, in: Proceedings of the 11th Annual Battery Conference on Applications and Advances, California, January 9–12, 1996, pp. 146. [23] A. Killis, J.F. LeNest, H. Cheradame, A. Gandini, Macromol. Chem. 183 (1982) 2835–2845. [24] D.W. Kim, J.K. Park, M.S. Gong, H.Y. Song, Polym. Eng. Sci. 34 (17) (1994) 1305–1313. [25] F.M. Gray, in: F.M. Gray (Ed.), Solid Polymer electrolytes, VCH, New York, 1991, p. 95. [26] L.M. Ding, Polymer 38 (16) (1997) 4267–4273. [27] L.M. Ding, J. Shi, C.Z. Yang, S.J. Dong, Polym. J. 29 (5) (1997) 410–416. [28] L.A. Utracki, Polymer Alloys and Blends, Carl Hanser, Munich, FRG, 1990. [29] O. Inganas, Br. Polym. J. 20 (1988) 223–233. [30] D.W. Kim, J.K. Park, H.W. Rhee, Solid State Ion. 83 (1996) 49–56. [31] G Di Marco, M. Lanza, M. Pieruccini, Solid State Ion. 89 (1996) 117–125. [32] W. Wieczorek, J.R. Stevens, J. Phys. Chem. 101 (9) (1997) 1529–1534. [33] J.M. Tarascon, M. Armand, Nature 414 (2001) 359–367. [34] G.P. Pandey, R.C. Agarwal, S.A. Hashmi, J. Power Sources 190 (2009) 563–572. [35] G.G. Kumar, S. Sampath, Solid State Ion. 160 (2003) 289–300. [36] R.M. Dell, Solid State Ion. 134 (2000) 139–158. [37] J. Barker, M.Y. Saidi, J.L. Swoyer, Electrochem. Solid State Lett. 6 (3) (2003) A53–A55. [38] J. Vondrák, J. Reiter, J. Velicka, M. Sedlaříkova, Solid State Ion. 170 (2004) 79–82. [39] D. Fautex, M.D. Lupien, C.D. Robitaille, J. Electrochem. Soc. 134 (1987) 2761–2767. [40] S.G. Greenbaum, Y.S. Pak, M.C. Wintersgill, J.J. Fontanella, J.W. Schultz, C. G. Andeen, J. Electrochem. Soc. 135 (1988) 235–238.

211

[41] S.G. Greenbaum, K.J. Adamic, Y.S. Pak, M.C. Wintersgill, J.J. Fontanella, Solid State Ion. 28–30 (1988) 1042–1046. [42] P.G. Bruce, C.A. Vincent, J. Chem. Soc. Faraday Trans. 89 (17) (1993) 3187–3203. [43] J.B. Wagner, C. Wagner, J. Chem. Phys. 26 (1957) 1597–1601. [44] M. Watanabe, S. Nagano, K. Sanui, N. Ogata, Solid State Ion. 28–30 (1988) 911–917. [45] R.C. Agrawal, S.A. Hashmi, Ionics 13 (2007) 295–298. [46] S. Arup Dey, S. Karan, S.K. De, Solid State Commun. 149 (2009) 1282–1287. [47] Jie Xu, Qi Xiong, Guijie Liang, Xiaolin Shen, Hantao Zou, Weilin Xu, IonPolymer interactions in SmCl3(H2O)6 doped poly(ethylene oxide) electrolytes, J. Macromol. Sci. Part B: Phys. 48 (4) (2009) 856–866, http://dx.doi.org/10. 1080/00222340902958778. [48] Y. Takahashi, H. Tadokoro, Macromolecules 6 (1973) 672–675. [49] M. Morita, F. Araki, N. Yoshimoto, M. Ishikawa, H. Tsutsumi, Solid State Ion. 136–137 (2000) 1167–1173. [50] X.G. Li, I. Kresse, J. Springer, J. Nissen, Y.L. Yang, Polymer 42 (2001) 6859–6869. [51] Changhua Liu, Chaobo Xiao, Hui Liang, J. Appl. Polym. Sci. 95 (2005) 1405–1411. [52] S. Rajendran, O. Mahendran, Ionics 7 (2001) 463–468. [53] K. Kiran Kumar, M. Ravi, Y. Pavani, S. Bhavani, A.K. Sharma, V.V.R. Narasimha Rao, J. Non-Cryst. Solids 358 (2012) 3205–3211. [54] Sireerat Intarakamhang, Preparation, characterization and molecular modeling of poly(ethylene oxide)/poly(vinyl pyrrolidone) montmorillonite nanocomposite solid electrolytes, Thesis adviser: Asst. prof. Visit Vao-Soongnern, 2005, pp. 54–175, ISBN: 974-533-520-7. [55] W. Wieczorek, J.R. Stevens, J. Phys. Chem. B 101 (1997) 1529–1534. [56] S. Ahmad, H.B. Bohidar, S. Ahmad, S.A. Agnihotry, Polymer 47 (2006) 3583–3590. [57] G. Hsiue, W. Kuo, Y. Huang, R. Jeng, Polymer 41 (2000) 2813–2825. [58] Z. Wen, M. Wu, T. Itoh, M. Kubo, Z. Lin, O. Yamamoto, Solid State Ion. 148 (2002) 185–191. [59] L. Fan, Z. Dang, C.W. Nan, M. Li, Electrochim. Acta 48 (2002) 205–209. [60] S. Mitra, A.R. Kulkarni, Solid State Ion. 154–155 (2002) 37–43. [61] S.A.M. Noor, A. Ahmad, I.A. Talib, M.Y.A. Rahman, Ionics 16 (2010) 161–170. [62] P.P. Chu, M.J. Reddy, J. Power Sources 115 (2003) 288–294. [63] S. Zang, Jim Y. Lee, L. Hong, J. Power Sources 134 (2004) 95–102. [64] P.P. Chu, M.J. Reddy, H.M. Kao, Solid State Ion. 156 (2003) 141–153. [65] T.M.W.J. Bandara, B.-E. Mellander, I. Albinsson, M.A.K.L. Dissanayake, H.M.J. C. Pitawala, J. Solid State Electrochem. 13 (2009) 1227–1232. [66] S.A. Saq’an, A.S. Ayesh, A. Zihlif, Opt. Mater. 24 (2004) 629–636. [67] W.M. Groenewoud, Characterization of polymers by thermal analysis, Eerste Hervendreef 32, 5232 JK 's Hertogenbosch, Elsevier, the Netherlands, 2001, pp. 13. [68] M.J. Reddy, P.P. Chu, Solid State Ion. 149 (2002) 115–123. [69] M. Jaipal Reddy, Peter. P. Chu, J. Siva Kumar, U.V. Subba Rao, J. power sources 161 (2006) 535–540. [70] J. Przyluski., K. Such, H. Wycislik, W. Wieczorek, Synth. Met. 35 (1990) 241–247. [71] A.M. Stephan, Yuria Saito, N. Muniyandi, N.G. Renganathan, S. Kalyanasundaram, R. Nimma Elizabeth, Solid State Ion. 148 (2002) 467–473. [72] Chun-Chen Yang, J. Power Sources 109 (2002) 22–31. [73] Ana F. Nogueira, M.A.S. Spinace, W.A. Gazotti, E.M. Girotto, Maco-A. De Paoli, Solid State Ion. 140 (2001) 327–335. [74] Wu Xu, Kok Siong Siow, Zhiqiang Gao, Swee Yong Lee, J. Electrochem. Soc. 146 (12) (1999) 4410–4418. [75] C.S. Ramya, T. Savitha, S. Selvasekarapandian, G. Hiran Kumar, Ionics 11 (2005) 436–441. [76] V.M. Mohan, P.B. Bhargav, V. Raja, A.K. Sharma, V.V.R. Narasimha Rao, Soft Mater. 5 (1) (2007) 33–46. [78] M. Jaipal Reddy, T. Sreekanth, M. Chandrasekhar, U.V. Subba Rao, J. Mater. Sci. 35 (2000) 2841–2845. [79] Amita Chandra, Suresh Chandra, J. Phy. D: Appl. Phys. 27 (1994) 2171–2179. [80] A.K. Sen, S. Son, Anionic transport in polymeric solid electrolytes, in: St. Andrews (Ed.), Proceedings of the First International Symposium on Polymer Electrolyte, AD-A223 013, D 5742-CH-02 Scotland, 1987, pp. 29.1–29.4. [81] B.L. Papke, M.A. Ratner, D.F. Shriver, J. Electrochem. Soc. 129 (8) (1982) 1694–1701. [82] S. Lanfredi, P.S. Saia, R. Lebullenger, Solid State Ion. 146 (2002) 329–339. [83] K.M. Kim, K.S. Ryu, S.-G. Kang, S.H. Chang, I.J. Chung, Macromol. Chem. Phys. 202 (2001) 866–872. [84] S.S. Sekhon, H.P. Singh, Solid State Ion. 169 (2002) 152–153. [85] Masayoshi Watanabe, Ionic conduction of polymer electrolytes and future applications, in: St. Andrews (Ed.), Proceedings of the First International Symposium on Polymer Electrolyte, AD-A223 013, D 5742-CH-02, Scotland, 1987, pp. 6.1–6.4. [86] H.J. De Bruin, J. Am, Ceram. Soc. 14 (1978) 20. [87] V.M. Mohan, V. Raja, A.K. Sharma, V.V.R. Narasimha Rao, Ionics 12 (2006) 219–226. [88] P. Balaji Bhargav, V. Madhu Mohan, A.K. Sharma, V.V.R. Narasimha Rao, Ionics 13 (2007) 173–178. [89] H. Yuankang, C. Zhusheng, Z. Zhiyi, in: B.V.R. Chowdari (Ed.), Materials for Solid State Ionics, World Scientific, Singapore, 1986, p. 333. [90] Ch.V. Subba Reddy, Xia Han, Quan-Yao Zhu, Li-Qiang Mai, Wen Chen, Eur. Polym. J. 42 (2006) 3114–3120.