Nano-size LiAlO2 ceramic filler incorporated porous PVDF-co-HFP electrolyte for lithium-ion battery applications

Electrochimica Acta 52 (2007) 4987–4993

Nano-size LiAlO2 ceramic filler incorporated porous PVDF-co-HFP electrolyte for lithium-ion battery applications N.T. Kalyana Sundaram, A. Subramania ∗ Advanced Materials Research Lab, Department of Industrial Chemistry, Alagappa University, Karaikudi 630 003, India Received 11 December 2006; received in revised form 17 January 2007; accepted 29 January 2007 Available online 6 February 2007

Abstract The optimized composition of PVdF-co-HFP–LiAlO2 based micro-porous nano-composite polymer electrolyte membranes (MPNCPEMs) was prepared with a preferential polymer dissolution process. Nitrogen adsorption isotherms and SEM micrographs showed that the enhanced ionic conductivity of polymer electrolyte was due to increase in pore-size, surface area and pore density, results an increase in the electrolyte uptake. The ac-impedance spectroscopy showed that the room temperature ionic conductivity of PVdF-co-HFP–LiAlO2 based polymer electrolyte membranes increased with the removal of PVA content and attained the maximum ionic conductivity of 8.12 × 10−3 S cm−1 . The prepared MPNCPEM of high ionic conductivity was subjected into LSV study. Finally, the electrode/electrolyte interfacial resistance was evaluated by monitoring the impedance response at different time intervals. © 2007 Published by Elsevier Ltd. Keywords: Ceramic fillers; Ionic conductivity; Polymer gels; Li-ion battery; BET-analysis

1. Introduction The development of polymeric systems with high ionic conductivity is one of the main goals in polymer research. This is due to their potential applications as electrolytes in solid-state batteries [1–3]. The main advantage of polymer electrolytes is favourable mechanical properties, ease of fabrication of thin film of desirable size and an ability to form effective electrode–electrolyte contacts. Armand et al. [4] have claimed that the crystalline complexes formed from alkali-metal salts with polyethylene oxide (PEO) are capable of demonstrating significant ionic conductivity and have highlighted their possible application as battery electrolytes. Nowadays, for all solidstate lithium batteries, poly(vinylidene fluoride-co-hexafluoro propylene) (PVdF-co-HFP) has been considered to be one of the most promising candidates [5–7]. At an ambient temperature, the PVdF-co-HFP based polymer electrolytes exhibited good dimensional stability. However, the PVdF-co-HFP polymer electrolytes showed practical ionic conductivity only at higher temperatures, and their melting points, and at such high

∗

Corresponding author. E-mail address: [email protected] (A. Subramania).

0013-4686/$ – see front matter © 2007 Published by Elsevier Ltd. doi:10.1016/j.electacta.2007.01.066

temperature, they exists in a quasi-liquid state and become very flexible, and therefore showed very poor dimensional stability. A dimensional weak polymer electrolyte film easily causes a short circuit between a cathode and an anode when it is applied to all solid-state lithium battery. Therefore, it is still an effective research topic to develop dimensionally stable polymer electrolytes with high ionic conductivity especially at room temperature. Studies have been made primarily on the enhancement of the ionic conductivity at room temperature via various approaches such as blends, co-polymers, comb-shape polymers, crosslinked networks and incorporation of ceramic fillers onto the polymer matrix. According to Scrosati and co-workers [8,9] and Wieczorek et al. [10] the Lewis acid–base interaction plays a vital role in the enhancement of ionic conductivity, electrochemical stability and interfacial stability of the composite polymer electrolytes. Moreover, the porous texture has proved an effective approach to improve the over all performances of polymer electrolyte, which usually showed good mechanical performances with enhanced ionic conductivity at room temperature besides the activation process requires critical moisture control only during the last activation step because of high water sensitivity of lithium salt [11–15]. The LiAlO2 can be used as best ceramic fillers for polymer electrolyte [16–22]. Hence,

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an attempt is made to prepare PVdF-co-HFP–LiAlO2 based micro-porous nano-composite polymer electrolyte membranes (MPNCPEMs) prepared with a preferential polymer dissolution process and the results are described herein. 2. Experimental 2.1. Materials and its pretreatment PVdF-co-HFP with an average molecular weight of greater than 500,000 (Aldrich, USA), PVA with an average molecular weight of greater than 70,000 (Aldrich, USA), LiAlO2 with an particles of less than ∼50 nm (Aldrich, USA) and LiClO4 (E-Merck, Germany) were dried under vacuum oven at 100 ◦ C under 10−3 Torr pressure for 48 h. Ethylene carbonate (EC) and diethyl carbonate (DEC) (Across organic, Belgium) were purified by distillation under reduced pressure. 2.2. Optimization of nano-sized LiAlO2 ceramic filler onto a PVdF-co-HFP matrix First, the PVdF-co-HFP co-polymer was dissolved in acetone. To this, different wt.% of nano-sized LiAlO2 ceramic filler (0, 2, 4, 6, 8 and 10 wt.%) were added. The resultant viscous solutions were spread as film on a glass substrate using doctor blade. Then, the nano-composite polymer films (NCPFs) were dried at 50 ◦ C in a vacuum oven under 10−3 Torr pressure for 6 h to remove any further traces of acetone. The thickness of the films were obtained in the range of 50–80 ␮m. Finally, the resultant nano-composite polymer films were subjected into differential scanning calorimetric analysis (DSC) to optimize the LiAlO2 ceramic fillers content onto a PVdF-co-HFP matrix. 2.3. Preparation of PVdF-co-HFP–LiAlO2 based MPNCPEMs The optimized composition ratio of PVdF-co-HFP with nanosized ceramic filler (LiAlO2 ) were taken and dissolved in DMF. To this different wt.% of PVA (0, 5, 10, 15 and 20 wt.%) were also added. The resultant viscous solutions were spread as films on a glass substrate using doctor blade. These resultant NCPFs were dried at 80 ◦ C in a vacuum oven under 10−3 Torr pressure for 6 h to remove any further traces of DMF. The thickness of the films was obtained in the range of 50–80 ␮m. The prepared cast NCPFs was immersed in a pool of excess of deionized water at 60 ◦ C to remove PVA from PVdF-co-HFP matrix to form microporous structure with different distribution ratios. These microporous nano-composite polymer membranes (MPNCPMs) were dried in a vacuum oven at 80 ◦ C under 10−3 Torr pressure for at least 6 h. Finally, the MPNCPMs were activated by soaking in 1 M LiClO4 electrolyte solution containing 1:1 (v/v) ratio of EC and DEC for 18 h to get MPNCPEMs. 2.4. Physical characterization SEM studies were carried out with JOEL-scanning electron microscope (JSM-35CF) to find out the surface morphology of

polymer film/membrane. The porosity (P) was determined using the following equation [23]: P=

ma /ρa ma /ρa + mp /ρp

(1)

where ma is the weight of MPNCPMs after impregnation with 1butanol and mp is the weight of MPNCPMs before impregnation with 1-butanol. Similarly, ρa and ρp are density of 1-butanol and the MPNCPMs, respectively. Swelling behaviour of the MPNCPMs was studied inside the glove box. The extent of swelling (Sw) of MPNCPMs was determined to investigate its dependence on micro-porous structure. The percentage of swelling was determined using Eq. (2) [24]: Sw =

W − W0 × 100 W0

(2)

where W0 is the weight of dried membrane and W is the weight of swelled membrane. The Brunauer–Emmett–Teller (BET) pore-size distributions of polymer membranes were determined using a continuousflow gas (N2 ) adsorption apparatus (Model ASAP 2000, Micromeritics Instrument (Corp.)). DSC experiments of NCPFs and MPNCPEMs were carried out using a Dupont-DSC analyzer (Model TA-2000) over a temperature range of 0–200 ◦ C at a scan rate of 10 ◦ C/min under N2 atm. The samples were sealed in Al crucible inside the glove box. The sealed samples were taken out from the glove box at the time of DSC experiments. The thermograms were base line corrected and calibrated against Indium metal. The crystallinity of the MPNCPEMs was calculated using the following equation from the DSC curves [25]:   Hm × 100 (3) Xc (%) = φ Hm φ

where Hm is the crystalline melting heat of pure PVdF (104.7 J/g) and Hm is the heat of fusion of MPNCPEMs. It can be calculated from the integral area of the base line and each melting curve. 2.5. Electrochemical characterization The ionic conductivity measurements were performed by sandwiching the MPNCPEM in between two stainless steel electrodes using HIOKI 3522-50 LCR meter over a frequency range of 1 mHz to 100 kHz at a scan rate of 1 mV/s in various temperatures ranging from 298 to 353 K. The electrochemical stability of MPNCPEM of high ionic conductivity was evaluated with cell featuring a stainless steel (SS) as working electrode and lithium as counter and reference electrodes by linear sweep voltammetry at 25 ◦ C using an EG&G Electrochemical analyzer (Model-6310) in the scan rate of 1 mV/s was performed to study the system under investigation. The cell was assembled in a glove box under argon atmosphere. The interfacial stability of PVdF-co-HFP based gel polymer electrolyte (GPE), nano-composite polymer electrolyte

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(NCPE) and MPNCPEM of high ionic conductivity system were evaluated. A typical cell consist of polymer electrolyte was sandwiched between two lithium-unblocking electrodes to form a symmetrical Li/electrolyte/Li cell used for test [26]. The cell was assembled and sealed in an argon filled glove box. Lithium electrode/electrolyte interfacial resistance was evaluated by monitoring the impedance response using an EG&G Electrochemical analyzer (Model-6310). The cell was stored under open circuit condition at a fixed temperature during the whole testing time. The electrode formulation consist of 80 wt.% LiCoO2 (from commercial market), 10 wt.% of carbon black, and 10 wt.% binder. The load of LiCoO2 was about 10 mg cm−2 . The prepared electrode pellets were dried at 120 ◦ C under vacuum for 48 h. The charge–discharge behaviour of the carbon/MPNCPEM/LiCoO2 cell was studied at the current density of 0.25 mA cm−2 . 3. Results and discussion 3.1. Optimization of PVdF-co-HFP–LiAlO2 based NCPFs DSC analysis of PVdF-co-HFP with different wt.% of LiAlO2 (0, 2, 4, 6, 8 and 10 wt.%) during the heating scan are displaced in Fig. 1(a)–(e). The endothermic peak obtained at 150.0 ◦ C is attributed to the melting point of semicrystalline PVdF-co-HFP. It can be seen from Fig. 1(a)–(e) that the melting temperature, heat of fusion of PVdF-co-HFP matrix decreased slightly, resulted crystallinity decreases (Fig. 2) when LiAlO2 ceramic filler is incorporated into the PVdF-co-HFP polymer matrix up to 6 wt.%. This may be attributed to interaction of polymer backbone and LiAlO2 ceramic filler, which gives disorder structure [27,28]. Hence, the crystallinity of PVdF-co-HFP matrix is decreased. Therefore, the melting temperature (Tm ) of semicrystalline PVdF-co-HFP decreased from 150.0 to 143.6 ◦ C when the content of LiAlO2 increased from 0 to 6 wt.% to

Fig. 1. DSC curve of NCPFs containing different wt.% of LiAlO2 content: (a) 0 wt.% of LiAlO2 ; (b) 2 wt.% of LiAlO2 ; (c) 4 wt.% of LiAlO2 ; (d) 6 wt.% of LiAlO2 ; (e) 8 wt.% of LiAlO2 ; (f) 10 wt.% of LiAlO2 .

Fig. 2. Crystallinity and heat of fusion of NCPFs containing different wt.% of LiAlO2 content.

the polymer matrix. Beyond this concentration, the melting temperature increased from 142.4 to 144.8 ◦ C. The increase in melting temperature (Tm ) of PVdF-co-HFP–LiAlO2 matrix can be attributed to decreasing the flexibility of PVdF-co-HFP polymer chain segments which increasing the heat of fusion (Hm ) and crystallinity (Xc ) as shown in Fig. 2. It is due to the steric hindrance caused by both PVdF-co-HFP and LiAlO2 will induce crystalline site in the polymer matrix, leading to lower segmental mobility [27]. It suggests that the incorporation of 6 wt.% of LiAlO2 into the PVDF-co-HFP matrix has been taken as an optimum composition for the preparation of MPNCPEMs. 3.2. SEM studies The SEM photographs were taken for the optimized composition of PVdF-co-HFP–LiAlO2 based nano-composite polymer matrix is subjected into before and after removal of 20 wt.% of PVA content (Fig. 3(a) and (b)). It can be seen from Fig. 3(a) that the before removal of PVA content in the PVdF-co-HFP–LiAlO2 matrix, could not observe any porous texture besides the LiAlO2 is homogeneously dispersed in the PVdF-co-HFP–PVA matrix. The MPNCPM obtained by the removal of 20 wt.% of PVA in the PVdF-co-HFP–LiAlO2 matrix shows the uniform porous structure with the pore diameter of 305 nm as shown in Fig. 3(b). It reveals that the polymer membrane prepared with the preferential polymer dissolution process showed a denser structure with uniform porosity. From this study, we found that the porosity of the polymer film depends on the preferential removal of polymer used. In addition, the micro-pores in the PVdF-coHFP–LiAlO2 membrane are well interconnected with strong pore walls besides LiAlO2 ceramic fillers can be seen in the holes of membrane, making down surface of membrane and there is no LiAlO2 aggregation in the upper surface, indicating that the nano-sized powder is homogeneously dispersed in the polymer matrix.

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Fig. 4. Porosity of MPNCPMs as a function of removal of different wt.% of PVA.

MPNCPM, we used the gas adsorption/desorption method. The average pore-size, BET surface area and porosity of the membranes are listed in Table 1 as a function of different wt.% of removal of PVA content. From this table it is observed that the larger the area and pore-size, the higher the porosity of the membrane will be. There, the sequence of the porosity increases with removal of PVA content in the polymer matrix which shows the same tendency as that of the electrolyte uptake.

Fig. 3. SEM photographs of NCPFs subjected into (a) before removal of 20 wt.% of PVA content and (b) after removal of 20 wt.% of PVA content.

3.3. Porosity and swelling behaviour The porosity of optimized composition of PVdF-coHFP–LiAlO2 based MPNCPMs was measured using Eq. (1). The porosity of MPNCPMs is shown in Fig. 4. From this figure it is observed that the removal of PVA content in the PVdF-co-HFP–LiAlO2 –PVA matrix create the porous polymer membrane. The increase in removal of PVA content causes an increase in porosity which is evident from SEM images is shown in Fig. 3. Hence, the uptake of electrolyte solution by PVdFco-HFP–LiAlO2 based MPNCPM increases with increase in removal of PVA content is shown in Fig. 5. To further understand the porous structure of the PVdF-co-HFP–LiAlO2 based

Fig. 5. Swelling of MPNCPMs as a function of removal of different wt.% of PVA.

Table 1 Micro-structure-related properties of PVdF-co-HFP–LiAlO2 based MPNCPM as a function of removal of different wt.% of PVA content Removal of PVA in wt.%

BET average pore diameter (nm)

BET surface area (m2 g−1 )

Porosity (%)

Ea (kJ/mol)

0 5 10 15 20

65 102 166 235 305

41.0 61.4 174.7 228.9 336.1

38.05 63.62 73.60 81.65 86.95

18.27 18.17 18.12 17.87 17.62

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may suggest that the crystallinity of polymer matrix is changed when liquid electrolyte is added into the polymer matrix. 3.5. Ionic conductivity studies

Fig. 6. DSC curve of MPNCPEMs as a function of removal of different wt.% of PVA: (a) 0 wt.% of PVA; (b) 5 wt.% of PVA; (c) 10 wt.% of PVA; (d) 15 wt.% of PVA; (e) 20 wt.% of PVA.

3.4. DSC analysis The thermal properties of optimized composition of PVdFco-HFP–LiAlO2 based MPNCPEMs were tested with DSC analysis. The DSC curves of MPNCPEMs are shown in Fig. 6(a)–(e) as a function of the removal of PVA content. The thermogram of PVdF-co-HFP–LiAlO2 based MPNCPEM exhibits an endothermic peak at 141.5 ◦ C. Further, the depression of melting temperature of PVdF-co-HFP–LiAlO2 matrix is observed when the electrolyte uptake of the polymer matrix increases (Fig. 6). In addition, the heat of fusion (Hm ) and crystallinity (Xc ) of MPNCPEM decreased as the electrolyte content increases (Fig. 7). The drop in heat of fusion (Hm ) and crystallinity (Xc ) of MPNCPEMs indicating that the increase in amorphous phase of PVdF-co-HFP–LiAlO2 matrix, due to interaction polymer backbone and liquid electrolytes [29,30]. It

Fig. 7. Crystallinity (Xc ) and heat of fusion (Hm ) of MPNCPEMs as a function of removal of different wt.% of PVA.

Fig. 8 displays the ionic conductivity of optimized composition of PVdF-co-HFP–LiAlO2 based MPNCPEMs as a function of different wt.% of removal of PVA content. It is worth to note that the ionic conductivity exhibits a same in tendency of porosity and the liquid electrolyte uptake are shown in Figs. 4 and 5. It is observed from the figure that an increase in removal of PVA content in the PVdF-co-HFP–LiAlO2 polymer matrix, increase the ionic conductivity. Porosity, electrolyte uptake and ionic conductivity of optimized composition of PVdF-co-HFP–LiAlO2 matrix is about 86.95%, 121.00% and 8.12 × 10−3 S cm−1 at room temperature, respectively, when the MPNCPEM is obtained by the removal of 20 wt.% of PVA content. Fig. 8 shows an Arrhenius plots of PVdF-co-HFP–LiAlO2 based MPNCPEMs. It is quite obvious from the figure that the ionic conductivity of MPNCPEMs increases with increase in temperature and the removal of PVA content. Nevertheless, the ionic conduction mainly depends on the entrapped liquid phase in a fully interconnected pore structure. These curves appear linear, so the apparent activation energy for the ions transport (Ea ) are obtained using the Arrhenius model σ = σ 0 exp(−E/RT), where R, T, σ and σ 0 are gas constant, temperature, the ionic conductivity of MPNCPEMs and the pre-exponential factor, respectively. According to this equation, the activation energy for the ions transport can be calculated from the slope of imitated straight line and are given in Table 1. From this table it is suggested that the activation energies of all samples are almost equal and that the variation in average pore-size does not influence the activation energy for ionic conduction. Again the conductivity data show a similar trend to that in the pore-size, porosity

Fig. 8. Arrhenius plots of MPNCPEMs as a function of removal of different wt.% of PVA. () 0 wt.% of PVA; () 5 wt.% of PVA; () 10 wt.% of PVA; () 15 wt.% of PVA; () 20 wt.% of PVA

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seems to be associated with the growth of passivation layer on the lithium electrodes and the degradation of the physical contact between the polymer electrolyte and lithium electrodes [31,32]. 3.8. Charge–discharge studies Charge–discharge curves for a carbon/MPNCPEM/LiCoO2 cell cycled between 3.0 and 4.2 V at a current density of 0.25 mA cm−2 are given in Fig. 11. The cell delivers a discharge capacity of 148.0 and 142.5 mAh g−1 with the coulombic efficiency of 99.86% and 98.95% on the first and the 25th cycle, respectively. The cell shows good capacity retention due to the good interfacial contact in the composite electrode with MPNCPEM.

Fig. 9. Linear sweep voltammetry curve of MPNCPEM of high ionic conductivity system.

and electrolyte uptake of the membrane, the higher the ionic conductivity will be. 3.6. Electrochemical stability studies One of the important parameters in the characterization of PVdF-co-HFP–LiAlO2 based MPNCPEM is the electrochemical stability window. This parameter can be determined by means of a linear sweep voltammetry (LSV). The resulting voltammograms for SS/MPNCPEM/Li cell is shown in Fig. 9. The onset voltage for anodic current is determined at around 5.2 V versus Li/Li+ which is assumed to be the decomposition voltage of the polymer electrolyte. 3.7. Interfacial stability Interfacial stability with electrode is an essential factor to guarantee acceptance performance in the electrochemical devices. The interfacial stability between the electrolyte and the lithium electrodes was investigated by monitoring the impedance response of a symmetrical Li/electrolyte/Li cell for the period of 20 h. Fig. 10(a)–(c) shows the impedance spectra for the cell based on PVdF-co-HFP based GPE, NCPE and MPNCPEM of high ionic conductivity systems as a function of storage time. The interfacial resistance of PVdF-co-HFP based GPE and NCPE is rapidly increased with various storage times is shown in Fig. 10(a) and (b). However, the interfacial resistance between MPNCPEM and lithium electrode is relatively stable during the same storage time, its value ranges from 137.66 to 482.64  is shown in Fig. 10(c). The initial interfacial resistance value for the MPNCPEM is also much lower than NCPE besides GPE. This may be due to better adhesion between the PVdF-co-HFP–LiAlO2 based MPNCPEM and lithium electrodes. The mechanism of resistance decrease with increasing time needs to be further investigated. The higher interfacial resistance of the GPE and NCPE during the storage time

Fig. 10. The interfacial resistances of PVdF-co-HFP based (a) GPE, (b) MPGPEM and (c) MPNCPEM.

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Fig. 11. Charge–discharge behaviour of carbon/LiCoO2 test coin cell using as prepared MPNCPEM of high ionic conductivity system.

4. Conclusions The addition of LiAlO2 onto PVdF-co-HFP to obtain microporous nano-composite polymer electrolyte membranes with a preferential polymer dissolution process has high ionic conductivity in the range of 8.12 × 10−3 S cm−1 at room temperature and its corresponding membrane porosity and electrolyte uptake was about 86.95% and 121.00%, respectively. The ionic conductivity has gone through a maximum when the removal of PVA content in the polymer matrix took place. The electrochemical stability window of PVdF-co-HFP–LiAlO2 based MPNCPEM is about 5.2 V versus Li/Li+ . The MPNCPEM has lowered the interfacial resistance when lithium metal was used as the anode. To conclude, the PVdF-co-HFP–LiAlO2 based MPNCPEM can be used as a promising candidate for high voltage rechargeable lithium-ion polymer batteries. Acknowledgement The authors gratefully acknowledge the DST (New Delhi), for the financial support. References [1] J. Xi, X. Qiu, X. Ma, M. Cui, J. Yang, W. Zhu, L. Chen, Solid State Ionics 176 (2005) 1249.

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