Mixed anion effect on the ionic transport behavior, complexation and various physicochemical properties of ionic liquid based polymer gel electrolyte membranes

Abstract

Li-ion conducting polymer gel electrolyte membranes (PGEMs) containing ionic liquid (IL), 1-butyl-3-methylimidazolium tetrafluroborate BMIMBF₄, polymer poly(vinylidene fluoride-co-hexafluoropropylene) PVdF-HFP and lithium bis(trifluoromethanesulfonyl)imide LiTFSI salt (having different anion i.e. BF₄⁻ and TFSI⁻) have been synthesized and characterized by various techniques. The results show that the synthesized PGEMs have good free-standing characteristics, good thermal stability (300–400 °C) and also have a wide electrochemical window (ECW) ∼4.0–4.20 V. The conductivity increases with increasing amount of IL, and attains a value of 3.2 × 10⁻³ S cm⁻¹ at room temperature for the PGEMs containing higher loadings of IL. A high total ionic transference number (∼0.99) and cationic transference number (t_Li⁺ ∼ 0.33) for the PGEMs containing higher loadings of IL have been obtained.

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Introduction

Polymer gel electrolyte membranes (PGEMs) are promising alternatives to be used as electrolyte membranes in various technological applications like high energy density rechargeable batteries, fuel cells, solar cells and electrochromic windows.^1–6 Usually polymer electrolytes for rechargeable battery applications are formed by using different polymers like PEO, PMMA, PVA, PAN and PVdF-HFP (poly(vinylidene fluoride-co-hexafluoropropylene)) etc. mixed with different alkali salts like LiBF₄, LiClO₄, NaCl, LiPF₆, LiTFSI, NaTFSI etc. with conventional organic plasticizers such as ethylene carbonate (EC), propylene carbonate (PC), dimethylcarbonate (DMC), diethylcarbonate (DEC) etc.^7–13 Among the above mentioned polymers, PVdF-HFP based polymeric membranes have emerged as a promising host matrix for the preparation of PGEMs. The role of the polymer matrix is mainly to store the liquid electrolyte and provide good mechanical support. The properties which make PVdF-HFP a promising matrix are its good hydrophobicity, excellent mechanical strength, electrochemical stability, ability to entrap large amount of a liquid electrolyte while retaining good mechanical and thermal stability, flexibility and transparency. PVdF-HFP also has a high dielectric constant (∼8.4) that helps in dissociation of charge carriers.^14–17 PVdF-HFP based PGEMs were first developed by Bellcore for lithium ion conducting polymeric membranes.¹⁸ Although, the PGEMs prepared with conventional organic plasticizers have high ionic conductivity but there are numerous shortcomings associated with these membranes such as flammability and volatility that makes them incompatible for the high-temperature range of operation.^19–21

In order to suppress the aforementioned problems associated with the conventional organic plasticizers we have incorporated different amounts of the ionic liquid (IL) in polymer electrolyte (i.e. PVdF-HFP + 20% LiTFSI). Ionic liquids are recently discovered materials and have been subjected to wide investigation owing to their exceptional properties like negligible vapour pressure, non-flammability, high ionic conductivity, wide electrochemical window, good chemical and thermal stability etc. ILs are molten salts having melting temperature below 100 °C and some of the ILs are present in liquid state even at room temperature and are known as room temperature ionic liquids (RTILs). Unlike inorganic salts that require solvent for dissociation into cation and anion, ionic liquids do not require any solvent for dissociation and are entirely composed of dissociated organic cations and inorganic/organic anions.^22–26

In our earlier studies, we have prepared PGEMs, PVdF-HFP/LiBF₄/BMIMBF₄ and PVdF-HFP/LiTFSI/BMIMTFSI having the same anion (i.e. salt and IL are having a common anion, BF₄⁻ and TFSI⁻ respectively).^27,28 It has been found that in these systems, there is less or no chance of contact and cross-contact ion pair formation hence the enhancement in conductivity was generally observed with IL content.^9,29 But in the present study, we have investigated the various physicochemical properties of polymer gel electrolyte membranes (PGEMs) (PVdF-HFP + 20% LiTFSI) + x% BMIMBF₄, where x = 10, 20, 30, 40, 50, 60 and 70. The IL and salt used in the present study 1-butyl-3-methylimidazolium tetrafluoroborate, BMIMBF₄ and lithium bis(trifluoromethanesulfonyl)imide LiTFSI salt have different anions i.e. [BF₄⁻] and [TFSI⁻] respectively which makes our system mixed anion system. It is well documented in the literature that the mixed-anion systems (i.e., in which a dopant salt and added IL having a different anion) have a chance to form contact (i.e. ion pair formation between cations and anions belonging to the respective dopant material, BMIM⁺…BF₄⁻ and Li⁺…TFSI⁻) and cross contact ion pair formation (i.e., ion pair formation between cations and anions belonging to the different dopant material, Li⁺…BF₄⁻ & BMIM⁺…TFSI⁻) which do not take part in conduction mechanism.^9,29 To the best of our knowledge, results are not available for PVdF-HFP based mixed anion systems. Keeping in view above consideration, we tried to address the question of extent of increment/decrement in the conductivity with varying concentration of ionic liquid. This observation is expected to give insights for the practical use of the synthesized membranes. Therefore, it would be interesting to study such a system. Moreover, this system may possibly prove useful for applications in electrochemical devices especially in rechargeable batteries. In this article, we report the synthesis and structural, thermal, mechanical and electrical transport behaviour of polymeric gel membranes on (PVdF-HFP + 20% LiTFSI) + x% BMIMBF₄. We have also addressed some more issues in this study, such as whether the incorporation of the IL in polymer electrolyte i.e. (PVdF-HFP + 20% LiTFSI) would change the crystallinity, thermal stability, melting temperature (T_m), complexation of PGEMs. Various characterization techniques such as scanning electron microscopy (SEM), X-ray diffraction (XRD), differential scanning calorimetry (DSC), thermogravimetric analysis (TGA), Fourier transform infrared (FTIR) spectroscopy, complex impedance spectroscopy and cyclic voltammetry (CV) have been used to characterize the prepared polymer gel electrolyte membranes and are discussed in Results and discussion section one by one.

Experimental

Materials

The starting materials poly(vinylidene fluoride-co-hexafluoropropylene) (P(VdF-HFP), molecular weight = 400 000 g mol⁻¹), LiTFSI (purity > 99.9%) salt, and the IL BMIMBF₄ (purity > 99%) were procured from Sigma Aldrich (Germany). The IL was dried in vacuum at about 10⁻⁶ Torr for 2 days before use.

Synthesis of the polymeric gel membranes

Conventional solution casting technique has been used for the preparation of polymer gel electrolyte membranes. For this, a desired amount of polymer PVdF-HFP was dissolved in acetone under stirring at 50 °C until a clear solution was obtained. After that, appropriate amount of LiTFSI salt was added and again stirred for 3–4 hours at 50 °C. After complete dissolution of salt with the polymer PVdF-HFP, desired amount of IL was added and again stirred for 4–5 hours to obtain a homogeneous mixture. The viscous slurry so obtained was poured into polypropylene Petri dishes for casting. After complete evaporation of the solvent, free-standing rubbery films of polymeric gel membranes containing different amounts of ionic liquid having thickness ∼200–300 μm were obtained. All the membranes are mechanically stable and highly flexible in nature (see Fig. 1). Samples were handled in an argon filled glove box (MBRAUN ABstar) with H₂O and O₂ ≤ 0.5 ppm.

Fig. 1 A typical photograph of prepared polymer gel electrolyte membrane.

X'Pert PRO X-ray diffractometer (PANalytical) with CuKα radiation (λ = 1.54 Å) in the range 2θ = 10° to 80° was used to record the X-ray diffraction profiles of the polymer gel electrolyte membranes. Scanning electron microscope (model Quanta C-200) was used to examine the surface morphology of the polymeric gel electrolyte membranes.

Thermal analysis was carried out by differential scanning calorimetry using Mettler DSC 1 system in the temperature range −110 to 160 °C at a heating rate of 10 °C min⁻¹ and thermogravimetric analysis (TGA) under continuous purging of nitrogen using Mettler DSC/TGA 1 system. The FTIR spectra of the polymeric gel membranes were recorded with a Perkin-Elmer FTIR spectrometer (Model RX 1).

Bulk elastic modulus (E) of the prepared PGEMs at room temperature was investigated using pulse-echo technique. Radio frequency pulses were sent by the pulser/receiver (Olympus model 5900PR) to excite a 6 MHz piezoelectric transducer (D6HB-10) to generate longitudinal ultrasonic waves. The transducer used for transmitting as well as receiving ultrasonic waves was coupled to the disc-shaped membrane. The return echo was received by the pulser/receiver and both the echo pulse and the input pulse were displayed on a 500 MHz Agilent digital storage oscilloscope DSO5052A. The transit time of the echo pulse was recorded, and velocity of propagation of ultrasonic waves in polymeric membranes was calculated using the relation v = 2d/t. The E values of the samples were calculated using the relation E = v²ρ, (where v = velocity of the longitudinal wave and ρ = density of the samples (i.e. ρ = mass/volume)).

Ionic conductivity of the polymeric gel membranes was measured by complex impedance spectroscopy technique using NOVO Control impedance analyzer in the frequency range 1 Hz–40 MHz. The bulk resistance was determined from the complex impedance plots. The electrical conductivity (σ) can be calculated by using the following relation:

where, l = thickness of the sample, A = cross-sectional area of the sample (usually in disc shaped) and R_b = bulk resistance obtained from complex impedance plots. For temperature dependent conductivity measurements, disc shaped samples were sandwiched between two stainless steel electrodes and then whole assembly was kept in a temperature controlled oven.

The total ionic transport number (t_ion) was determined by d.c. polarization technique. For this, a voltage of 10 mV was applied across the disc shaped samples (sandwiched between two stainless steel electrodes) and the resultant current was monitored as a function of time. The combination of a.c./d.c. technique was used to calculate cationic transport number (i.e. t_Li⁺) of the PGEM containing higher amount of IL. In this technique, Li/PVdF-HFP + 20% LiTFSI + 70% BMIMBF₄/Li cell was subjected to polarization by applying a voltage ΔV = 10 mV, for 1 hour and resulting currents were calculated (i.e. initial and final current). The cell resistances were also measured before and after the polarization using a.c. impedance spectroscopy.

The cyclic voltammetric studies were carried out using electrochemical analyzer with an AUTOLAB PGSTAT 302N controlled by NOVA 1.10.4 software version (Metrohm Auto Lab) to estimate the ‘electrochemical stability window’ of the polymer electrolyte membrane.

Results and discussion

SEM study

Fig. 2[a]–[c] shows the surface morphology of the pristine PVdF-HFP, (PVdF-HFP + 20 wt% LiTFSI) + 20 wt% BMIMBF₄ and (PVdF-HFP + 20 wt% LiTFSI) + 70 wt% BMIMBF₄. Fig. 2[a] shows the crystalline grains with lamellar structure of the pristine PVdF-HFP membrane. On incorporation of 20 wt% of LiTFSI salt and 20 wt% of IL together in the PVdF-HFP polymer, the size of the grains decreased (see Fig. 2[b]). When the concentration of IL varies from 10 to 70 wt% in the polymer electrolyte the grains size starts decreasing which makes the prepared PGEMs more amorphous and flexible with no crystalline grains (see Fig. 2[c]). Therefore, from above discussions we conclude that the surface morphology is uniform and amorphous phase present throughout the membranes in the presence of IL in the system. The result suggests that the polymer PVDF-HFP has superior solvent retention ability as compared to other existing polymers like PEO, PVA, PVP etc. that makes PVdF-HFP based polymer electrolyte membranes more amorphous.

Fig. 2 SEM micrograph of the (a) pristine PVdF-HFP, (PVdF-HFP + 20 wt% LiTFSI) + x wt% BMIMBF₄ (b) x = 20 and (c) x = 60.

Mechanical stability test

The elastic modulus (E) of the prepared PGEMs was calculated (description is given in the Experimental section). The value of E of pure PVdF-HFP, PVdF-HFP + 20% LiTFSI (as calculated in our earlier study^27,28), and (PVdF-HFP + 20% LiTFSI) + x wt% of BMIMBF₄ are listed in Table 1. The decrease in E values of the PGEMs with the addition of salt and IL in PVdF-HFP matrix was due to the plasticization effect of the IL.

Table 1 The value of elastic modulus of P(VdF-HFP) + 20% LiTFSI + x wt% of IL for different values of ‘x’ obtained by ultrasonic measurement

PGEMs (PVdF-HFP + 20% LiTFSI) + x wt% BMIMBF4	Elastic modulus (dyne cm^−2)
Pure PVdF-HFP	14.56 × 10^10
x = 0	7.5 × 10^10
x = 20	7.0 × 10^10
x = 40	6.54 × 10^10
x = 70	5.89 × 10^10

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XRD study

The X-ray diffraction profile of pristine PVdF-HFP and (PVdF-HFP + 20% LiTFSI) + x wt% BMIMBF₄ (where x = 0, 20, 40, 50 and 60) are shown in Fig. 3. From Fig. 3[a], it is found that the pristine PVdF-HFP consist five crystalline peaks at 2θ = 16.75°, 18.21°, 20.04°, 26.71°, and 38.83° that belong to the crystalline phase of α-PVDF. As reported earlier, the sharp crystalline peaks related to LiTFSI salt appeared at 2θ = 14.1, 15.9, 18.6, 18.9 and 21.4.^30–32 From Fig. 3 (curve b), we have found that when salt (20 wt%) was mixed in the polymer PVdF-HFP matrix, three peaks related to the crystalline phase of PVdF-HFP at 2θ = 16.75, 18.21, 26.71 and approximately all the crystalline peaks related to LiTFSI salt vanish and only one peak of LiTFSI at 2θ = 14.11 (see Fig. 3 (curve b)) remains. From Fig. 3 (curve c), it has been found that upon incorporation of lower amount of IL (∼20 wt%) in the polymer electrolyte PVdF-HFP + 20% LiTFSI, the remaining peak of LiTFSI disappeared and reduction in the crystalline peaks of PVdF-HFP was also observed. The intensity of the above mentioned crystalline peaks of PVdF-HFP continuously reduced with the increasing content of IL in the PGEMs due to the enhanced amorphicity of the prepared membranes.^27,28

Fig. 3 XRD profile of (a) pristine PVdF-HFP (b) PVdF-HFP + 20 wt% LiTFSI, (PVdF-HFP + 20 wt% LiTFSI) + x wt% BMIMBF₄ (c) x = 20 (d) x = 40 (e) x = 50 and (f) x = 60.

Ionic conductivity study

In battery application, the polymer electrolytes act as a separator that provides a medium through which ions can transport from anode to cathode and vice versa. Hence, the ionic conductivity of the polymer electrolyte is a very important parameter for the practical use in solid state lithium ion battery applications. The room temperature ionic conductivity of the polymer electrolyte i.e. PVdF-HFP + 20 wt% LiTFSI was found to be ∼1.01 × 10⁻⁷ S cm⁻¹ (ref. 31) which is relatively low for applications. A rapid increment of almost two orders was observed upon incorporation of small (i.e. 10% of IL) amount of IL. It was found that the ionic conductivity gradually increases with increasing IL content in PGEMs and reaches ∼3.2 × 10⁻³ S cm⁻¹ for higher loading of IL at room temperature (see Fig. 4[A]). Fig. 4[B] shows the Nyquist plots (measured at room temperature) of the two prepared PGEMs containing 10% and 20% of IL (similar nature was found for all the prepared PGEMs). The suppressed semicircle (in the high-frequency region) followed by an inclined spike (in the low-frequency region) was observed because of parallel combination of the double layer capacitance C_dl and bulk resistance (R_b) that signifies the diffusion of ions. The angle of the depressed semicircle and the inclination angle of the straight line can be attributed to the occurrence of distributed macroscopic materials properties (i.e. constant phase element (CPE)) (see inset of Fig. 4[B]).

Fig. 4 [A] The composition and temperature dependent ionic conductivity of the PGEMs (PVdF-HFP + 20 wt% LiTFSI) + x wt% BMIMBF₄ (a) x = 10, (b) x = 20, (c) x = 30 (d) x = 40 (e) x = 50, (f) x = 60 and (g) x = 70. [B] Typical Nyquist plots of (PVdF-HFP + 20% LiTFSI) + x wt% BMIMBF₄ (a) x = 10 and (b) x = 20 at room temperature and its equivalent circuit (inset of the figure).

From Fig. 4[B] it can be seen that upon increasing the amount of IL (i.e. only from 10 to 20%) in the membranes the bulk resistance of the system reduced considerably and it continuously decreased with the increasing amount of IL which resulted increment in conductivity upon inclusion of IL in the membranes. In the present study, we have found monotonic increase in the ionic conductivity on increasing the IL content except for the membrane containing 60 wt% of IL (after certain temperature, see Fig. 4[A]). This result is beneficial for practical applications. Therefore, we concluded that the presence of LiTFSI salt in the membranes does not promote the formation of contact/cross-contact ion pairs because it has low lattice energy and also has a poor tendency to form ion-pairs. The LiTFSI salt is also known to have a high degree of salt dissociation because of the presence of strong electron-withdrawing SO₂CF₃ groups at both sides of the imide anion. Many researchers have also used LiTFSI salt to prepare polymer gel electrolyte membranes for battery applications and they have concluded that LiTFSI salt works as a plasticizer for polymer matrix and creates free-volume that enhances ionic mobility and hence the ionic conductivity^33–35 (as ionic mobility and ionic conductivity are directly proportional to each other). This is not only because of the choice of salt it's also because of the polymer used (i.e. PVdF-HFP in the present case). It is expected that polymer PVdF- and its copolymer PVdF-HFP does not form complexes with the cation of the salts/or ILs. These are known for their inertness, hydrophobicity and usually they are chemically stable.^15–17 Thus, the PGEMs containing polymer PVdF-HFP and LiTFSI salt show high ionic conductivity.

In accordance with the results available in literature, the ionic conductivities obtained are of the same order of magnitude. For example, Hofmann et al.³⁶ reported an ionic conductivity of 1.04 ± 0.3 mS cm⁻¹ for 0.5 mol kg⁻¹ Li-TFSA in MPPyrr-TFSA/PVdFHFP. Navarra et al. presented an ionic conductivity of 0.3 mS cm⁻¹ for 0.2 M Li-TFSA in a N-butyl-N-ethylpyrrolidinium TFSA-PVdF-HFP (30 wt%) matrix at room temperature.³⁷ Ye et al. reported an ionic conductivity of 0.3 mS cm⁻¹ for 0.5 M LiTFSI in MPPyrr-TFSA/PVdF-HFP (30 wt%) matrix at 23 °C.³⁸ Here, we found room temperature ionic conductivity ∼3.2 × 10⁻³ S cm⁻¹ for the PGEM having higher loading of IL (i.e. 70%). Fig. 4[A] shows the temperature dependent ionic conductivity and we have found that as we increase the temperature, conductivity of prepared PGEMs increases and follows an Arrhenius type thermally activated process that can be expressed as given below:

σ = σ_o exp(−E_a/kT) (2)

where, σ_o = pre-exponential factor, E_a = activation energy, k = Boltzmann constant and T = temperature (in kelvin). The σ vs. 1/T plots (see Fig. 4[A]) show rapid jump in the conductivity values at a definite temperature which is the melting temperature (T_m) of polymer at which the amorphicity of the membranes increases significantly. For most polymer electrolytes, it is well documented in literatures that the in crystalline phase, the polymer chain is hard or rigid and thus ionic mobility is almost negligible therefore crystalline phase has lower conductivity than amorphous phase. The melting temperature (T_m) of the prepared PGEMs changes with the concentration of IL as examined by the DSC results (see Fig. 6) which are discussed later in this investigation. The activation energy (E_a) was calculated from Fig. 4[A] (region 1) and found to decrease with the increasing concentration of IL in PGEMs that shows that IL provides easier ionic transport in the system hence increases the ionic conductivity.

The enhancement in the ionic conductivity of the prepared PGEMs on increasing IL concentration is due to the increase in number of mobile charge carriers and their mobility. This can be explained by the general conductivity relation

where, n_i is the number of charge carriers, q_i is the charge of mobile charge carrier and μ_i is the mobility of charge carriers. According to eqn (3), the enhancement in the ionic conductivity of the polymer electrolytes can be achieved by increasing number of charge carriers and/or the mobility.

The total ionic transport number (t_ion) of the prepared PGEM containing higher amount of IL was calculated (see Fig. 5[A]) using eqn (4) as given below:

t_ion = (i_T − i_e)/i_T (4)

where, i_T is the total current and i_e is residual current. The t_ion was found to be ∼0.99, which shows that the total conductivity is the contribution of all the four charge entities (i.e. BMIM⁺, Li⁺, BF₄⁻ and TFSI⁻) present in the membranes. Thus, it is very necessary to calculate the cationic transport number (t_Li⁺).

Fig. 5 [A] DC polarization curves of symmetric cells: (a) SS|PGEM containing 70 wt% of IL|SS with applied voltage of 100 mV recorded at room temperature. [B] d.c. polarization curve at the applied voltage of 10 mV, and inset of [B] is the a.c. impedance plot before and after polarization of the cell (i.e. Li|PGEM containing 70 wt% of IL|Li) at room temperature. [C] Linear sweep voltagram (LSV) of cell: Li|PGEM containing 70 wt% of IL|LiMn₂O₄ at room temperature with scan rate of 10 mV s⁻¹.

The cationic transport number (t_Li⁺) of the prepared PGEMs is obtained by the combined a.c./d.c. techniques (as discussed earlier in Experimental section) using a Li/Li electrochemical cell comprising two Li metal electrodes. The transference numbers as determined by the Bruce–Vincent method³⁹ are shown in Fig. 5[B]. In this technique, a constant voltage ΔV = 10 mV was applied to polarize the cell (Li|PGEM|Li) and the resulting current (i.e. initial value (I_o) to a final value i.e. steady state value (I_ss)) w.r.t time was examined. The resistances of the cell before (R_o) and after (R_ss) the polarization at room temperature were also examined by a.c. impedance spectroscopy and are given in the inset of Fig. 5[B]. The expression given below has been used to evaluate the cationic transport number (t_Li⁺)

t_Li⁺ = (I_ss(ΔV − I_oR_o))/(I_o(ΔV − I_ssR_ss)) (5)

The cationic transport number (t_Li⁺) at room temperature for the PGEM containing 70% of IL has been found to be ∼0.33. Egashira et al.⁴⁰ reported that Li⁺ ion transport number of the system PEO-PMA/LiTFSI/HTMATFSI was 0.10–0.16. Kumar et al.⁴¹ found that Na⁺ ion transport number was ∼0.23 of the system PVdF-HFP/NaTf/EMITf.

The electrochemical stability window of prepared PGEMs was also investigated by linear sweep voltammetry (LSV). The electrochemical stability window of the prepared electrolytes is generally characterized by the reduction potential of its cation and the oxidation potential of its anion. The electrochemical stability window is the difference between the anodic and cathodic limiting potential.^42,43 The linear sweep voltammetry (LSV) curve of the cell Li|PGEM|LiMn₂O₄ at room temperature with the scan rate of 10 mV s⁻¹ are shown in Fig. 5[C]. The electrochemical stability window has been found ∼4.0–4.20 V. The ECW value sounds good for the practical application of prepared PGEMs especially in Li-ion conducting rechargeable battery application and shows good reversibility of the lithium ion insertion and extraction. No other signature was found in the voltage range thus confirming that no side reaction has occurred.

Thermal studies by DSC and TGA

Fig. 6 shows the DSC thermograms of pristine PVdF-HFP and polymer gel electrolyte membranes, (PVdF-HFP + 20 wt% LiTFSI) + x wt% BMIMBF₄ (where x = 10, 20, 30, 40, 50, 60 and 70). The melting peak of pristine PVdF-HFP is observed at 145 °C. From Fig. 6, it can be seen that the melting temperature (T_m) of the prepared membranes decreased with increasing IL content.

Fig. 6 DSC thermogram of (a) pristine PVdF-HFP and prepared PGEMs (PVdF-HFP + 20 wt% LiTFSI) + x wt% BMIMBF₄ (b) x = 10, (c) x = 20, (d) x = 30 (e) x = 40 (f) x = 50, (g) x = 60 and (h) x = 70.

It was expected that degree of crystallinity (X_c) decreases with IL content as evidenced from SEM and XRD analysis. Therefore, degree of crystallinity (X_c) was calculated using the relation given below:

where, ΔH_m = melting heat involved in the phase transition and = melting heat of 100% crystalline PVdF-HFP (i.e. for crystalline PVdF-HFP is 104.7 J g⁻¹).⁴⁴ The value of X_c decreases with increasing IL content (see Table 2).

Table 2 Degree of crystallinity (X_c) of the prepared PGEMs obtained by DSC

PGEMs ((PVdF-HFP + 20% LiTFSI) + x wt% BMIMBF4)	% degree of crystallinity (Xc)
x = 10	31.12
x = 20	22.55
x = 30	19
x = 40	18.9
x = 50	13.03
x = 60	11.9
x = 70	10

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For better understanding the role of IL in the membranes, we proposed a phenomenological model as shown in Fig. 7. Fig. 7 shows (a) semi crystalline (i.e. crystalline + amorphous) polymer membrane, (b) polymer electrolyte (polymer + salt), (c) PGEMs containing lower amount of added IL and (d) PGEMs containing higher amount of added IL. Fig. 7 explains that on the addition of LiTFSI salt, the polymer chain became flexible. Further on incorporation of IL in PGEMs reduces the intermolecular interaction between the polymer chains and hence increases the mobility and flexibility of polymer chain segment which increases the amorphicity of the membranes. Number density of charge carriers also increased upon increasing IL content that results increase in ionic conductivity of the membranes.

Fig. 7 A phenomenological model which explains the behaviour of ionic conductivity and decrease in degree of crystallinity of the prepared PGEMs upon increasing IL content.

Fig. 8 shows the thermogravimetric (TGA) and their first derivative (DTGA) curves of prepared PGEMs (PVdF-HFP + 20 wt% LiTFSI) + x wt% BMIMBF₄ (where x = 10, 20, 30, 40, 50, 60 and 70). From our recent studies,^24,25 we have found that (i) both pristine PVdF-HFP and pristine IL show single step decomposition and their corresponding decomposition peaks appeared at ∼470 °C and 465 °C respectively and (ii) the polymer electrolyte i.e. PVdF-HFP + 20 wt% LiTFSI exhibited three step decomposition mechanism with decomposition peaks at ∼468 °C (decomposition temperature (T_d) of polymer, PVdF-HFP), 340 °C (may be related to the T_d of complex formed between polymer with cation of Li-salt) and 310 °C (T_d of the LiTFSI salt). It has also been found by us in our earlier studies that when polymer PVdF-HFP was mixed with ILs (like BMIMBF₄ and BMIMTFSI) there is possibility of forming complex entities in which some part of IL complexed with polymer backbone and remaining IL was present as such in the membrane. But when IL was mixed in polymer electrolyte (i.e. PVdF-HFP containing LiTFSI salt), there is less chance to form complex as we have found in this system. From DTGA curves, it has been found that at lower IL content, the T_d of polymer PVdF-HFP, IL and polymer–IL complexed cannot be distinguished. While, at higher IL content (above than 30 wt%), multi decomposition peaks have been found that may be related to decomposition of polymer–salt complex (marked as *), decomposition of polymer (marked as β) and decomposition of polymer–IL complex (marked as α). But the intensity of the peaks related to decomposition of complex of IL with polymer, decomposition of complex of salt with polymer and the decomposition peak of polymer itself starts decreasing with increasing IL content and peaks almost disappear at higher content of IL and also intensity of decomposition peak of free IL (marked as γ) starts growing becomes maximum for higher content of IL. It may also be remarked that thermal stability of the prepared membranes also increased with increasing IL content that makes the prepared membranes thermally stable. Therefore, the polymer gel electrolyte membranes based on ILs are the suitable choice for practical application.

Fig. 8 TGA thermogram of PGEMs (PVdF-HFP + 20 wt% LiTFSI) + x wt% BMIMBF₄ (a) x = 10, (b) x = 20, (c) x = 30 (d) x = 40 (e) x = 50, (f) x = 60 and (g) x = 70.

Ion–polymer interaction by FTIR study

FTIR spectroscopy has been used to investigate the possible ion–polymer interaction. The FTIR spectra of (PVdF-HFP + 20% LiTFSI) + x wt% IL containing different amounts of ionic liquid is divided into two different regions 1000–450 cm⁻¹ and 3200–2800 cm⁻¹ given in Fig. 9[A] and [B] respectively. The FTIR spectra of pristine PVdF-HFP and pristine IL along with their assignments are given in insets of Fig. 9[A] and [B] respectively. These two regions are of our particular interest for the present discussion because we have found significant changes in the peak positions.

Fig. 9 [A] and [B] FTIR spectra of PGEMs (PVdF-HFP + 20 wt% LiTFSI) + x wt% BMIMBF₄ (a) x = 10, (b) x = 20, (c) x = 30 (d) x = 40 (e) x = 50, (f) x = 60 and (g) x = 70 in the spectral range 1000–400 cm⁻¹ and 3200–2800 cm⁻¹ respectively. Inset of [A] and [B] shows the FTIR spectra of pristine PVdF-HFP and pristine BMIMBF₄.

Since polymer PVdF-HFP is semi-crystalline in nature therefore its spectra contains vibrational bands related to crystalline as well as amorphous phase of polymer. The vibrational bands related to the crystalline phase (α-phase) of the polymer PVdF-HFP are observed at 489, 532, 613, 760, 795, and 974 cm⁻¹ whereas the vibrational bands related to amorphous phase (β-phase i.e. marked in red circle in the inset of Fig. 9[A]) of the polymer are observed at 840 and 879 cm⁻¹.^45–51 From Fig. 9[A] it can be seen that, almost all the intense bands which are related to the crystalline phase (i.e., α-phase 489, 532, 613, 760, 795, and 974 cm⁻¹) of pristine PVdF-HFP start disappearing and/or become weak whereas the peaks belonging to the amorphous phase (β-phase 840 and 879 cm⁻¹) (as shown by the dotted portion in Fig. 9[A]) become prominent upon incorporation of IL in the polymer PVdF-HFP/polymer electrolyte (PVdF-HFP + 20 wt% LiTFSI). Further, it has also been found that the intensity of some IL related peaks keep on increasing upon increasing content of IL (marked as * in Fig. 9[A]). This shows that IL increases the amorphicity of the polymer film as found from SEM/XRD and DSC studies discussed earlier.

From our earlier studies we have found that this region can give the information about the complexation due to incorporation of IL in the polymeric matrix. This region is further divided into two sections; section (A) and section (B). Section (A) is related to the C–H stretching vibrations of butyl chain of IL (also those of polymer backbone stretching) and section (B) is related to the imidazolium cation ring of IL (see Fig. 9[B]).

If complexation between the IL and polymer occurs, new peak should arise in this region and/or some significant change in the peak position of the band should occur. To see the complexation of IL and polymer we closely inspect the two regions marked section (A) and section (B) (in Fig. 9[B]) in the spectral range 3200–2800 cm⁻¹ one by one.

From section (A) of Fig. 9[B], it can be seen that C–H stretching vibrations related to the butyl chain of IL and that of polymer do not show any significant change. This shows that the butyl chain is not complexing with the polymer backbone as found in our earlier systems. In Fig. 9[B] section (B), bands are related to the imidazolium cation ring of IL. To know whether the imidazolium cation ring of IL complexes with the polymer backbone or not, we closely investigated the section [B] of Fig. 9[B] and carried out a detailed deconvolution of the spectra in the spectral range 3200 to 3070 cm⁻¹. The deconvolution was done with the help of Peakfit software⁵² using multiple Gaussian peaks to extract the exact peak positions of the prepared PGEMs. The deconvoluted spectra of pristine IL and prepared PGEMs containing different amounts of IL (with the value of square of regression coefficient i.e. r² = 0.999) are given in Fig. 10. The deconvoluted FTIR spectra of pristine IL (see Fig. 10[a]) consists four strong peaks at 3174, 3158, 3126 and 3104 cm⁻¹. From Fig. 10[b]–[d], it can be seen that there are no additional peaks appearing which indicate complexation. Therefore, from Fig. 10 it can be concluded that when IL, BMIMBF₄ and LiTFSI salt both are present in the membranes based on polymer PVdF-HFP there is no chance complexation as discussed in earlier section. The polymer matrix has a cross linked network structure, which acts as a container for the liquid electrolyte. Therefore, in the present study, polymer PVdF-HFP can be considered as a container for IL and salt. Salt dissociates into cations and anions which are free to move in the polymer network with the help of IL and executes ion transfer under external potential drive, which gives the ion conduction of the polymer electrolytes. The aforementioned structure of ion conductive polymer electrolytes has the advantage of high ionic conductivity comparable to the liquid electrolytes as they overcome the problem of leakage, portability, ease in preparation, synthesized in desired shapes and sizes, prevent electrodes from corrosion etc. which are associated with the liquid electrolytes, and at the same time, possessing sufficient mechanical strength and flexibility arising from the polymer matrix. Thus, on the basis of the above discussions it can be concluded that, both IL and LiTFSI salt increase the number of charge carriers by providing more free ions in prepared PGEMs which play a significant role in enhancing the conductivity of the system. Further, on increasing IL content we have found decrease in peak intensity of peaks related to the crystalline phase of polymer PVdF-HFP indicating the reduction in crystallinity of the PGEMs in the presence of IL.

Fig. 10 Deconvoluted FTIR spectra of PGEMs (PVdF-HFP + 20 wt% LiTFSI) + x wt% BMIMBF₄ (a) x = 20, (b) x = 40, (c) x = 60 and (d) x = 70 for CH stretching vibrational mode of imidazolium cation ring of IL in the region 3200–3070 cm⁻¹.

Conclusions

Lithium ion conducting PGEMs based on the PVdF-HFP/LiTFSI/BMIMBF₄ have been synthesized and characterized. The prepared PGEMs have high room temperature ionic conductivity (∼3.2 × 10⁻³ S cm⁻¹) comparable to the liquid electrolytes and are free standing and flexible in nature. The prepared PGEMs can also overcome the problem associated with the liquid electrolytes like leakage, portability and can be synthesized in desired shapes and sizes, they also prevent electrodes from corrosion etc. The pristine PVdF-HFP possess a semi-crystalline structure and its amorphicity increases with the addition of LiTFSI salt and BMIMBF₄ ionic liquid. The ionic conductivity of the polymeric gel membranes was found to increase as the amount of IL increased in the membranes. A high total ionic transference number ∼0.99 signifying that the conductivity is predominantly ionic in nature and the cationic transference number (t_Li⁺) ∼ 0.33 with a wider electrochemical window (ECW) ∼4.0–4.20 V for the PGEMs containing higher loading of IL (∼70 wt% of IL) have been obtained.

Acknowledgements

One of us R. K. S. gratefully acknowledge financial support from BRNS-DAE, Mumbai and DST, New Delhi, India to carry out this work.

References

1. F. M. Gray, Polymer Electrolytes: Fundamental and Technological Applications, VCH Publishers, New York, 1997.
2. N. Mahmood, M. Islam, A. Hameed, S. Saeed and A. Nawaz Khan, J. Compos. Mater., 2013, 1.
3. P. G. Bruce and C. A. Vincent, Polym. Electrolytes, 1993, 89, 3187.
4. P. Novák, K. Müller, K. S. V. Santhanam and O. Haas, Chem. Rev., 1997, 97, 207.
5. W. H. Meyer, Polymer Electrolytes for Lithium-Ion Batteries, Adv. Mater., 1998, 10, 439.
6. N. Mahmood and Y. Hou, Adv. Sci., 2014, 1, 1400012.
7. A. S. N. Smara, M. Z. Kufian, S. R. Majid and A. K. Arof, Electrochim. Acta, 2011, 57, 91.
8. P. E. Stallworth, S. G. Greenbaum, F. Croce, S. Slanet and M. Salomon, Electrochim. Acta, 1995, 40, 2137.
9. S. K. Chaurasia, R. K. Singh and S. Chandra, J. Raman Spectrosc., 2011, 42, 2168.
10. A. Martinelli, A. Matic, P. Jacobsson, L. Borjesson, A. Fernicola, S. Panero, B. Scrosati and H. Ohno, J. Phys. Chem. B, 2007, 111, 12462.
11. K. Tsunemi, H. Ohno and E. Tsuchida, Electrochim. Acta, 1983, 28, 591.
12. H. Ohno, Macromol. Symp., 2007, 551, 249.
13. X. Tian, B. Zhu and Y. Xu, J. Membr. Sci., 2005, 248, 109.
14. P. L. Kuo, C. H. Tsao, C. H. Hsu, S. T. Chen and H. Hsu, J. Membr. Sci., 2016, 499, 462.
15. D. H. Lim, J. Manuel, J. H. Ahn, J. K. Kim, P. Jacobsson, A. Matic, J. K. Ha, K. K. Cho and K. W. Kim, Solid State Ionics, 2012, 225, 631.
16. G. P. Pandey, R. C. Agrawal and S. A. Hashmi, J. Power Sources, 2009, 190, 563.
17. I. Osada, H. D. Vries, B. Scrosati and S. Passerini, Angew. Chem., Int. Ed., 2016, 55, 500.
18. J. M. Tarascon, A. S. Gozdz, C. Schmutz, F. Shokoohi and P. C. Warren, Solid State Ionics, 1996, 86, 49.
19. S. Ramesh and O. P. Ling, Polym. Chem., 2010, 1, 702.
20. J. Y. Song, Y. Y. Yong and C. C. Wan, J. Electrochem. Soc., 2000, 14, 73219.
21. G. B. Appetecchi, F. Croce and B. Scrosati, Electrochim. Acta, 1995, 40, 991.
22. Y. S. Ye, J. Rick and B. J. Hwang, J. Mater. Chem. A, 2013, 1, 2719.
23. D. R. MacFarlane, N. Tachikawa, M. Forsyth, J. M. Pringle, P. C. Howlett, G. D. Elliott Jr, J. H. Davis, M. Watanabe, P. Simon and C. A. Angell, Energy Environ. Sci., 2014, 7, 232.
24. Y. Lu, S. S. Moganty, J. L. Schaefer and L. A. Archer, J. Mater. Chem., 2012, 22, 4066.
25. K. Yin, Z. Zhang, L. Yang and S. I. Hirano, J. Power Sources, 2014, 258, 150.
26. Y. Lu, K. Korf, Y. Kambe, Z. Tu and L. A. Archer, Angew. Chem., 2014, 126, 498.
27. Shalu, S. K. Chaurasia, R. K. Singh and S. Chandra, J. Appl. Polym. Sci., 2015, 132, 41456.
28. Shalu, V. K. Singh and R. K. Singh, J. Mater. Chem. C, 2015, 3, 7305.
29. S. K. Chaurasia, R. K. Singh and S. Chandra, Vib. Spectrosc., 2013, 68, 190.
30. E. M. Fahmi, A. Ahmad, N. N. M. Nazeri, H. Hamzah, H. Razali and M. Y. A. Rahman, Int. J. Electrochem. Sci., 2012, 7, 5798.
31. D. Saikia and A. Kumar, Electrochim. Acta, 2004, 49, 2581.
32. D. Saikia, C. C. Han and Y. W. Chen-Yang, J. Power Sources, 2008, 185, 570.
33. F. Alloin, A. D'Aprea, N. El Kissi, A. Dufresne and F. Bossard, Electrochim. Acta, 2010, 55, 5186.
34. Z. Wang, W. Gao, X. Huang, Y. Mo and L. Chen, Electrochem. Solid-State Lett., 2001, 4, A132.
35. D. Benrabah, J. Y. Sanchez and M. Armand, Solid State Ionics, 1993, 60, 87.
36. A. Hofmann, M. Schulz and T. Hanemann, Electrochim. Acta, 2013, 89, 823.
37. M. A. Navarra, J. Manzi, L. Lombardo, S. Panero and B. Scrosati, ChemSusChem, 2011, 4, 125.
38. H. Ye, J. Huang, J. J. Xu, A. Khalfan and S. G. Greenbaum, Li ion conducting polymer gel electrolytes based on ionic liquid/PVDF-HFP blends, J. Electrochem. Soc., 2007, 154, A1048.
39. P. G. Bruce and C. A. Vincent, J. Electroanal. Chem., 1987, 225, 1.
40. M. Egashira, H. Todo, N. Yoshimoto and M. Morita, J. Power Sources, 2008, 178, 729.
41. D. Kumar and S. A. Hashmi, Solid State Ionics, 2010, 181, 416.
42. S. Fang, Y. Tang, X. Tai, L. Yang, K. Tachibana and K. Kamijima, J. Power Sources, 2011, 196, 1433.
43. K. Yin, Z. Zhang, X. Li, L. Yang, K. Tachibana and S. Hirano, J. Mater. Chem. A, 2015, 3, 170.
44. C. W. Liew, Y. S. Ong, J. Y. Lim, C. S. Lim, K. H. Teoh and S. Ramesh, Int. J. Electrochem. Sci., 2013, 8, 7779.
45. X. Tian, X. Jiang, B. Zhu and Y. Xu, J. Membr. Sci., 2006, 279, 479.
46. M. Kobayashi, K. Tashiro and H. Tadokor, Macromolecules, 1975, 8, 158.
47. Z. Li, G. Su, D. Gao, X. Wang and X. Li, Electrochim. Acta, 2004, 49, 4633.
48. S. Abberent, J. Plestil, D. Hlavata, J. Lindgren, J. Teganfeldt and A. Wendsjo, Polymer, 2001, 42, 1407.
49. K. Wang, H. Lee, R. Cooper and H. Liang, Appl. Phys. A: Mater. Sci. Process., 2009, 95, 435.
50. R. D. Simoes, A. E. Job, D. L. Chinaglia, V. Zucolotto, J. C. Camargo-Filho, N. Alves, J. A. Giacometti Jr, O. N. Oliveira and C. J. L. Constantino, J. Raman Spectrosc., 2005, 36, 1118.
51. Shalu, S. K. Chaurasia, R. K. Singh and S. Chandra, J. Phys. Chem. B, 2013, 117, 897.
52. Peakfit is a commercially available product from Seasolve Software Inc.

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