Ind. Eng. Chem. Res. 2010, 49, 1281–1288
1281
Effect of Ball Milling on Electrochemical Properties of PVdF-HFP Porous Membranes Applied for DMFCs G. Gnana kumar,† Chang Su So,† A. R. Kim,† Kee Suk Nahm,*,†,‡ and R. Elizabeth Specialized Graduate School of Hydrogen and Fuel Cell Engineering and School of Chemical Engineering and Technology, Chonbuk National UniVersity, Jeonju 561-756, Republic of Korea, and Department of Physics, Lady Doak College, Madurai 625002, India
Poly(vinylidene fluoride-co-hexafluoropropylene) (PVdF-HFP) porous polymer electrolyte membranes were prepared by the solvent-nonsolvent evaporation technique. The particle size and molecular weight of the polymer were greatly varied with the ball-milling technique, and the electrochemical properties of the polymer electrolyte membranes were examined as functions of these variations. The high hydrophobic nature of the host polymer is responsible for the high-temperature application and low methanol permeation. A hydrophilic nature was imparted to the fabricated porous membranes through activation with silica sulfuric acid liquid electrolyte. A decrease in the molecular weight effectively reduces the viscosity and is responsible for high ion migration. In addition, a decrease in the particle size lead to an effective acidification, so that the acidified polymer could effectively promote the ionic conductivity further. A much lower fuel permeation associated with the high ionic conductivity of the ball-milled membranes facilitates the performance of direct methanol fuel cells and supports their potential application therein. 1. Introduction For future energy demands, the direct methanol fuel cell (DMFC) has been considered as a prevailing energy source that provides emission-free clean energy, low costs, and fuel safety.1,2 However, the low power density derived from DMFCs other than proton-exchange membrane fuel cells (PEMFCs) impedes their commercialization.3 The power density of DMFCs can be enhanced only by the proper polymer electrolyte membranes and the desired characteristics of polymer membranes are as follows: (i) high ionic conductivity, (ii) low fuel permeation, (iii) low thickness to minimize the membrane resistance, (iv) low swelling capability to minimize the interface resistance, (v) chemical stability, and (vi) low cost.4 So far, many efforts have been made to develop polymer electrolyte membranes that fulfill these requirements. However, most polymer electrolyte membranes exhibit a high ionic conductivity associated with high methanol permeability.5 An increase in the methanol permeability leads to poisoning of cathode catalysis, increased reaction overpotential due to the mixed potential, loss of fuel, and emissions of low-concentration toxic materials.6,7 In addition, a high concentration of fuel leads to a higher fuel permeability, which results in a lower fuel cell performance. These problems can be effectively tackled only through the membrane, not through the catalysts. Although catalysts can reduce the methanol permeability to a minimal extent, the host matrix of the membrane electrode assembly (i.e., the polymer electrolyte membrane) can effectively lower the methanol permeability and can enhance the overall efficiency of DMFCs. For effective ionic conductivity, hydrophilic particles are preferred, whereas for lower molecular transportation, hydrophobic particles are favored. In general, hydrophilic polymers cannot control the high molecular permeation that results in high fuel crossover.8 This directs research efforts toward hydrophobic * To whom correspondence should be addressed. E-mail: nahmks@ chonbuk.ac.kr. Fax: +82 63 270 2306. † Specialized Graduate School of Hydrogen and Fuel Cell Engineering, Chonbuk National University. ‡ School of Chemical Engineering and Technology, Chonbuk National University.
polymers such as poly(vinylidene fluoride-co-hexafluoropropylene) (PVdF-HFP) for the betterment of physical stability and lower molecular permeation. However, the highly hydrophobic nature of such polymers does not provide effective pathways for ion transportation. Ion migration is enhanced by the addition of hydrophilic acidic sites and is further enhanced by the inclusion of hygroscopic fillers. Recently, porous polymer electrolyte membranes have become attractive in the fuel cell field because of their limited ionic conductivity and other electrochemical properties. The electrochemical properties of porous membranes are widely dependent on three parameters, namely, (i) extended porosity for the high electrolyte absorption, (ii) pore interconnectivity for the efficient transportation of charge carriers, and (iii) pore homogeneity for low mass concentration gradients.9 Although the porous nature of such membranes constrains ionic conductivity, it also leads to high fuel permeability and low electrolyte retention. Efforts have been taken to address the aforementioned problems from the perspective of porogenic agents and suitable electrolytes.10-12,1 However, so far, no efforts have been made to modify the polymer host (PVdF-HFP) itself. If the modification is carried out on the host matrix instead of the aforementioned guest materials, effective results could be obtained. Hence, we report herein a study on a porous polymer electrolyte membrane and its influence on fuel cell performance. 2. Experimental Section 2.1. Membrane Preparation. Poly(vinylidene fluoride-cohexafluoropropylene) (PVdF-HFP) (Kynar Flex 2801, Atofina, Japan) copolymer was ball-milled for 5, 10, 15, and 20 h at 500 rpm using a Fritsch Pulverisette 6 planetary mono mill ballmilling machine. For comparison, the derived polymer was also utilized without any modification. The appropriate amount of pure or ball-milled PVdF-HFP copolymer was dissolved in a volatile solvent (acetone) and stirred for 1 h at 60 °C. The immiscible nonsolvent tetrahydrofuran (THF) was added, and the resulting solution was magnetically stirred for 3 h. The concentration of PVdF-HFP
10.1021/ie901008k 2010 American Chemical Society Published on Web 12/18/2009
1282
Ind. Eng. Chem. Res., Vol. 49, No. 3, 2010
Table 1. Nomenclature, Porosity, Methanol Permeability, and Selectivity Ratio of the Fabricated Porous Membranes methanol permeability (cm2/s) porosity (%) sample pure PVdF-HFP/THF ball-milled (20 h) PVdF-HFP/THF
nomenclature PHP BMPHP
before activation 63.45 55.64
after activation 13.45 4.53
in acetone solution was maintained at 10 wt %, and that of the nonsolvent was maintained at 5 wt %. The resultant solution was cast as a membrane on a glass substrate. The prepared membranes were kept at 120 °C under 10-3 Torr pressure for 36 h in a vacuum for the evaporation of the solvent and nonsolvent, and the membranes were activated with a 10 wt % solution of silica sulfuric acid and dried. The properties of the fabricated membranes are presented in Table 1. 2.2. Preparation of Silica Sulfuric Acid. Silica particles (90 nm in size) were prepared according to the procedure described elsewhere.13 A 500 mL suction flask equipped with a constantpressure dropping funnel and a gas outlet tube were utilized for the preparation of silica sulfuric acid. The suction flask was charged with 2.5 g of silica gel, and 11 g of 1 M chlorosulfuric acid was added dropwise over a period of 30 min at room temperature. HCl gas was immediately evolved from the reaction vessel as follows SiO2-OH + CISO3H f SiO2-OSO3H + HCIv
methanol concentration
(1)
After the completion of the reaction, the mixture was shaken for 30 min and subjected to further experiments. 2.3. Morphological Characterizations. Transmission electron microscopy (TEM) images of the polymer were obtained using a BIO-TEM instrument (H-7650, Hitachi, Tokyo, Japan). Morphological images of the pure and ball-milled PVdF-HFP polymer were obtained with an XE-150 advanced scanning probe microscope (Park Systems Corp., Suwon, Korea) equipped with a Z scanner in noncontact mode operation at a scan rate of 0.6 Hz. Molecular weights of the studied polymers were obtained by gel permeation chromatography using an HLC-8320 EcoSEC system (Tosoh Corporation, Tokyo, Japan). Morphological images of the prepared membranes were obtained by scanning electron microscopy (SEM) using a JSM5410LV instrument (JEOL, Tokyo, Japan) equipped with energy-dispersive X-ray (EDX) spectroscopy. 2.4. Porosity, Electrolyte Uptake, and Leakage. The porosities of the polymer electrolyte membranes were determined by n-butanol absorption as described elsewhere.14 The membranes were activated with a 10 wt % solution of silica sulfuric acid for a couple of hours for the determination of electrolyte uptake. After the excess solution at the surface of the polymer electrolyte had been removed, the membranes were dried and weighed. Liquid uptake and electrolyte leakage were calculated using the equations liquid uptake (%) ) [(Wf - Wd)/Wd] × 100 electrolyte leakage (%) ) [(Wi - Wf)/(Wi - Wd)] × 100 where Wd is the weight of the dried polymer membrane and Wi and Wf are the initial and final (equilibrium) weights, respectively, of the membrane after absorption of the acid aqueous solution. 2.5. Structural and Physical Characterization. The FTIR spectra of the membranes were recorded between 400 and 4000 cm-1 in transmittance mode using a Jasco (Tokyo, Japan)
conductivity (mS/cm) 34.65 56.72
2M
4M -7
1.2 × 10 9.5 × 10-8
6M -7
3.5 × 10 1.8 × 10-7
4.9 × 10-6 2.8 × 10-6
FT-IR-300 spectrophotometer. The thermal behavior of the membranes was examined on a Perkin-Elmer (Wellesley, MA) instrument under a nitrogen atmosphere at a heating rate of 20 °C/min from 30 to 800 °C. 2.6. Electrochemical Characterization. The ionic conductivity of the fabricated membranes was examined using a BekkTech (Loveland, CO) conductivity test cell in conjunction with a PGZ 301 Dynamic electrochemical impedance spectroscopy (EIS) voltammeter (Radiometer Analytical SAS, Lyon, France).15,16 Methanol permeation rates were measured using a house permeation cell at 30 °C as described elsewhere.16 Direct methanol fuel cell performance was tested on a Scribner (Southern Pines, NC) 850C compact fuel cell test station. Membrane electrode assemblies (MEA) were fabricated by using 1:1 Pt-Ru (1.5 mg/cm2) and Pt (0.5 mg/cm2) as the anode (methanol oxidation) and cathode (oxygen/air reduction) catalysts, respectively. A mixture of distilled water, Nafion ionomer, and isopropanol was sprayed onto the commercially available electrodes by the spray method. The impregnated electrodes were dried at 60 °C for 30 min. Then, the prepared composite membrane was sandwiched between the electrodes and hot pressed at 100 °C with a pressure of 100 kg/cm2 for 3 min. The DMFC electrochemical performances of the fabricated membrane electrode assemblies were evaluated with a singlecell fixture having an active area of 5 cm2 (ElectroChem Inc., Woburn, MA). Methanol solution (2, 4, or 6 M) was fed into the anode at a flow rate of 4 mL min-1 by a peristaltic pump, and nonhumidified air was fed into the cathode at a flow rate of 200 mL min-1 with a back-pressure of 20 psi. The current-voltage (I-V) data were collected at 80 °C by measuring cell voltages that were stabilized for at least 3 min after application of current in potentiostatic mode. 3. Results and Discussion 3.1. Morphological Properties of the Polymer. The surfaceto-bulk ratio of nanosized materials is much greater than that of coarse materials. The large fraction of atoms present at the surface of a nanomaterial dominates its surface properties. Indeed, grain-size reduction is one of the main factors for the enhancement of electrochemical properties. Ball milling is an effective technique for reducing the particle size from micrometers to nanometers. The shaking of steel balls imparts kinetic energy to the individual balls within the walls of the vial. The particle size gets reduced by the kinetic energy imparted through collisions among the balls. As the milling time increases, the particle size decreases to the point that there are lattice defects in the molecular structure of the powder. Figure 1 shows TEM images of pure PVdF-HFP and 20-h-ball-milled polymer powders. The size of pure PVdF-HFP polymer particles ranges between 4 and 6 µm, whereas the ball-milling technique effectively reduces the size of the polymer particles from micrometers to nanometers (i.e, in the range of 190-350 nm). The micro- and nanostructures of polymer are fundamental to a number of modern technological applications including polymer electrolyte membranes which can greatly affect the physical and electrochemical behaviors. Figure 2 exhibits the
Ind. Eng. Chem. Res., Vol. 49, No. 3, 2010
Figure 1. TEM images of (a) pure and (b) ball-milled PVdF-HFP polymer powder.
AFM images of PVdF-HFP copolymer and 20-h-ball-milled copolymer. Figure 2a clearly reveals the bulky size of the obtained polymer. Meanwhile, the ball-milled polymer contains nanometric-sized particles, as shown in Figure 2b. Effective collisions of the balls disperses the bulkier particles and results in homogeneously distributed nanosized particles (Figure 2b.i inset). The molecular weights of 0-, 5-, 10-, 15-, and 20-h-ball-milled polymers were determined by gel permeation chromatography to be 401100, 396138, 226199, 171166, and 165311 g/mol, respectively. Strong collisions of the balls with the obtained polymer effectively reduced its molecular weight, and the effect of molecular weight on the electrochemical properties is discussed in a later section. 3.2. Morphological Images of the Membranes. Figure 3 reveals the SEM images of PHP and BMPHP membranes (see Table 1 for membrane nomenclature). The morphology of the polymer electrolyte membrane plays a decisive role in electrolyte absorption through the porous matrix and, therefore, in the electrochemical properties. The outward expansion of polymer segments gives a porous nature to the membrane upon the complete evaporation of the nonsolvent THF (Figure 3). The highly uniform pores developed in the fabricated membranes are responsible for the high electrolyte absorption and high ion transport. Similar pores were observed for the studied PHP and BMPHP membranes. In addition, a thin layer formed over the porosity for the BMPHP membrane, and the exact mechanism involved in its formation is yet to be understood. 3.3. Porosity. Although a high porosity is responsible for the high ionic conductivity of polymer electrolyte membranes, it results in high fuel permeability, low electrolyte retention, physical instability, and high swelling.17 Thus, an imbalance arises between the required parameters of polymer electrolyte membranes. Therefore, a balance must be maintained, which can be done by optimizing the porosity. Table 1 lists the porosities of the fabricated membranes. A decrement in the porosity was observed for the BMPHP membrane compared to the PHP membrane. Although a decrease in porosity was observed for the BMPHP membrane, this membrane exhibited higher electrochemical properties than the PHP membrane, as discussed later. The decrease in the porosity of the BMPHP membrane is attributed to the thin layer that formed over its porosity. 3.4. Electrolyte Absorption. Electrolyte absorption is an essential factor for porous membranes that ensures ionic conductivity and functionalization.18 The porous PVdF-HFP membrane act as a host matrix, and electrolyte absorption effectively creates functionalization. For the prepared porous membranes, interconnected pores facilitate high liquid electrolyte absorption and high ion transportation.19 Figure 4 illustrates the electrolyte absorption of the fabricated membranes as a function
1283
of time. The electrolyte uptake increases gradually with time and becomes constant after 18 h. Although the BMPHP membrane exhibited a lower porosity than the PHP membrane, it exhibited a high liquid electrolyte absorption because of the more strongly interconnected pore structures obtained through the ball-milling effect. From this result, it is clear that porosity alone cannot determine the effectual electrolyte absorption; influential factors such as viscosity and particle size also play crucial roles in electrolyte absorption. The compatibility of the electrolyte and lower tortuosity in the prepared porous membranes influences their high electrolyte uptakes. 3.5. Electrolyte Leakage. For porous polymer electrolyte membranes, electrolyte leakage is a significant threat that degrades the durability of the membranes.1,20 In general, an increase in the porosity favors a high electrolyte uptake and also leads to a high electrolyte leakage. Even though a high electrolyte uptake was obtained, a minimal electrolyte leakage maintained for the fabricated PVdF-HFP membranes (Figure 4) influenced its lifetime durability. The minimal electrolyte leakage obtained for the fabricated membranes is attributed to their interconnected pore structures. 3.6. Thermal Analysis. Fuel cells have to be operated at higher temperatures. An increase in the fuel cell operating temperature eliminates the use of a reformer and greatly enhances the electrode reaction kinetics; improves the CO tolerance, which promotes a light weight; and contributes to efficient green-power-generating fuel cells.21-23 In general, the hydrophobic nature of the host polymer satisfies the thermal properties, whereas the inclusion of a hydrophilic nature enhances the electrochemical properties and also deteriorates the corresponding membrane’s thermal stability. In a few cases, the host polymer itself is hydrophilic in nature and cannot effectively satisfy the higher-temperature applications because of its lower thermal stability. The above criterion influenced the selection of hydrophobic PVdF-HFP as the host matrix for this study. As depicted in the TGA spectrum of Figure 5, the BMPHP membrane before electrolyte activation experienced a single weight loss at around 420-500 °C due to the thermal degradation of PVdF-HFP units. However, the silica sulfuric acid activated membrane exhibited two major weight losses at around 330 and 420-500 °C, which are attributed to the degradation of sulfonic acid units and the random scission of PVdF-HFP units, respectively.1 From the thermal analysis, it is clear that silica sulfuric acid was entrapped in the PVdF-HFP polymer matrix. In addition, a higher thermal degradation of sulfonic acid units observed for the BMPHP membrane ensured the higher electrolyte uptake of the membrane as evidenced from electrolyte uptake values. Thermal analysis of the nonactivated pure and ball-milled powder membrane is shown in the inset of Figure 5. The ball-milled powder membrane experienced a weight loss at lower temperatures than the pure PVdF membrane, which is attributed to the smaller polymer particle size. A decrease in the polymer particle size increases the relative surface area of the membrane. An increase in the surface area of the samples increases the heat-transfer area, which, in turn, increases the heat-transfer rate. The increased heat-transfer rate allows the individual particles to reach the furnace temperature more rapidly than larger particles. From this explanation, it is clear that the particle size of the polymer plays a decisive role in the membrane, even though the polymer is completely dissolved in the solvent. Figure 5 shows the DSC profile of the BMPHP polymer electrolyte membrane before and after activation with the silica
1284
Ind. Eng. Chem. Res., Vol. 49, No. 3, 2010
Figure 2. AFM images of (a) pure and (b) ball-milled PVdF-HFP polymer powder Inset: Highly magnified image of ball-milled PVdF-HFP polymer powder.
Figure 3. Morphological images of (a) PHP and (b) BMPHP polymer electrolyte membranes.
Figure 5. Thermal gravimetric analysis (TGA) and differential scanning calorimetry (DSC) spectra of BMPHP polymer electrolyte membrane (s) before and ( · · · ) after activation with silica sulfuric acid.
Figure 4. Electrolyte absorption and electrolyte leakage for the (9) BMPHP and (b) PHP polymer electrolyte membranes.
sulfuric acid electrolyte solution. The melting temperature of the activated BMPHP membrane is lower than that of the nonactivated BMPHP membrane. The bimodal peak that is present for the nonactivated membrane vanishes for the activated membrane and results in a lower melting temperature. The entrapment of the silica sulfuric acid liquid electrolyte ruptures the crystal nature of the host polymer, which leads to a partial distortion of the crystalline structure of the PVdF-HFP membrane. A strong interaction between the hydrophobic polymer chains and the hydrophilic liquid electrolyte solution reduces the crystallinity, which, in turn, increases the capability to absorb a significant amount of water molecules. The influence of the
Figure 6. IR spectra of BMPHP polymer electrolyte membranes (a) before and (b) after activation with silica sulfuric acid.
ceramic property of the silica sulfuric acid electrolyte solution impedes the polymer electrolyte membrane’s additional melting temperature. 3.7. FT-IR Spectroscopy. Figure 6 provides structural confirmation of ball-milled PVdF-HFP and silica sulfuric acid entrapped ball-milled PVdF-HFP membranes. The IR spectrum of the bare BMPHP membrane (before electrolyte activation) is assigned as follows (Figure 6a): The deformed vibration of
Ind. Eng. Chem. Res., Vol. 49, No. 3, 2010
Figure 7. Ionic conductivity in terms of temperatures measured at 100 RH for the (9) PVdF-HFP (20-h-ball-milled)/THF, (b) PVdF-HFP (15-h-ballmilled)/THF, (2) PVdF-HFP (10-h-ball-milled)/THF, (1) PVdF-HFP (5h-ball-milled)/THF, and (() PVdF-HFP (not ball-milled)/THF polymer electrolyte membranes activated with silica sulfuric acid.
CH2 groups appears at the frequency of 1403 cm-1, with the weakening interaction between the H atoms of the CH2 groups and the F atoms of the CF2 groups. The asymmetric and symmetric stretching vibrations of CF2 are found at 1085 and 1190 cm-1, respectively. The band observed at 883 cm-1 ensures the amorphous phase of the copolymer. The C-F wagging mode is assigned to the peak observed at around 613 cm-1. The R phase of PVdF-HFP is found at 766 and 976 cm-1, whereas the β and γ phases are assigned to the bands observed at 489 and 842 cm-1.24,25,1 For the silica sulfuric acid entrapped BMPHP membrane, the observed peaks are attributed as follows: The hydronium ion (H3O+) formation in the membrane is found at 3416 cm-1. This peak should not be of water itself, but is attributed to hydronium ion, which can combine with coexisting RSO3- to produce RSO3H and release water with the progress of dehydration in the membrane.15 The absorption of water molecules is observed at 1640-1650 cm-1. A strong peak observed at 1101 cm-1 is due to the sulfonated silica particles. S-O bands of silica sulfuric acid are found at 610 and 572 cm-1.26 The peak at 476 cm-1 is assigned to the Si-O-Si and -O-Si-O stretches.27,28 From this IR spectrum, it is clear that silica sulfuric acid electrolyte has been completely entrapped in the porous PVdF-HFP polymer matrix and confirms the structural configuration of the composite membrane. Thus, a strong interaction occurred between the host polymeric units and the silica sulfuric acid that guaranteed the durability of the fabricated porous membranes. 3.8. Ionic Conductivity. The ionic conductivity of the fabricated membranes is shown in Figure 7. In general, for porous membranes, three phases are found: (i) host polymer matrix, (ii) liquid electrolyte solution stored in the pores, and (iii) gel type of electrolyte formed by the swelling of the PVdFHFP membrane by liquid electrolyte.29 For nonporous membranes, the transport of the liquid electrolyte from the surface layers to the core of the film is tedious. However, for porous membranes, liquid electrolyte easily swells the surface layer; enters the pores, which were produced for high electrolyte absorption; and swells the polymer chains. For the pore-free PVdF-HFP membrane, the chance of liquid electrolyte entrapment is ruled out, which is responsible for the lower ionic conductivity. The pores created from the evaporation of the
1285
nonsolvent THF for the fabricated membranes effectively creates pathways for high electrolyte absorption. The ionic conductivity achieved in this study, which is higher than that of a previous report,1 is attributed to the better compatibility between the porous membrane and the silica sulfuric acid electrolyte than other conventional liquid electrolytes. In addition, silica sulfuric acid has both acidic and hygroscopic functionalities, which effectively humidifies the membrane and ensure high-temperature operation. As discussed, an increase in the ball-milling time effectively reduces the molecular weight of the corresponding polymers. In general, high-molecular-weight polymers exhibit higher viscosities, and viscosity decreases with decreasing molecular weight. The viscosity (η) of polymer electrolytes is inversely proportional to mobility (µ) (µ ) q/6πηr, where q is the carrier charge and r is the radius of the carrier ions), which depends on the molecular weight of the polymer. Hence, polymers with different molecular weights can affect conductivity behaviors. A high viscosity obtained for a higher-molecular-weight polymer membrane results in a lower mobility (mobility is inversely related to viscosity), which decreases the conductivity. The higher viscosity attained for higher-molecular-weight samples arrests the segmental motion and free ion concentration of the polymer, which, in turn, increases the physical and thermal properties. In contrast, a decrease in the molecular weight increases the flexibility of the polymeric backbone, which, in turn, collectively enhances the ion migration and ion conductivity of the corresponding membranes.30 At the same time, lowermolecular-weight polymers result in lower thermal stabilities compared to higher-molecular-weight samples. However, the obtained thermal spectrum (Figure 5) clearly indicates that the lower-molecular-weight BMPHP polymer membrane still can be used for higher-temperature fuel cell applications. In addition, the acidification of microsized particles can become tedious, or larger-size particles cannot be completely acidified. The acid molecules can easily penetrate into the nanosized ball-milled particles and acidify the polymer particles to a larger extent. When solvated, the acid provides additional protons and is responsible for the high ionic conductivity of BMPHP membrane. Meanwhile, the acid molecules cannot easily penetrate into the bulkier PVdF-HFP polymer because of its larger area, which results in a lower ionic conductivity. As shown in Figure 7, an increase in the ball-milling time effectively increases the ionic conductivity, which is attributed to reductions in both the polymer particle size and the viscosity. Figure 8 displays surface and cross-sectional morphological images of the fabricated BMPHP porous membranes before and after liquid electrolyte activation. Phase separation did not occur in the fabricated membrane, as revealed by the cross-sectional image. From the SEM images, it is clear that the ball-milled membrane exhibited high liquid electrolyte absorption and that the electrolyte completely covered the porous nature of the membrane and was responsible for the high ionic conductivity. Figure 8 shows a cross-sectional image of the BMPHP porous membrane after the post-test analysis. It is clear that there was no significant leaching of the silica sulfuric acid for the prepared membranes because of the strong attachment with the PVdFHFP polymer. Although a few pores are visible after the posttest analysis, they are present in a negligible amount and could not affect the electrochemical properties. Figure 9a shows the EDX spectrum of the ball-milled powder membrane. The presence of carbon and fluorine elements indicates the presence of the host polymer matrix, PVdF-HFP. The lack of other peaks in the spectrum confirms the high purity
1286
Ind. Eng. Chem. Res., Vol. 49, No. 3, 2010
Figure 8. (I) Surface morphology and (II) cross-sectional images of BMPHP membranes (a) before and (b) after activation of an electrolyte and (c) after the post-test analysis (II only).
Figure 9. EDX spectra of the BMPHP membrane (a) before and (b) after activation with an electrolyte.
of the ball-milled powder. Figure 9b shows the EDX spectrum of the ball-milled PVdF-HFP membrane activated with silica sulfuric acid. The entrapment of silica sulfuric acid in the porous membrane is confirmed by the presence of silica, sulfur, and oxygen elements. After activation with the liquid electrolyte solution, porosities were calculated for the fabricated membranes and are included in Table 1. The BMPHP membrane exhibited a large decrement in the porosity than the PHP membrane (Table 1) and ensured its high electrolyte absorption. The durability of the ionic conductivity of the fabricated membranes was measured in terms of time (Figure 10). Although a drop in the ionic conductivity occurred over time, it did not lead to a dramatic drop as occurs for conventional porous polymer electrolyte membranes. After 11 days, the ionic conductivity values became constant, which confirms the durability of the membranes. 3.9. Methanol Permeability. The methanol permeability of the fabricated membranes is given in Table 1. The highly porous nature of the membrane creates a high methanol molecule permeation that degrades the DMFC performance. An increase in the methanol concentration gradually increases the methanol permeability of the fabricated membranes up to 4 M, and beyond that concentration, the methanol permeability rapidly increased.
Figure 10. Ionic conductivity of (a) PHP and (b) BMPHP membranes measured at 100% RH and 100 °C with respect to time.
The lower methanol permeabilities obtained for the fabricated membranes are attributed to the hydrophobic nature of the PVdF-HFP copolymer.1 The BMPHP membrane exhibited a
Ind. Eng. Chem. Res., Vol. 49, No. 3, 2010
Figure 11. DMFC performances of (a) PHP and (b) BMPHP polymer electrolyte membranes activated with silica sulfuric acid measured at a 2 M concentration of methanol and 80 °C.
1287
However, the path of travel of the methanol molecules becomes tedious for the BMPHP membrane because of the narrower transportation channels, which, in turn, results in a lower methanol permeability. In addition, polymer particles acidified to a greater extent induce additional protons through solvation with water and promote the ionic conductivity of the BMPHP membrane. Complete acidification of the polymer chains, decreased viscosity of the polymeric solution, and high electrolyte uptake collectively enhance the ion transport. Meanwhile, complete coverage of the porosity by the liquid electrolyte and formation of a layer over the surface of the membrane are responsible for the decrement in methanol permeation for the BMPHP membrane compared to the PHP membrane. Thus, an increase in the ion transport and an effective blockage of the molecular permeation of BMPHP membrane evokes the overall performance of direct methanol fuel cells. An increase in the methanol concentration effectively increases the DMFC efficiency up to 4 M concentration by providing a higher number of protons. However, beyond a concentration of 4 M, the DMFC efficiencies of the prepared membranes decreased as a result of higher fuel permeation. The effortless preparation, lack of environmental hazards, low costs, higher ionic conductivity, lower fuel permeability, and comparable DMFC performances of the fabricated membranes influence their prospective application over Nafion membranes, so that the fabricated membranes can be considered as a potential competitor for the commercially available membranes. 4. Conclusions Pure PVdF-HFP and ball-milled PVdF-HFP polymer electrolyte membranes were fabricated for use in direct methanol fuel cells. A thin layer formed over the porosity of the polymer matrix for the BMPHP membrane. A decrease in the particle size and molecular weight of the copolymer effectively absorbs or allows high acidification and increases the segmental motion of ions, respectively. The overall porous nature of the membrane was fully covered by effective liquid electrolyte absorption, which lowered the methanol permeability of the BMPHP membrane, and the prefermenting performance obtained from the ball-milled PVdF-HFP membrane satisfies the need for polymer electrolytes in the DMFC field.
Figure 12. DMFC performances of (a) PHP and (b) BMPHP polymer electrolyte membranes activated with silica sulfuric acid, measured at a 4 M concentration of methanol and 80 °C.
lower methanol permeability than the PHP membrane. The BMPHP membrane effectively promotes absorption of the liquid electrolyte, which, in turn, completely covers the porous nature of the membrane. In addition, the thin layer formed over the porosity of the ball-milled PVdF-HFP membrane was also a main reason for the observed lower methanol permeability. 3.10. DMFC Performance. The proton conductivity-methanol permeability selectivity ratio effectively determines the overall performance of direct methanol fuel cells. The BMPHP membrane exhibited a higher selectivity ratio than the PHP membrane. Figures 11 and 12 display the DMFC performances of the fabricated membranes at 2 and 4 M concentrations of methanol. An increase in the methanol concentration increases the fuel cell efficiency of the fabricated membranes, as shown in Figures 11 and 12. Maximum power densities of 44 and 32 mW/cm2 and current densities of 178 and 101 mA/cm2 were obtained for the BMPHP and PHP membranes, respectively, at a 4 M concentration of methanol (Figure 12). Methanol molecules can easily permeate through the PHP membrane.
Acknowledgment This work was supported by the Ministry of Knowledge Economy (MKE) through the Specialized Graduate School program. Literature Cited (1) Gnana kumar, G.; Nahm, K. S.; Elizabeth, R. N. Electro chemical properties of porous PVdF-HFP membranes prepared with different nonsolvents. J. Membr. Sci. 2008, 325, 117. (2) Hacquard, A. Improving and understanding direct methanol fuel cell (DMFC) Performance. M.Sc. Thesis. Worcester Polytechnic Institute, Worcester, MA, 2005. (3) Okamoto, K.; Yin, Y.; Yamada, O.; Islam, M. N.; Honda, T.; Mishima, T.; Suto, Y.; Tanaka, K.; Kita, H. Methanol permeability and proton conductivity of sulfonated co-polyimide membranes. J. Membr. Sci. 2005, 258, 115. (4) Neburchilov, V.; Martin, J.; Wang, H.; Zhang, J. A review of polymer electrolyte membranes for direct methanol fuel cells. J. Power Sources 2007, 169, 221. (5) Heinzel, A.; Barraga´n, V. M. A review of the state-of-the-art of the methanol crossover in direct methanol fuel cells. J. Power Sources 1999, 84, 70.
1288
Ind. Eng. Chem. Res., Vol. 49, No. 3, 2010
(6) Yamauchi, A.; Ito, T.; Yamaguchi, T. Low methanol crossover and high performance of DMFCs achieved with a pore-filling polymer electrolyte membrane. J. Power Sources 2007, 174 (1), 170. (7) Prashantha, K.; Park, S. G. Nanosized TiO2-filled sulfonated polyethersulfone proton conducting membranes for direct methanol fuel cells. J. Appl. Polym. Sci. 2005, 98, 1875. (8) Park, Y. S.; Yamazaki, Y. Low water/methanol permeable Nafion/ CHP organic-inorganic composite membrane with high crystallinity. Eur. Polym. J. 2006, 42 (2), 375. (9) Cao, J. H.; Zhu, B. K.; Xu, Y. Y. Structure and ionic conductivity of porous polymer electrolytes based on PVDF-HFP copolymer membranes. J. Membr. Sci. 2006, 281, 446. (10) Croce, F.; Hassoun, J.; Tizzani, C.; Scrosati, B. Nanoporous composite, low cost, protonic membranes for direct methanol fuel cells. Electrochem. Commun. 2006, 8, 1125. (11) Yamaguchi, T.; Kuroki, H.; Miyata, F. DMFC performances using a pore-filling polymer electrolyte membrane for portable usages. Electrochem. Commun. 2005, 7, 730. (12) Yamamoto, D.; Munakata, H.; Kanamura, K. Development of low cross-over composite membrane composed of porous polyimide and polymer electrolyte for direct methanol fuel cells. In Meeting Abstracts, 214th ECS Meeting, October 12-October 17, 2008, Honolulu, HI; The Electrochemical Society: Pennington, NJ, 2008; MA2008-02, B8 - PEM Fuel Cells 8, Abstract 1109. (13) Gnana kumar, G.; Senthilarasu, S.; Lee, D. N.; Kim, A.; Kim, P.; Nahm, K. S.; Lee, S. H.; Elizabeth, R. N. Synthesis and characterization of aligned SiO2 nanosphere arrays: Spray method. Synth. Met. 2008, 158, 684. (14) Gnana kumar, G.; Lee, D. N.; Kim, A.; Kim, P.; Nahm, K. S.; Elizabeth, R. N. Structural and transport properties of porous PVdF-HFP electrolyte membranes modified with an inorganic filler. Compos. Interfaces 2008, 15 (7-9), 731. (15) Gnana kumar, G.; Lee, D. N.; Kim, P.; Nahm, K. S.; Elizabeth, R. N. Poly(Vinylidene fluoride-co-hexa fluoropropylene)/Poly vinyl alcohol porous membranes for the application of fuel cells. J. Polym. Res. 2009, 16, 55. (16) Gnana kumar, G.; Uthirakumar, P.; Nahm, K. S.; Elizabeth, R. N. Fabrication and electro chemical properties of poly vinyl alcohol/para toluene sulfonic acid membranes for the applications of DMFC. Solid State Ionics 2009, 180, 282. (17) Ramya, K.; Dhathathreyan, K. S. Direct methanol fuel cells: determination of fuel crossover in a polymer electrolyte membrane. J. Electroanal. Chem. 2003, 542, 109. (18) Gnana kumar, G.; Kim, P.; Nahm, K. S.; Elizabeth, R. N. Structural characterization of PVdF-HFP/PEG/Al2O3 proton conducting membranes for fuel cells. J. Membr. Sci. 2007, 303 (1-2), 126.
(19) Li, Z. H.; Zhang, H. P.; Zhang, P.; Wu, Y. P.; Zhou, X. D. Macroporous nanocomposite polymer electrolyte for lithium-ion batteries. J. Power Sources 2008, 184, 562. (20) Kim, T.; Kang, I. K.; Cho, G.; Park, K. New Method for the Preparation of Solid Polymer Electrolyte Based on Poly(vinylidene fluorideco-hexafluoropropylene). Korean J. Chem. Eng. 2005, 22 (2), 234. (21) Li, Q.; He, R.; Jensen, J. O.; Bjerrum, N. Approaches and recent development of polymer electrolyte membranes for fuel cells operating above 100 °C. J. Chem. Mater. 2003, 15, 4896. (22) Xiao, G.; Li, Q.; Hjuler, H. A.; Bjerrum, N. J. Hydrogen oxidation on gas diffusion electrodes for phosphoric acid fuel cells in the presence of carbon monoxide and oxygen. J. Electrochem. Soc. 1995, 142, 2890. (23) Oetjen, H. F.; Schmidt, V. M.; Stimming, U.; Trila, F. Performance data of a proton exchange membrane fuel cell using H2/CO as fuel gas. J. Electrochem. Soc. 1996, 143, 3838. (24) Aravindan, V.; Vickraman, P. Characterization of SiO2 and Al2O3 incorporated PVdF-HFP based composite polymer electrolytes with LiPF3(CF3CF2)3.. J. Appl. Polym. Sci. 2008, 108, 1314. (25) Li, Z.; Su, G.; Gao, D.; Wang, X.; Xiaoping, L. Effect of Al2O3 nanoparticles on the electrochemical characteristics of P(VDF-HFP)-based polymer electrolyte. Electrochim. Acta 2004, 49, 4633. (26) Tezuka, T.; Tadanaga, K.; Hayashi, A.; Tatsumisago, M. Inorganicorganic hybrid membranes prepared from 3-aminopropyltriethoxysilane and sulfuric acid as anhydrous proton conductors. Solid State Ionics 2007, 178, 705. (27) Gnana kumar, G.; Senthilarasu, S.; Lee, D. N.; Kim, A.; Kim, P.; Nahm, K. S.; Lee., S. H.; Elizabeth, R. N. Synthesis and characterization of aligned SiO2 nanosphere arrays: Spray method. Synth. Met. 2008, 158, 684. (28) Li, L.; Li, G.; Smith, R. L.; Inomata, H. Microstructural evolution and magnetic properties of NiFe2O4 nanocrystals dispersed in amorphous silica. Chem. Mater. 2000, 12, 3705. (29) Zhang, P.; Zhang, H. P.; Li, G. C.; Li, Z. H.; Wu, Y. P. A novel process to prepare porous membranes comprising SnO2 nanoparticles and P(MMA-AN) as polymer electrolyte. Electrochem. Commun. 2008, 10 (7), 1052. (30) Kumar, R.; Sekhon, S. S. Effect of molecular weight of PMMA on the conductivity and viscosity behavior of polymer gel electrolytes containing NH4CF3SO3. Ionics 2008, 14, 509.
ReceiVed for reView June 22, 2009 ReVised manuscript receiVed October 31, 2009 Accepted December 4, 2009 IE901008K
