In this study, a polypropylene (PP) mesh was used to prepare proton- and Li+ conducting composite membranes for fuel cells and lithium rechargeable batteries, respectively. For the preparation of Li+ conducting membrane, polypropylene mesh was first immersed in an electrolyte solution, which was composed of LiBF4 and ethylene carbonate. Then the swollen membrane was immersed in an acetone solution of polyethylene glycol diacrylate (PEGDA), polyvinylidenefluoride-co-hexafluoro-propylene and photoinitiator. Finally, PP fabric was taken out from the solution and exposed to UV irradiation. Furthermore, proton conducting membranes were prepared by immersing the PP mesh into a mixture of vinyl phosphonic acid, PEGDA and photoinitiator. Afterwards, samples were cured under UV light. PP-reinforced membranes designed for fuel cell applications exhibited a room temperature conductivity of 3.3×10-3 mS/cm, while UV-cured electrolyte for Li batteries showed ionic conductivities in the range of 1.61×10-3–5.4×10-3 S/cm with respect to temperature. In addition, for lithium-doped composite polymer electrolyte (CPE), the electrochemical stability window was negligible below 4.75 V vs. Li/Li+. It is concluded that lithium-doped CPE has suitable electrochemical stability to allow the use of high-voltage electrode couples.
During the last three decades, ion-conducting polymer electrolytes have continued to attract great attention from both scientific and industrial perspectives due to their growing application potentials in electronic devices such as rechargeable batteries, displays, sensors, fuel cells, etc. (1, 2). Work on polymer/salt mixtures has increased during recent years due to the growing interest in these systems as fast ion-conducting systems. Several types of lithium-ion conducting polymer electrolytes have been prepared as perfect homogenous mixtures of the components (3).
Polymer electrolyte membranes (PEMs) are promising advanced materials for fuel cells and Li polymer (LiP) battery applications as well as for super capacitors (4, 5). When compared to liquid electrolytes, solid polymer electrolytes have several advantages such as favorable mechanical properties, ease of fabrication into thin films of desirable size, and the ability to form effective electrode-electrolyte contacts (5). PEMs used in fuel cells are semipermeable membranes that serve as a solid barrier that separates fuel from reactant gases, while allowing the transportation of protons (6). PEMs for fuel cells should provide some requirements such as high proton conductivity, high water uptake above 100°C, durability, excellent thermal and mechanical properties, and low cost (7).
Lithium rechargeable batteries, which use lithium intercalation compounds as the positive and negative materials, consist of four main parts: anode, cathode, electrolyte and separator (8). The positive electrode (anode) is composed mainly of a metal oxide such as lithium cobalt oxide (LiCoO2), while the most widely used cathode material is graphite. Liquid electrolytes, gel electrolytes, polymer electrolytes and ceramic electrolytes are the four widely used electrolyte materials that conduct Li ions between cathode and anode materials (8, 9). For instance, gel electrolytes are generally composed of a polyvinylidenefluoride (PVDF) or polyvinylidenefluoride-co-hexafluoro-propylene (PVDF-HFP) polymer, a lithium salt (LiBF4, LiPF6, etc.) and a carbonate solvent. But their mechanical strength is still not efficient to allow high-speed battery manufacturing that would employ the lamination and packing processes used in the plastic industry. Thus, in the literature, there have been several recent reports where microporous polyolefin separator films were used also (10–13). A separator is a porous membrane whose main function is to separate anode and cathode materials to prevent electrical short circuits, and, thus, it plays a critical role in all battery designs (14). The separator that should be used in LiP should have good mechanical and dimensional stability, good insulating properties, uniform thickness and porosity, etc. (8, 9, 14).
The photopolymerization process is of great interest in industrial applications because of its superior features over conventional heat curing systems, such as lower energy consumption, less environmental pollution, lower process costs, high chemical stability and very rapid curing even at ambient temperatures (15–17). Because of these advantages, the photopolymerization process has been used widely in the coating and paint sectors (18). To our best knowledge, there are only a few studies that describe the preparation of both H+ and Li+ conducting membranes (19, 20). Panero et al. (19) and Nasef et al. (20) used electron beam for the grafting of styrene-swollen PVDF membranes. Then these membranes were either sulfonated via chlorosulfonic acid to prepare proton conducting membranes or immersed into an electrolyte solution to introduce Li-ion conductivity. However, the route they chose to prepare conductive membranes was somehow difficult due to the use of expensive electron beam equipment and to inevitable problems arising from the harsh conditions in the sulfonation process.
In this study, we also prepared both proton- and lithium-ion conducting membranes. We used a polypropylene mesh as a reinforcing material and utilized UV curing technology because of its ease of application, environmental friendliness and fast production rate. The goal of this study was to develop a fast, dynamic and easy way of preparing both proton-conducting and ionic-conducting membranes for fuel cells and lithium ion battery applications. This paper presents the preparation and characterization of UV-cured PVDF-HFP/polyethylene glycol diacrylate (PEGDA)/vinyl phosphonic acid (VPA) and LiBF4-doped PVDF-HFP/PEGDA composite polymer electrolytes reinforced by polypropylene fabric. The thermal properties of composites were evaluated by thermogravimetric analysis (TGA). The mechanical properties of proton- and ionic conducting membranes were analyzed by tensile measurements. The ionic and proton conductivities of membranes were measured as a function of temperature.
Polyvinylidenefluoride-co-hexafluoro-propylene (Mw 130,000 g/mol), PEGDA (Mw 258 g/mol) and VPA were purchased from Aldrich (St. Louis, USA) and used as received. The radical photoinitiator, 1-hydroxycyclohexyl phenyl ketone (Irgacure-184), was obtained from Ciba Specialty Chemicals (Basel, Switzerland). Acetone, ethylene carbonate (EC) and lithium tetrafluoroborate (LiBF4) were purchased from Merck (Germany). LiBF4 and EC were dried under vacuum before being used and then stored in a home-designed glove box until use. The polypropylene mesh (PP) was obtained from a local supplier (Turkey). Deionized water obtained from a Milli Q-water purification system was used.
Li+-ion conducting membranes were prepared by immersing the polypropylene fabric into the electrolyte solution of LiBF4 in EC. Strips cut from the porous polypropylene fabric were first immersed in an electrolyte solution (1 mol/l), which was composed of LiBF4 and EC, and kept in this solution for 3 h. In contrast, a 5 wt.% acetone solution of PVDF-HFP was prepared. Then, to this solution, PEGDA at a weight percent of 5% with respect to the amount of PVDF-HFP and the photoinitiator at a weight percent of 3% based on the total weight of PEGDA were added. And the mixture was stirred until a clear and stable solution was obtained. Swollen strips were then soaked into this PVDF-HFP/PEGDA solution for 1 min. Fabrics were taken out from the UV curable solution and immediately irradiated with a high-pressure UV lamp for 5 min. A homogeneous free-standing membrane was obtained. The thickness of the resulting membranes was found to be ∼580 μm.
Similarly to the preparation of the Li+ conducting membrane, strips of polypropylene fabric were immersed into a UV curable solution. The photocurable solution was prepared by mixing the VPA, PEGDA and photoinitiator. The weight ratio between PEGDA and VPA was 5:1, and the photoinitiator was added at 3 wt.% of the amount of VPA. Polypropylene strips were kept in this solution for 1 min, and the swollen membranes were then cured under UV irradiation for 5 min. The thickness of the resulting membranes was found to be about 0.5 cm.
FT-IR spectrum was recorded on a Perkin Elmer Spectrum 100 ATR-FTIR spectrophotometer (Perkin Elmer, Waltham, MA, USA). Scanning electron microscopy (SEM) imaging of the membranes was performed on a Philips XL30 ESEM-FEG/EDAX (Philips, Eindhoven, The Netherlands). The specimens were prepared for SEM by freeze fracturing in liquid nitrogen and applying a gold coating. The mechanical properties of the conductive membranes were determined by standard tensile stress-strain tests to measure modules, ultimate tensile strength and elongation at break. Standard tensile stress-strain experiments were performed at room temperature on a materials testing machine (Z010/TN2S, Zwick GmbH&Co. KG, Ulm, Germany), using a crosshead speed of 5 mm/min. The reported values were the average of at least three measurements. Thermogravimetric analyses of the UV-cured membranes were performed using a Perkin-Elmer Thermogravimetric analyzer (Pyris 1 TGA model, Perkin Elmer, Waltham, MA, USA). Samples were run from 30 to 750°C with a heating rate of 10°C/min under air atmosphere. Proton and Li+ conductivity measurements were recorded using a Gamry Potentiostat/Galvanostat/ZRA (Gamry Series G 750, Warminster, PA, USA) with the Gamry Framework Software (EIS300). Membrane samples were kept in distilled water for 24 h at room temperature prior to testing. Membranes were fixed between two platinum electrodes in a Teflon frame (BekkTech Conductivity Clamp), and proton conductivities were measured as a function of temperature. Conductivity measurement of fully hydrated membranes was carried out with the cell immersed in distilled water. The proton conductivity, σ, in these experiments was calculated from the following equation:
where L is the distance between the two electrodes; T and W are the thickness and width of the membrane, respectively; and R is the resistance value measured. The membrane resistances were obtained from Nyquist plots by extrapolating the impedance data to the real axis on the high frequency side. Lithium conductivity measurements were performed by sandwiching the polymer electrolyte between two stainless steel (SS-type 304, 0.025 mm thick, Alfa Aesar, Karlsruhe, Germany) electrodes. Measurements were carried out under argon atmosphere in a home-designed glove box (Ercom Kompresör, Kartal/Istanbul, Turkey). The electrodes were connected to a Gamry Potentiostat/Galvanostat/ZRA instrument equipped with the Gamry Framework Software (EIS300). The conductivity was calculated from this resistance, the thickness of the membrane and the area of the electrodes according to the equation below:
where σ is the ionic conductivity, Rb is the bulk resistance, λ is the film thickness and A is the surface area of the electrode. Linear sweep voltammetry (LSV) measurement was recorded using a Gamry Potentiostat/Galvanostat/ZRA with the Gamry Framework Software (PHE200, Physical Electrochemistry Software, Warminster, PA, USA). LSV experiment was performed to investigate the electrochemical stability window of the Li-doped composite polymer electrolyte using stainless steel (SS) as the working electrode and lithium foil (Sigma-Aldrich, Steinheim, Germany) as the reference and counter electrodes. Cell assembly was carried out in argon atmosphere inside a home-designed glove box. The scanning rate was 1 mV/s.
This work combines the properties of polypropylene mesh with the fast and easy production of the UV curing technology to fabricate novel conductive composite membranes for fuel cells and Li-ion batteries. The properties of the PP mesh used in this study are given in Table 1. Figure 1 shows the structures of all the monomers used and also depicts the preparation of conductive membranes.
|Sample||Type||Thickness (μm)||Pore size (μm)||Electrolyte uptakea (%)|
aOne mole per liter of LiBF4 in EC.
The structures of the composites were investigated by ATR-FTIR. The ATR-FTIR spectra of neat PP mesh (A), Li+ conducting membrane (B) and proton conducting membrane (C) can be seen in Figure 2. As can be seen in Figure 2A, in the neat PP mesh spectrum, a strong peak appears at about 2950–2800 cm-1 due to aliphatic –CH stretching vibrations.
In both spectra (B and C), the peaks at around 2920–2850 cm-1 are due to the aliphatic –CH stretching vibrations of PEGDA. Also, the ester carbonyl absorption of PEGDA can be seen at 1725 cm-1. In the spectrum (Figure 2B), the peak at 1070 cm-1 is related to the BF- anion in the polymer composite system (21). In the spectrum of Figure 2C, the peak at 1144 cm-1 was due to –P=O vibrations and the absorption bands at 1025 and 970 cm-1 were ascribed to the P–OH stretching bands. Previously, it was reported that cation-exchange resin containing the VPA group shows a strong peak at about 1194 and 965 cm-1 (22). The absence of the peaks corresponding to the vinylic and acrylic groups at around 1635–1610 cm-1 indicates that membranes were successfully prepared with high conversions of PEGDA and VPA.
The morphology of the composites was determined by scanning electron spectroscopy (SEM), as shown in Figure 3. As can be seen in Figure 3A, typical SEM images of neat PP mesh fragments revealed a smooth and very regular surface. The figure also shows the pore size and interstice distance, which are important mesh characteristics.
A cross-sectional SEM image of the proton conductive composite is shown in Figure 3B. The adhesion between the PP mesh filament and the polymer matrix is crucial for the mechanical properties and the thermal stability of the PP fabric-reinforced composites. Thus, the PP fabric filament should be well penetrated into the polymer matrix or should be surrounded by the polymer matrix. As can be seen in Figure 3B, it is evident from the SEM images that neat PP mesh and PP mesh coated and well surrounded by the polymer matrix were obtained.
The TGA technique was used to investigate the thermal oxidative stability of these polypropylene fabric-reinforced conductive composite materials. TGA thermograms of the membranes are given in Figure 4, and the results are shown in Table 2. In this work, proton conducting membranes were prepared by using VPA. In Table 2, it can be seen that the first weight loss temperature (5 wt.%) of neat polypropylene fabric is 272°C, while it is 214°C and 217°C for the proton conducting membrane and lithium-ion conducting membrane, respectively. It can be seen from the results that conductive membranes start to degrade before the neat PP fabric. This situation can be attributed to the loss of absorbed moisture, unreacted monomers and low-molecular-weight volatile compounds (photoinitiator, solvent, etc.). The char yields were also collected at 750°C. It can be clearly seen that 16 wt.% of the proton conducting membrane remains at 750°C, while neat PP fabric and lithium-ion conducting membrane completely degraded before 600°C. This result shows that the proton conducting membrane is thermally stable and suitable for high-temperature fuel cell applications.
|Sample||Tensile modulus (MPa)||Tensile strength (N)||Elongation at break (%)||T5a (°C)||Tmax (°C)||Char (%)|
|Proton conducting membrane||185±15||89±7||6±2||214||375||16|
|Li-ion conducting membrane||26±7||91±11||138±34||207||350||–|
aT5 is the 5% weight-loss temperature.
The tensile strength, elongation at break and Young’s modulus of the composite membranes are shown in Table 2. Untreated PP fabric was found to have low modulus and a high elongation-at-break value. When samples were cured under UV light, the modulus of the conductive membranes increased due to the cross-linking of PEGDA. Furthermore, it can be seen that the modulus of the ionic conducting membrane is lower than that of the proton conducting membrane, which can be attributed to the presence of PVDF-HFP in lithium conducting membranes. The presence of PVDF-HFP decreased the UV curing efficiency. Also, due to their flexible structure, these membranes resulted in low modulus and high elongation-at-break values. The decrease in the tensile modulus of the composite polymer electrolytes by the incorporation of the Li salt is attributed to the intramolecular interaction between the chains of the polymer and the salt. Meanwhile, the flexibility of the membrane improved due to the plasticizer effect of EC. The plasticizing effect is related to the weakening of the dipole-dipole interaction between the PVDF-HFP chains. This weakening of the dipole-dipole interaction increases the flexibility of the composite polymer electrolyte (23). Thus, the elongation at break (%) value increased. In contrast, the mechanical strength of this Li-ion conducting membrane is comparable to that found in previous studies (24, 25). Moreover, VPA and PEGDA containing proton conducting membranes showed high modulus and, due to the restricted movement of the molecular chains, these proton conducting membranes had low elongation-at-break values.
The plot of H+ conductivity values vs. temperature for VPA containing PP-reinforced membranes is shown in Figure 5. PP-based proton conducting membranes designed for fuel cell applications exhibited a room temperature conductivity of 3.3×10-3 mS/cm at 100% humidification. As can be seen from the figure, proton conductivity increases with respect to temperature and reaches the value of approximately 1.0×10-2 mS/cm at 80°C.
A plot of the temperature dependence of lithium ion conductivity of PP-based composite membranes is shown in Figure 6. It can be seen from the plot that PVDF-HFP containing UV-cured membrane for Li-ion batteries showed ionic conductivities in the range of 1.61×10-3–5.4×10-3 S/cm with respect to temperature (20–80°C). Song et al. (26) found the ionic conductivity of LiPF6 containing a PEGDA/PVDF blend (5:5) to be 4 mS/cm at room temperature. Thus it can be said that the conductivity of our membranes is in good agreement with literature results.
An Arrhenius plot of the proton conducting membrane is given in Figure 7. The Arrhenius plot is related to the activation energy (Ea). Thus, the Ea for polymer electrolytes can be calculated by the Arrhenius equation:
where T is the temperature on the Kelvin scale, σ0 is a pre-exponential factor and R is the ideal gas constant. The Ea value calculated from the slope of the Arrhenius plot of VPA containing proton conducting membrane was determined as 6.90 kJ/mol. This value is lower than the activation energy of Nafion 115 membrane (9.04 kJ/mol). Thus proton transfer is much easier in PP-reinforced membranes (27).
The Arrhenius plot for Li+ conducting membranes is shown in Figure 7. The activation energy for these membranes was determined as 7.57 kJ/mol. A linear relationship was observed between lithium ion conductivity and temperature values.
Figure 8 shows the current-voltage response of the lithium-doped composite polymer electrolyte obtained in the voltage range between 3.5 and 5.5 V vs. Li at room temperature. Appetecchi et al. (28) reported that the liquid phase decomposed at the surface of the lithium, thereby severely affecting the electrochemical performance, such as the cyclability of the lithium electrode. Thus, the high stability of the composite polymer electrolytes (CPEs) may be due to the absence of any impurities, which is a welcome feature because it permits their use in high-voltage battery applications (29). The flat plateu in Figure 8 with no peaks in the 3.5–4.7-V voltage range confirms the high purity of the prepared composite polymer electrolyte. The current increases rapidly with a small change in the potential and so this shows anodic breakdown. In addition, the onset of the current during anodic scan, which is representative of the decomposition of the membrane, indicates an anodic breakdown voltage of approximately 4.75 V vs. Li.
The aim of this study was to prepare both Li+ conducting and proton conducting membrane for LiPs and fuel cells, respectively. In this study, a PP mesh was chosen as a reinforcing material for its low price, good thermal and mechanical properties, and porous structure. Proton conducting membranes showed high thermal stability due to the presence of VPA. Due to the cross-linked structure of the resulting membranes, the mechanical properties of the PP fabric were enhanced. Proton conducting membranes designed for fuel cell applications exhibited a room temperature conductivity of 3.3×10-3 mS/cm. Moreover, Li-ion conducting membranes showed ionic conductivities in the range of 1.61×10-3–5.4×10-3 S/cm with respect to temperature (20–80°C). In addition, for lithium-doped CPE, the current responses are negligible below 4.75 V vs. Li/Li+. It is concluded that lithium-doped CPE has suitable electrochemical stability to allow the use of high-voltage electrode couples. It is evident from SEM images that neat PP mesh and PP mesh coated and well surrounded by polymer matrix were obtained.
Financial support from the Scientific and Technological Research Council of Turkey (research project number 209T121) is gratefully acknowledged.
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