Optimization of filler type within poly(vinylidene fluoride-co-trifluoroethylene) composite separator membranes for improved lithium-ion battery performance

Composites Part B 96 (2016) 94e102

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Optimization of filler type within poly(vinylidene fluoride-co-trifluoroethylene) composite separator membranes for improved lithium-ion battery performance ~o Nunes-Pereira a, 1, Manab Kundu b, 1, Attila Go € ren a, c, Maria Manuela Silva c, Joa a, c, * b ndez a, d , Lifeng Liu , Senentxu Lanceros-Me Carlos M. Costa a

Centro/Departamento de Física, Universidade do Minho, Campus de Gualtar, 4710-057 Braga, Portugal International Iberian Nanotechnology Laboratory (INL), Av. Mestre Jose Veiga, 4715-330 Braga, Portugal Centro/Departamento de Química, Universidade do Minho, Campus de Gualtar, 4710-057 Braga, Portugal d gico de Bizkaia, 48160 Derio, Spain BCMaterials, Parque Científico y Tecnolo b c

a r t i c l e i n f o

a b s t r a c t

Article history: Received 21 January 2016 Accepted 14 April 2016 Available online 22 April 2016

Porous poly(vinylidene fluoride-co-trifluoroethylene) (P(VDF-TrFE)) based composite membranes filled with clays (montmorillonite, MMT), zeolites (Y zeolite, NaY), ceramics (barium titanate, BaTiO3) and carbonaceous (multiwalled-carbon nanotubes, MWCNT) fillers were prepared by solvent casting at room temperature. It is shown that, the thermal, mechanical and electrochemical properties of the membranes are not significantly affected by the presence of the fillers. On the other hand, the overall electrochemical behavior of the separator membranes improves with the inclusion of fillers, with respect to the pure polymer, as shown by the increase of the room temperature ionic conductivity for the composite membranes. Furthermore, cathodic half-cells based on composite membranes showed higher capacity retention and rate performance after 50 cycles than pristine polymer membranes. It is thus concluded that filler type deeply affects membrane separator performance in lithium-ion batteries, the P(VDF-TrFE) membrane with MMT filler being the one with the best performance among the evaluated fillers. © 2016 Elsevier Ltd. All rights reserved.

Keywords: A. Nano-structures B. Electrical properties B. Physical properties B. Chemical properties

1. Introduction Lithium-ion batteries are increasingly used in consumer electronics applications, including mobile phones, computers and toys, and in the automotive industry [1e4]. Industrial development and implementation is based on intensive R&D devoted to new and more efficient materials for the different lithium-ion battery components in order to increase energy and power density and safety and to reduce production costs [5,6]. Polymer electrolytes (PE) are applied in many types of lithium-ion batteries [7,8] since they are free from leakage, show flexible geometry and simple packing properties when used in energy storage systems [9,10]. The “classic” classification of PE was proposed by Gray in 1997 [11] and

* Corresponding author. Centro/Departamento de Física, Universidade do Minho, Campus de Gualtar, 4710-057 Braga, Portugal. E-mail address: cmscosta@fisica.uminho.pt (C.M. Costa). 1 Equal contribution. http://dx.doi.org/10.1016/j.compositesb.2016.04.041 1359-8368/© 2016 Elsevier Ltd. All rights reserved.

an updated reorganization of this class of materials [12] lead to the denomination of separator/electrolyte when lithium-ion battery applications are involved [13]. Thus PE are divided into different types: solid polymer electrolytes (SPE), gel polymer electrolytes (GPE) and composite polymer electrolytes (CPE) [12]. SPE consist in salts dissolved in a polymeric matrix. GPE are obtained in two steps: first the salts are dissolved into a polar or ionic liquid and then added to the host polymer to provide an adequate mechanical stability. CPE are very similar to SPE but with the inclusion of different (nano)fillers (inert oxide ceramic, molecular sieves, metallic, etc.) dispersed in a polymer matrix to improve the mechanical, thermal and/or electrochemical properties [14,15]. For lithium-ion battery applications, the separator/electrolyte allows the movement of the ions between anode and cathode during battery charge and discharge, and ensures the electronic insulation between them [16]. The main parameters that determine the performance of a separator are permeability, porosity/pore size, electrolyte absorption and retention, chemical, mechanical and thermal stability [13,17].

J. Nunes-Pereira et al. / Composites Part B 96 (2016) 94e102

Different morphologies of PE have been obtained through linear and branched polymers, block copolymers and plasticized polymers and organic/inorganic hybrids [18,19]. Among the most interesting polymers for battery PE [20], are poly(vinylidene fluoride) (PVDF) and copolymers (poly(vinylidene fluoride-cotrifluoroethylene) e P(VDF-TrFE); poly(vinylidene fluoride-cohexafluoropropene) e P(VDF-HFP) and poly(vinylidene fluorideco-chlorotrifluoroethylene) e P(VDF-CTFE)) [13]. PVDF and its copolymers show strong advantages when compared to other polymers, e.g., polyolefins, due to their high dipolar moment and dielectric constant, semi-crystallinity, chemical resistance, mechanical strength, good contact between electrode and electrolyte, possibility of porosity control through binary and ternary polymer/ solvent systems and high anodic stability due to the presence of strong electron-withdrawing groups (eCeFe) [13,21e23]. P(VDFTrFE) has been used to develop single SPE with improved thermal and electrochemical performance when compared to conventional commercial separators and also with respect to PVDF and P(VDFHFP), as it shows higher ionic conductivity, thermal stability and high cyclability in anodic (Li/SneC) and cathodic (Li/LiFePO4) halfcells [22]. P(VDF-TrFE) shows a lower degree of crystallinity when compared to PVDF and higher mechanical properties when compared to P(VDF-HFP), important requirements for battery separators [13]. P(VDF-TrFE) based CPE also show high flexibility and good electrochemical stability [15,24e28]. In CPE, fillers can participate or not in the ionic conduction process depending on their type. The fillers are typically selected in order to improve one or more properties of the separator, and the most used fillers are inert oxide ceramics (Al2O3, TiO2, etc.) [29], molecular sieves and zeolites [30], ferroelectric materials [31] and carbon based fillers [32]. In previous works, the optimization of the thermal, mechanical and electrochemical properties of CPE based on P(VDF-TrFE) separators was performed as a function of filler type and content and it was established the ideal content for each filler in order to obtain high ionic conductivity with suitable thermal and mechanical properties [24e27]. With respect to molecular sieves and zeolites, it was found that 4 wt% of the layered clay montmorillonite, MMT, in P(VDF-TrFE) optimizes the electrochemical stability, thermal, mechanical and electrical properties [24]. In a similar way, the optimal filler content of NaY zeolite was 16 wt%, with which the CPE exhibited the highest ionic conductivity [25]. For MWCNT the electrolyte uptake, ionic conductivity and thermal stability of the polymer membranes proved to be independent of MWCNT type and the optimized filler content for this application was 0.1 wt% [27]. Finally, with the dielectric BaTiO3 as a filler, the physicochemical properties of the composite membranes are more dependent on filler size than on filler content, the most suitable CPE for battery separator applications being the composite with 16 wt% of BaTiO3 with an average size of 500 nm [26]. Generically, the aforementioned filler types allow increasing the ionic conductivity when compared to the pristine copolymer matrix but, depending on the filler type, other properties are also improved. Thus, clays also improved the thermal stability and electrolyte uptake and carbonaceous fillers improved the interfacial stability of the electrodes [33]. There are different CPE reported in the literature for lithium-ion battery applications [33], but there is a lack of direct comparison of different filler types for a given pristine polymer matrix and, in particular, this has been never performed for P(VDF-TrFE). In this context, this work shows a comparative study of the electrochemical performance and battery cyclability of CPE prepared with the optimized filler contents previous determined for P(VDF-TrFE) [24e27]: 4 wt% of MMT, 16 wt% of NaY, 16 wt% of BaTiO3 500 nm and 0.1 wt% of MWCNT, thus allowing a direct comparison of filler type on separator performance.

95

2. Experimental section 2.1. Materials Poly(vinylidene fluoride trifluoroethylene) (P(VDF-TrFE)) (70/ 30) was acquired from Solvay. Montmorillonite K10 (MMT) was acquired from Aldrich; NaY zeolites (Si/Al ratio: 2.83; Nominal Cation Form: Sodium; Na2O weight: 13%; Unit cell size: 24.65 Å) from Zeolyst International; barium titanate particles (BaTiO3) with average size of 500 nm were obtained from Nanoamor and multiwalled-carbon nanotubes (MWCNT) Baytubes C 150P from Bayer Materials Science. Lithium bis(trifluoromethanesulfonyl) imide (LiTFSI), obtained from SigmaeAldrich, was dried under vacuum at 25
C for 48 h and propylene carbonate (PC) was purchased from Merck and used as received. Pristine polymer, MMT/P(VDF-TrFE), BaTiO3/P(VDF-TrFE), NaY/P(VDF-TrFE) and MWCNT/P(VDF-TrFE) will be called hereafter P(VDF-TrFE), MMT, NaY, BaTiO3 and MWCNT, respectively. Table 1 shows the optimized filler content into P(VDF-TrFE) and the corresponding surface area, as provided in the datasheets of the materials. 2.2. Membrane preparation The different membranes were prepared following the procedures previously described for P(VDF-TrFE) composites with MMT [24], NaY [25], BaTiO3 [26] and MWCNT [27]. In short, each filler was dispersed in N,N-dimethylformamide (DMF, from Merck) during 4 h in ultrasonic bath. The polymer was then added to the solution to obtain a concentration of 15% (w/w). The filler particles to polymer relative concentrations was 4 wt% for MMT, 16 wt% for NaY, 16 wt% for BaTiO3 and 0.1 wt% for MWCNT, as determined in previous works to be the best filler concentrations for CPE applications [24e27]. The solution was prepared at room temperature under magnetic stirring until complete polymer dissolution was obtained (~2 h). In order to prevent the formation of aggregates and to accelerate polymer dissolution, the solution temperature was increased 5
C above room temperature during the first 15 min. Finally, the solution was placed in a Petri glass dish and the solvent completely evaporated at room temperature in a gas extraction chamber for 15 days. Fig. 1 shows a schematic representation of the procedure, indicating processing time and temperature for each preparation step of the P(VDF-TrFE) composite membranes. 2.3. Electrolyte solution and uptake The electrolyte solution uptake was performed by immersing the membranes for 24 h in a 1 M LiTFSI in PC electrolyte solution, with an ionic conductivity (s0) of 6.5 mS/cm at 25
C. The uptake value was evaluated according to:

Uptake ¼

m  m0  100 m0

(1)

Table 1 Optimized filler content of the different P(VDF-TrFE) composites for CPE applications and corresponding surface area for each filler [34e37]. Filler

Optimized filler content (wt%)

Surface area (m2/g)

MMT NaY BaTiO3 MWCNT

4 16 16 0.1

220e270 900 2.0e2.2 220

96

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Fig. 1. Schematic representation of the experimental procedure for the preparation of the P(VDF-TrFE) composite membranes.

where m0 is the mass of the dry membrane and m is the mass of the membrane after immersion in the electrolyte solution. 2.4. Characterization techniques Membranes were coated with a thin gold layer using a sputter coating (Polaron, model SC502 sputter coater) and the morphology analyzed using a scanning electron microscope (SEM) (Leica Cambridge apparatus at room temperature). The average pore size (f) was determined on the basis of the diameter of 40 pores, using the SEM images at 200 magnification and the Image J software. The porosity of the membranes (3) was determined with a pycnometer by the following procedure: the pycnometer was filled with ethanol and its mass was measured (m1); the mass of the sample was measured (m0) and then immersed in ethanol; after the sample being completely soaked in ethanol, ethanol was added to fill completely the pycnometer and the mass of the sample plus the pycnometer was measured (m2); finally, the sample was removed from the pycnometer and the residual weight of the pycnometer with ethanol was determined (m3). Ethanol was used due to its low density and easy soaking by the sample. The porosity of the membrane was calculated according to Ref. [38]:

3¼

m2  m3  m0 m1  m3

(2)

The porosity of each membrane was obtained as the average of the values determined in six measurements. 2.5. Electrochemical evaluation The ionic conductivity was evaluated by an Autolab PGSTAT-12 (Eco Chemie) set-up in a frequency range from 500 mHz to 65 kHz. The samples were placed in a constant volume support equipped with gold blocking electrodes located within a Buchi TO 50 oven. The temperature, measured by a type K thermocouple, was varied between 20 and 120
C. The ionic conductivity (si) was calculated for each heating cycle according to:

si ¼

d Rb A

rffiffiffiffiffiffiffiffi s0 3

si



si ¼ s0 exp

Ea RT

 (5)

where s0, is the pre-exponential factor, Ea is the apparent activation energy for ion transport, R is the gas constant (8.314 J/mol K) and T is the temperature.

2.6. Lithium cell manufacturing and testing The cathode was prepared from a slurry based on CeLiFePO4 (Phostech Lithium, Lda), carbon black (Super P, Timcal Graphite & Carbon, Switzerland) and PVDF (Binder, Solef 5130, Solvay) in a weight of 8:1:1 in N-methyl-1-pyrrolidone (NMP) solvent. After stirring, the slurry was casted on an aluminum foil through the doctor-blade technique and dried at 100
C for 4 h in a conventional oven (Binder ED23 oven). The active mass loading of the cathode material was ~2.5 mg cm2. The separator membranes were immersed in the electrolyte solution for 10 min within an argon filled glove box (JACOMEX, Germany) with the moisture and oxygen levels at ~1.0 ppm. Then, using the swollen membranes as separators with a diameter of 14 mm, 2032 coin-type Li/CeLiFePO4 half-cells were prepared. The diameter of the lithium anode and the CeLiFePO4 cathode discs was 10 mm. Galvanostatic cycling was carried out at room temperature in the voltage range from 2.0 V to 4.0 V at a current density between 0.1C and 2C using a Biologic VMP3 station. The activation cycle was carried out at 0.1C (17 mA g1).

3. Results and discussion 3.1. Microstructural characteristics, thermal and mechanical properties

(3)

where Rb is the bulk resistance, d is the thickness and A is the area of the sample. The tortuosity (t), ratio between the effective capillarity and the thickness of the sample, was determined by:

t¼

where s0 is the electrical conductivity of the liquid electrolyte, si is the electrical conductivity of the membrane and the electrolyte at room temperature and 3 is the porosity of the membrane. The temperature dependence of the ionic conductivity follows the Arrhenius equation:

(4)

It has been previously shown that P(VDF-TrFE) membranes prepared under suitable solvent evaporation conditions lead to a pore distribution adequate for lithium ion battery separator applications [22]. The characteristic porous microstructure has been explained by the spinodal decomposition of the liquideliquid phase separation followed by polymer crystallization [38,39]. Fig. 2 shows the SEM cross section images of the pristine polymer (Fig. 2a) and the composites membranes (MMT, Fig. 2b; NaY, Fig. 2c; BaTiO3, Fig. 2d; MWCNT, Fig. 2e).

J. Nunes-Pereira et al. / Composites Part B 96 (2016) 94e102

97

Fig. 2. Cross section pictures of a) pristine P(VDF-TrFE), b) MMT, c) NaY, d) BaTiO3 and e) MWCNT composite membranes.

The pore distribution observed in the composite membranes (Fig. 2bee) is similar to that observed in the pristine polymer (Fig. 2a). Thus, the pore distribution is not affected by the presence of fillers, leading to a homogeneous microstructure, but influences the phase separation process and solvent evaporation kinetics of the polymer/solvent system as reflected in the different pore size and degree of porosity (Fig. 3). Fillers are incorporated into the polymer matrix structure, with the exception of zeolites that appear decorating the internal walls of the pores. Further, the polymer crystallizes in the polar b-phase both in the pristine polymer and the polymer composites [24e27]. Fig. 3 shows the average pore size, porosity and electrolyte solution uptake of the polymer membranes, showing that they are influenced by the inclusion of the fillers. The pore size (8 mm for the pristine polymer) increases slightly with the presence of MMT and

70

b)

a)

Porosity (ε) Uptake

400

60

Property (%)

Average pore size (μm)

80

NaY and shows a stronger increase with the addition of BaTiO3 and MWCNT to 43 and 60 mm, respectively. This is related with shape of the filler, the average pore size increasing more for spherical (BaTiO3) and cylindrical nanostructure (MWCNT). For aluminosilicate fillers (MMT and NaY) the average pore size shows a larger variation for the NaY three-dimensional structure (f ¼ 21 mm) than for MMT. Generally, the porous structure formation is related to the liquideliquid spinodal decomposition and solvent evaporation rate, which are affected by the presence of the fillers and their interaction with polymer and solvent. The degree of porosity is higher for the composite membranes than for the pristine polymer, reaching values up to 83%, with the exception of the NaY membrane, which shows a lower degree of porosity (~36%), attributed to the accumulation of the NaY particles on the pore walls. Further, the liquideliquid phase separation seems to be

50 40 30 20

300 200 100

10 0

P(VDF-TrFE) MMT

NaY

Samples

BaTiO3

MWCNT

0

P(VDF-TrFE) MMT

NaY

Samples

Fig. 3. a) Average pore size, b) porosity and uptake for the different composite samples [24e27].

BaTiO3

MWCNT

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affected by the higher surface area of the NaY particles (Table 1), the interaction of the filler with the polymer chains resulting in lower phase separation of the P(VDF-TrFE)/DMF system and lower porosity. For the lower surface area fillers (MMT, BaTiO3 and MWCNT e Table 1) the porosity remains nearly constant. The electrolyte solution uptake is also higher for the composite membranes than for the pristine polymer, with the exception of the BaTiO3 membrane which shows a decrease, due to reduced free space in the pore cavities. The higher uptake values for MMT and NaY filler composites is related with the degree of porosity and, in particular, with the large number of Lewis acid sites in the fillers, leading to a good affinity with the liquid electrolyte [40,41]. Thus, filler type, and the related membrane morphological variations, is more relevant than filler content for determining electrolyte uptake, since samples with the filler concentration of 16% for NaY and MMT show the highest and the lowest uptake value, ~440 and 343% respectively (Fig. 3). Polymer battery separators must show integrity at high temperature (>100
C) [17,42] and their mechanical properties have a significant effect on the battery manufacturing process, safe handling and stability [42]. The determination of the melting temperature and mechanical properties is thus essential to analyze these requirements, as well as for the evaluation of the degree of crystallinity which also affects separator conductivity. These parameters were evaluated in the DSC thermograms and in the stressestrain curves presented in previous works [24e27]. Table 2 compiles the degree of crystallinity the melting temperature and the mechanical properties (yield stress and elastic modulus) of the composite membranes investigated in this work. The degree of crystallinity decreases slightly within experimental error with the inclusion of NaY, remaining unchanged for MMT and BaTiO3 composites. It seems that defective crystallization of the polymer is induced at the interface with NaY [21], which is also the filler with the largest surface area (Table 1). Further, there is a slight increase of the degree of crystallinity for the MWCNT filler composite, which may be related to the nucleation effect ascribed to MWCNT and its lower concentration with respect to other fillers, the filler acting as a nucleation agent and not as a defect within the polymer structure [44]. The overall mechanical properties of the composite membranes show slight differences when compared to the pristine polymer, indicating that those properties are mainly determined by the porous microstructure of the membranes and not by the fillers. The composites with MMT filler show the highest values of yield stress and elastic modulus (3.7 and 81.4 MPa, respectively), thus becoming the filler with the larger mechanical reinforcement effect, for the used filler content. On the other hand, the composite membrane with BaTiO3 shows lower mechanical properties than the pristine polymer, which can be attributed to the lower surface area of these particles, acting as a defect rather than a mechanical reinforcement. 3.2. Ionic conductivity and battery performance The electrochemical properties of the battery separator are essential in determining its performance [17]. Impedance spectroscopy provides information of cell resistance over a wide

frequency range and it was performed after electrolyte solution uptake. The results are shown through the Nyquist plots (Fig. 4), imaginary impedance Z00 as function of real impedance Z0 . Fig. 4a shows the Nyquist plot in the measured frequency range and Fig. 4b shows a magnification of the effect of each filler in the high frequency range. The Nyquist plot is typically characterized by three distinct zones: a semicircle at high frequencies corresponding to the charge transfer process, a straight line at low frequencies, related to diffusion processes and the transition between both [45]. The Nyquist plots of Fig. 4a do not show the semicircle characteristic of the charge transfer process, but an improved charge transfer process due to the presence of fillers (insert of Fig. 4b), suggesting and enhancement in the number of free ions and/or their mobility. The representation of the ionic conductivity at high frequencies (Fig. 4b) shows the inclined straight-line typical of the blocking electrode capacitive behavior [45], which depends on the filler type present in the polymer membrane. The ionic conductivity of the composites shown in Fig. 4b increases, as a function of filler type, in the following way: MWCNT < BaTiO3 < NaY < MMT < P(VDF-TrFE). Table 3 shows the ionic conductivity (si), tortuosity (t) and activation energy (Ea), obtained from equations (3)e(5), respectively. The room temperature ionic conductivity (Table 3) of the composite membranes is in the same order of magnitude for the pure polymer and the composites, but the presence of the fillers improve the ionic conductivity of the pristine membrane due to filler/salt interactions [46]. In all composite membranes, the ionic conductivity value after 10 min of swelling reaches 104 S cm1, i.e., a suitable value for lithium-ion battery applications [47]. Regarding the different fillers, the ionic conductivity is slightly higher for BaTiO3 due to the high polarity of the ceramic filler. The tortuosity describes the average pore connectivity and can be determined from the ionic conductivity measurements [17]. The tortuosity of the composite membranes with NaY and BaTiO3 is lower than for the pristine polymer (Table 3), which indicates that these fillers improve the pore connectivity within the composite. On the other hand, the membrane with MMT shows a slightly higher tortuosity than the pristine polymer, indicating a decrease of the pore connectivity. Further, the composite membranes show lower activation energy than the pristine polymer (Table 3) which is associated with the increase of ionic mobility caused by the addition of the fillers. The NaY membrane shows the lowest value of 3.3 kJ/mol, and this markedly decrease is related to the stronger Lewis acid centers in the filler [25]. The temperature dependence of the ionic conductivity, calculated from equation (3), is shown in Fig. 5 for the different membranes. It is observed that the overall ionic conductivity is slightly higher in the composite membranes, suggesting that the fillers promote higher ionic conductivity through increasing ionic charge carriers as indicated by the larger uptake value. The BaTiO3 filler, on the other hand, do not increase the uptake, but it is a highly polar,

Table 2 Degree of crystallinity, melting temperatures and mechanical properties of P(VDF-TrFE) and corresponding composites [24e27]. Sample

cc ± 2 (%)

Tm ± 2 (
C)

Yield stress ± 0.1 (MPa)

Elastic modulus ± 0.1 (MPa)

Ref.

P(VDF-TrFE) 4% MMT 16% NaY 16% BaTiO3 0.1% MWCNT

28 27 23 27 31

145.6 145.8 146.3 146.2 147.1

2.1 3.7 1.9 1.8 1.7

40.0 81.4 76.0 34.1 69.2

[43] [24] [25] [26] [27]

J. Nunes-Pereira et al. / Composites Part B 96 (2016) 94e102

8000

200

a)

b)

150

P(VDF-TrFE) 4%MMT 16%NaY 16% BaTiO3

100

0.1%MWCNT

-Z'' (Ω)

-Z'' (Ω)

6000

4000 P(VDF-TrFE) 4%MMT 16%NaY 16% BaTiO3

2000

0.1%MWCNT

0

0

1000

2000

3000

4000

Z' (Ω)

99

50

0 20

30

40

Z' (Ω)

50

60

5000

Fig. 4. Nyquist plots at 25
C for the pristine polymer and the composites in the measured frequency range (a) and at high frequencies (b).

high dielectric constant filler promoting increase of the ionic conductivity. Furthermore, regardless of the sample, Fig. 5 shows a moderate increase of the electrical conductivity as a function of temperature, suggesting the increase in the number of charge carriers and/or of the ionic mobility, as generally expected in gel electrolytes [43,48]. The electrochemical stability window of the pristine polymer and the composite membranes was determined in the previous works [24e27], showing good anodic stability above

Table 3 Room temperature ionic conductivity (si), tortuosity (t) and activation energy (Ea) of the membranes filled with 1 M LiTFSI electrolyte solution; s0 ¼ 6.5 mS/cm at 25
C. Sample

si ± 5% (S/cm)

P(VDF-TrFE) 4% MMT 16% NaY 16% BaTiO3 0.1% MWCNT

3.2 3.6 3.9 4.5 3.3

    

104 104 104 104 104

t ± 5%

Ea ± 5% (kJ/mol)

37.8 38.9 24.4 34.1 38.7

9.6 7.7 3.3 6.4 9.1

Temperature (ºC) 90

75

Log σi (S/cm)

-3.0

60

45

30

15

-3.1

P(VDF-TrFE) 4%MMT 16%NaY 16aTiO3

-3.2

0.1%MWCNT

-3.3 -3.4 -3.5 3.1

3.2

3.3

3.4

-1

1000/T (K ) Fig. 5. Log si as a function of 1000/T for the pristine polymer and the composites after electrolyte uptake.

4.0 versus Li/Liþ, i.e., enough to be compatible with most of the common cathode materials used for lithium batteries [49]. In order to investigate the electrochemical performance of the composites membranes, a lithium metallic/composite membrane/ CeLiFePO4 battery (cathodic half-cell) was assembled and tested and the results are shown in Fig. 6. The cyclability of the composite membranes (Figs. 6b and 7) is evaluated in relation to the results for the pristine polymer. Fig. 6a shows selected voltage versus capacity profiles at different scan rates from 0.1C to 2C for the pristine polymer, demonstrating the typical flat charge/discharge potential plateau that reflects the reversible charge (lithium removal)discharge (lithium insertion) cycling behavior of the CeLiFePO4 cathode material. This flat plateau is observed at around ~3.4 V versus Li/Liþ until 2C that corresponds to the Fe2þ/Fe3þ redox reaction [50]. The first discharge capacity at C/10 is 144.4 mAh g1 and for 2C is 82.7 mAh g1. Fig. 6a also shows that the potential of the discharge plateau voltage decreases progressively as the scan rates increases, which is a result of polarization within the electrode material. This behavior is due to the slow charge transfer due to the electronic resistance of the electrode [51]. At 2C (Fig. 6b) the polarization effect is independent of filler type. The discharge capacity at 2C for the different composites membranes is: MMT: 103.1 mAh g1; NaY: 56.1 mAh g1; MWCNT: 82.3 mAh g1 and BaTiO3: 79.4.7 mAh g1; and for the pristine polymer P(VDF-TrFE): 83.5 mAh g1. The discharge capacity value is thus related with the ionic conductivity and uptake values of the different samples. The higher discharge value of the composite membrane with MMT when compared to the pristine polymer shows that MMT leads to faster and easier lithium ion transportation between the electrodes and lower interfacial resistance, indicating better compatibility and lower reactivity of the electrolyte with the lithium metal [52]. It is also observed that high surface area fillers, such as MMT and NaY, form a barrier layer at the electrode that effectively hinders electrodeelectrolyte reaction [53]. Fig. 7a compares the discharge capacities of the lithium-ion half-cells assembled with the different samples with C-rates increasing from 0.1 to 1.0 C every five cycles and ten cycles for 2C. Fig. 7a shows that the discharge capacity values of the cathodic half-cell decrease with increasing discharge current density, the half-cell with the MMT composite membrane showing always higher discharge capacity values than the pristine polymer at the several rates. The composite membranes with

100

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Fig. 6. Galvanostatic chargeedischarge curves for a) P(VDF-TrFE) at different scan rates and b) for all samples at 2C.

different filler types exhibit superior rate capability for C/10 and C/5 in comparison with the pristine polymer. The cycling performances of 0.1% MWCNT, 16% NaY and 16% BaTiO3 composite membranes at C/2, C, 2C are very similar to the pristine polymer. At the scan rate of 2C, the composite membrane with MMT holds 54% of the discharge capacity value in comparison to 0.1C, while the pristine polymer presents a value of 47%. Thus, the composite membrane with MMT shows larger discharge capacities over a wide range of discharge current densities, indicating that the clay microstructure (large surface area and high stability) positively influences the ionic transport and the reduction of the ohmic polarization [54]. The cycling performance of the cathodic half-cells with the different composite membranes and the pristine polymer up to 50 cycles are presented in Fig. 7b at 2C. Regarding the cycling performance at 2C (Fig. 7b), the cathodic half-cell containing the composite membranes with MMT demonstrates a higher discharge capacity value (85 mAh g1 after 50 cycles) compared to the other samples, which is attributed to the high ionic conductivity value and uptake. The other composite membranes also show stable cycling performance and good capacity retention. For this scan rate, regardless the composite membrane, the discharge efficiency approaches to 100% with a very stable specific capacity over 50 cycles, which is 50% of the theoretical specific capacity for composite membrane with MMT. The discharge capacity values versus cycle number were also measured at C5 for all samples (data not shown). The coulombic efficiency is about 100% for all samples during the 50 cycles at the discharge current rate of C5, indicating excellent reversibility and stability

upon chargeedischarge cycling. It is also observed that the discharge capacity value as a function of cycle number is stable for all composite membranes after 50 cycles, e.g. the composite membrane with MMT shows 142 mAh g1 against 129.4 mAh g1 for the pristine polymer. At this scan rate, the discharge capacity value decrease with increasing cycle number for the pristine polymer, which is attributed to the formation of the inert solid electrolyte interface (SEI) film in the interfacial layer between sample and electrodes [55]. Thus, this work allows the quantification of the effect of different fillers on the performance of lithium-ion battery membranes. Further, it is shown that the best composite membrane is the sample with 4% MMT, showing better cycling performance than the pristine polymer. 4. Conclusions P(VDF-TrFE) based composite membranes with different types of fillers have been processed by solvent casting. The addition of fillers is reflected on membranes morphological, thermal, mechanical and electrochemical properties in different ways, depending on filler shape and type. The average pore size increases by the addition of fillers, in particular for spherical (BaTiO3) and cylindrical (MWCNT) particles. The degree of porosity of the composite membranes increases slightly with the inclusion of MMT, BaTiO3 and MWCNT, reaching a maximum value of ~83% for the membrane with the lowest surface area filler, MMT. All composite membranes are thermally stable up to 100
C and the mechanical

Fig. 7. Rate performance from 0.1C to 2C of cathodic half-cells from the different composite membranes and the pristine polymer (a) and comparison of the cycling performance of the cathodic half-cells for 50 cycles at 2C (b).

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behavior is determined by the porous microstructure. The ionic conductivity is strongly improved by the addition of the fillers, reaching the maximum value of 4.5  104 S/cm for the BaTiO3 composite membrane. The cathodic half-cells with the composite membrane with 4% MMT deliver a capacity of 173 mAh g1 at 0.1C against 144 mAh g1 for the pristine polymer. Regardless of the scan rate, the composite membranes exhibit good rate and cycling capabilities after 50 cycles. Thus, all used fillers lead to specific improvements in the performance of the membrane, being MMT the best choice for battery applications due to the high ionic conductivity and better rate capability. Acknowledgments This work was supported by the Portuguese Foundation for Science and Technology (FCT) under grants SFRH/BD/66930/2009 (J.N.P.), SFRH/BD/90313/2012 (A.G.), SFRH/BPD/112547/2015 (C.M.C.). The authors thank Solvay, Timcal and Phostech for kindly supplying the high quality materials. The authors thank financial support from the Basque Government Industry Department under n de Bizkaia for the ELKARTEK Program. SLM thanks the Diputacio financial support under the Bizkaia Talent program. References [1] Tarascon JM, Armand M. Issues and challenges facing rechargeable lithium batteries. Nature 2001;414(6861):359e67. [2] Armand M, Tarascon JM. Building better batteries. Nature 2008;451(7179): 652e7. [3] Dahlin GR, Strøm KE. Lithium batteries: research, technology, and applications. Nova Science Publishers; 2010. [4] Lu L, Han X, Li J, Hua J, Ouyang M. A review on the key issues for lithium-ion battery management in electric vehicles. J Power Sources 2013;226:272e88. [5] Scrosati B, Garche J. Lithium batteries: status, prospects and future. J Power Sources 2010;195(9):2419e30. [6] Tarascon J-M. Key challenges in future Li-battery research. 2010. [7] Wright PV. Developments in polymer electrolytes for lithium batteries. MRS Bull 2002;27(08):597e602. [8] Meyer WH. Polymer electrolytes for lithium-ion batteries. Adv Mater 1998;10(6):439e48. [9] Vijayakumar G, Karthick SN, Paramasivam R, Ilamaran C. Morphology and electrochemical properties of P(VdF-HFP)/MgO-based composite microporous polymer electrolytes for Li-ion polymer batteries. Polym-Plastics Technol Eng 2012;51(14):1427e31. [10] Agrawal RC, Pandey GP. Solid polymer electrolytes: materials designing and all-solid-state battery applications: an overview. J Phys D Appl Phys 2008;41(22):223001. [11] Gray FM, Chemistry RSO. Polymer electrolytes. Royal Society of Chemistry; 1997. [12] Di Noto V, Lavina S, Giffin GA, Negro E, Scrosati B. Polymer electrolytes: present, past and future. Electrochim Acta 2011;57(0):4e13. [13] Costa CM, Silva MM, Lanceros-Mendez S. Battery separators based on vinylidene fluoride (VDF) polymers and copolymers for lithium ion battery applications. RSC Adv 2013;3(29):11404e17. [14] Popall M, Andrei M, Kappel J, Kron J, Olma K, Olsowski B. ORMOCERs as inorganiceorganic electrolytes for new solid state lithium batteries and supercapacitors. Electrochim Acta 1998;43(10e11):1155e61. [15] Murata K. An overview of the research and development of solid polymer electrolyte batteries. Electrochim Acta 1995;40(13e14):2177e84. [16] Huang X. Separator technologies for lithium-ion batteries. J Solid State Electrochem 2011;15(4):649e62. [17] Arora P, Zhang Z. Battery separators. Chem Rev 2004;104(10):4419e62. k P, Müller K, Santhanam KSV, Haas O. Electrochemically active polymers [18] Nova for rechargeable batteries. Chem Rev 1997;97(1):207e82. [19] Judeinstein P, Titman J, Stamm M, Schmidt H. Investigation of ion-conducting ormolytes: structureeproperty relationships. Chem Mater 1994;6(2):127e34. [20] Gray FM. Solid polymer electrolytes: fundamentals and technological applications. Wiley; 1991. mez [21] Costa CM, Firmino Mendes S, Sencadas V, Ferreira A, Gregorio Jr R, Go Ribelles JL, et al. Influence of processing parameters on the polymer phase, microstructure and macroscopic properties of poly(vinilidene fluoride)/ Pb(Zr0.53Ti0.47)O3 composites. J Non-Cryst Solids 2010;356(41e42):2127e33. ndez S. [22] Costa CM, Rodrigues LC, Sencadas V, Silva MM, Rocha JG, Lanceros-Me Effect of degree of porosity on the properties of poly(vinylidene fluorideetrifluorethylene) for Li-ion battery separators. J Membr Sci 2012;407e408: 193e201.

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