Novel Proton Exchange Membranes Based on Sulfonated Poly(acrylonitrile-co-glycidyl methacrylate)/Poly(vinyl chloride) Composite

Abstract

In this study, novel proton exchange membranes (PEMs) based on a composite of sulfonated polyacrylonitrile (SPAN), sulfonated polyglycidyl methacrylate (SPGMA), or sulfonated poly(acrylonitrile-co-glycidyl methacrylate) (SP(AN-co-GMA))/polyvinyl chloride (PVC) were developed to be used for direct methanol fuel cells (DMFCs). After polymerization and sulfonation of the prepared polymers, the polyelectrolyte membranes were prepared by the casting and solvent evaporation technique for sulfonated homo- or co-polymers with polyvinyl chloride (PVC) composites. The resulting membranes were characterized by Fourier infrared and Raman spectral analyses, X-ray diffractometry, and scanning electron microscopy. The findings of this study reveal that both the thermal stability and ion exchange capacity of the composite membranes based on sulfonated copolymers were higher than that of their corresponding composites based on sulfonated homopolymers. In this context, the weight loss percentage of the prepared composite polyelectrolyte membranes did not exceed 12% of their initial weights. The IEC of all the composite membranes ranged from 0.18 to 0.48 meq/g. Thus, the IEC value increased with the increasing proportion of the glycidyl methacrylate comonomer. Moreover, the prepared PEMs based on SP(AN-co-GMA)/PVC composites showed lower methanol permeability (8.7 × 10⁻⁷ cm²/s) than that of the Nafion membranes (3.39 × 10⁻⁶ cm²/s). Therefore, these prepared PEMs are a good candidate for DMFCs applications.

1. Introduction

The use of fossil fuels results in an increasing amount of greenhouse gas emissions, toxic smoke, and dust, which harm human health and the environment. It is urgent to develop electric vehicles powered by renewable energy sources, such as hydrogen-based fuel cells, in place of conventional fuel vehicles [1]. Fuel cells that use proton exchange membranes have made tremendous signs of progress in the previous decade as clean energy candidates. Due to their simplicity and high-power density, PEMFC-powered cars have become increasingly popular. However, due to their elevated cost and limited operating temperature, they are not widely used. High temperature and low humidity are two desirable conditions for a good PEM [2,3]. A high-water uptake, low methanol uptake, and low methanol crossover (MCO) should be present in the membrane, especially at temperatures more than 800C, for enhanced crossover (CO) tolerance. The ionic conductivity should be good, as should the chemical and mechanical endurance. As a result, the production costs should be as low as possible [2,4,5]. Nafion is a perfluorosulfonic acid membrane and has a rigid and hydrophobic PTFE backbone, providing good chemical and mechanical stability. Due to the significant phase separation between the hydrophobic backbone and hydrophilic sulfonic acid groups at the end of the side chain when the membrane is hydrated, Nafion possesses a very high proton conductivity. This is due to an increased methanol crossover and a higher manufacturing cost. To increase the performance, a polymer composite and composite membranes have been synthesized [6,7].

Due to its good film-forming properties, in addition to good mechanical strength and chemical and thermal stability, as well as the low cost of PVC, improving the ion conductivity of PVC has been suggested in this work by its combination with functional polymers, i.e., sulfonated homopolymers (SPAN and PGMA) and copolymers (SP(AN-co-GMA). These novel proton-conducting polymer composite membranes based on SPAN/PVC, SPGMA/PVC, and SP(AN-co-GMA)/PVC have been prepared for the first time and their structural and functional properties have been characterized. Moreover, the impact of AN:GMA monomeric ratios has been studied.

2. Materials and Methods

2.1. Materials

Polyvinyl chloride (PVC) (molecular weight~48,000) fine powder (Belami Fine Chemicals, Mumbai, India), acrylonitrile (AN) (purity 99%) (Acros organics, Waltham, MA, USA), glycidyl methacrylate (GMA) (purity 95%) (Acros organics, USA), tetrahydrofuran (THF) (Panreac, EU), N, N dimethylformamide (DMF) (Panreac, EU), ethanol (International Company, Suez, Egypt), methanol (El Salam Company, Suez, Egypt), and sulfuric acid (United Company, Cairo, Egypt) were used.

2.2. Preparation of Homo- and Co-Polymers

Monomeric solutions of AN or GMA (10%, v/v) were prepared in an aqueous ethanolic solution (50%) with 0.01M of potassium persulphate (KPS) as an initiator. The pre-prepared monomeric solutions were heated in a shaking water bath at 55 °C for 4 h to accomplish the polymerization process. Afterwards, the obtained homo-polymers were filtrated and periodically rinsed with an aqueous ethanol solution to remove the unreacted monomer and then dried in a conventional oven at 70 °C for 4 h, as shown in Figure 1. By the same aforementioned technique, the co-polymers (AN-co-GMA) were prepared using varying AN:GMA monomeric ratios (2:1, 1:1, or 1:2) [8,9].

2.3. Sulfonation of Homo- and Co-Polymers

Fine powders of the prepared homo-polymers (PAN and PGMA) or co-polymers (AN-co-GMA) were dispersed into sodium sulfite (Na₂SO₃) (4% w/v in an aqueous ethanolic solution; 30% v/v). The resultant dispersions were heated at 80 °C for 3 h, as shown in Figure 2. Then, the resulting sulfonated polymers were rinsed with distilled water several times to eliminate the residue of the unreacted sodium sulphide and dried in an oven at 70 °C for 24 h [10,11].

2.4. Preparation of Polyelectrolyte Membranes (PEMs)

Polyvinyl chloride (PVC) (0.6 gm) was dissolved into THF (10 mL). Then, sulfonated homo-polymers (SPAN or SPGMA) or sulfonated copolymers (SPAN-co-GMA) prepared from their corresponding co-polymers preprepared at varying AN:GMA co-monomer ratios of 2:1, 1:1, or 1:2 (0.15 gm) were dissolved in DMF (5 mL) via stirring. Then, the resulting solutions of the PVC and sulfonated homo- and co-polymers were blended using a mechanical stirrer. Afterwards, the membrane-forming solutions were cast onto a glass plate and allowed to dry at ambient temperature. The dried membranes were immersed in sulfuric acid (5 M) at room temperature for 24 h to activate their surface. Finally, the activated membranes were immersed in deionized water for another 24 h to remove the excess sulfuric acid [12]. The dried membranes were manually removed from the petri dishes and conditioned in a humidity chamber at 25 ± 2 °C with 50 ± 2% relative humidity in an environmental chamber before further mechanical testing.

2.5. Characterization

2.5.1. Structural Properties

FTIR and Raman scattering spectral analyses were conducted to explore the chemical structure of the polymeric substances and to confirm the success of the synthesis process. These analyses were performed using a Shimadzu FTIR-8400 S (Shimadzu, Kyoto, Japan) and a Raman scattering spectrometer (SENTERRA-Bruker, Germany) [13,14]. Moreover, X-ray diffractometry was performed to study the crystalline structure of the prepared polymeric substances using an X-ray diffractometer (Shimadzu XRD-7000, Japan). The crystallinity degree of such polymeric substances and their composites was calculated according from the following equation:

% Crystallinity = A_Crystalline × 100/A_Total

where A_Crystalline represents the area of the crystalline regions and A_Total represents the total area under the peak [15,16].

2.5.2. Thermal Stability

The thermal stability of the resulting polymeric materials was assessed using thermogravimetric analysis. This analysis was performed using a thermogravimetric analyzer (Shimadzu TGA-50, Shimadzu, Kyoto, Japan) by measuring the weight loss of the heated substances during the gradual increase in the temperature at a constant heating rate (10 °C/min) under N₂ gas. The specimens were heated in the temperature range from ambient to 700 °C [15].

2.5.3. Mechanical Strength

The tensile strength of the polymer electrolyte membranes was determined using a universal testing machine (Shimadzu UTM, Japan) and, as previously mentioned, the dried membranes were conditioned in a humidity chamber at 25 ± 2 °C with 50 ± 2% relative humidity before further mechanical testing. The specimens’ dimensions were 30 × 10 mm. The measurements were carried out at a constant speed of the cross-heads’ movement of 5 mm/min.

2.5.4. Morphological Features

The morphological features and microstructure of the polymeric materials and membranes were examined using a JEOL JSM-6360LA SEM (Jeol, Akishima, Japan). The SEM was operated at an acceleration voltage of 10–20 KV after the samples were mounted on stainless steel stubs with double-sided tape; a 10–20 nm thick layer of gold was sputtered on the samples [17,18].

2.5.5. Surface Roughness

The surface roughness of the PEMs was measured with a surface roughness meter SJ-20111 Portable Surface Testers from Japan. Software-Sj tools Double-sided tape was used to secure the samples on a glass slide. All of the data are the mean of ten measurements [19].

2.5.6. Wettability

The wettability of the resulting PEMs was assessed by measuring the contact angle of the water droplet with the polyelectrolyte membrane with a VCA 2500 XE coupled with a CCD camera and the ramé-hart DROP image series software (AST Products, Billerica, MA, USA) [20].

2.5.7. Ion Exchange Capacity

The IEC of the PEMs based on the PVC and sulfonated polymers composites is dependent on the number of sulphonic groups substituted on the functional reactive groups of the monomeric units. The weighed membrane samples were submerged in a 2 M NaCl solution for 24 h. Acid–base titration using 0.01 N sodium hydroxide and phenolphthalein as an indicator was used to determine the proton emitted from these samples. The IEC of the polyelectrolyte membrane was calculated by using the following equation.

$$\mathrm{I}\mathrm{E}\mathrm{C}(\mathrm{m}\mathrm{e}\mathrm{q}/\mathrm{g})=\frac{\mathrm{N}\times \mathrm{V}}{\mathrm{W}}$$

where N and V are the normality and volume titer of the sodium hydroxide solution, respectively, and W denotes the sample weight [21].

2.5.8. Water and Methanol Uptake

The polyelectrolyte membrane specimen was soaked in deionized water at room temperature for 24 h. Afterwards, free water on the membrane surface was absorbed with filter paper, and the membrane specimen was promptly weighed [3]. The water uptake was calculated by using the following equation.

$$W\%=\frac{W_{wet}-W_{dry}}{W_{dry}}\times 100$$

where W_wet and W_dry are the wet and dry membrane weights, respectively. By the same aforementioned technique, the methanol uptake was determined by replacing water with methanol [22].

2.5.9. Methanol Permeability

Methanol permeability is a measure of the membrane’s resistance to methanol crossover from cathode-to-anode. A diffusion glass cell composed of two similar compartments was used in this test, wherein the PEM specimen was tightly clamped between these compartments, A and B, which were filled with 2 M of a methanol solution and deionized water, respectively. The diffusion cell was stirred for one hundred minutes at room temperature. At an interval of twenty minutes, one ml of compartment B was taken to measure the amount of transmitted methanol using gas chromatography. The transmission curve was obtained by plotting methanol concentration as a function of time and the following equation was used to calculate the methanol permeability (P) [9]:

$$P=\frac{K\times V\times L}{A\times C_A}$$

where K is the slope of the transmission curve, V is the initial volume of deionized water, L is the membrane thickness, A is an area of permeation, and C_A is the initial concentration of methanol.

3. Results and Discussion

3.1. Structural Properties

3.1.1. FTIR Spectra

The FTIR spectra of PAN, SPAN, PGMA, SPGMA, P(AN-co-GMA), and SP(AN-co-GMA) are presented in Figure 3A. The FTIR spectrum of PAN revealed a characteristic absorption peak at 2256 cm⁻¹ due to C≡N stretching [23]. On the other hand, the FTIR spectrum of SPAN revealed an absorption peak at 1134 cm⁻¹ due to the symmetric vibration of O=S=O [24]. The FT-IR spectrum of PGMA revealed a strong characteristic absorption band at 1716 cm⁻¹ due to the stretching vibration of C=O belonging to methacrylate ester. Another characteristic peak appeared at 1153 cm⁻¹ due to the typical stretching vibration of C-O-C. Moreover, an absorption band appeared at 879 cm⁻¹ due to the epoxy ring. However, in the SPGMA spectrum, the intensity of the absorption band of the epoxy rings was lowered and shifted from 879 to 613 cm⁻¹. Additionally, an absorption peak belonging to the OH group was shifted from 3477 to 3443 cm⁻¹. These changes can be attributed to the opening of a fraction of epoxy rings via the sulfonation process [25]. In addition, a characteristic absorption band was observed at 1149 cm⁻¹ due to the stretching of the sulphonate groups. The FTIR spectrum of poly(AN-co-GMA) exhibited an absorption peak at 2261 cm⁻¹ due to the C≡N group of acrylonitrile [3]. Moreover, absorption peaks at 1701 and 3433 cm⁻¹ appeared as a result of the stretching of the carbonyl and hydroxyl groups, respectively [26]. The absorption band appeared at 1153 cm⁻¹ due to the vibration of the sulphonate groups, confirming sulfonation of poly(AN-co-GMA).

Figure 3B depicts the FTIR spectra of the PVC–SP(AN-co-GMA)/PVC composite. The spectrum of PVC displayed an absorption band at 642 cm⁻¹ due to C-Cl stretching. Moreover, absorption peaks at 1384 and 2939 cm⁻¹ were observed corresponding to the C-C and C-H vibrations, respectively [26]. On the other hand, the FTIR spectrum of PVC with sulfonated polymers revealed a slight shift for the adsorption band of C-Cl from 634 to 622 cm⁻¹ and a decrease in its intensity, which can be attributed to the reduction of the PVC proportion in the composites. Furthermore, new absorption peaks appeared at 1199 and 3443 cm⁻¹ due to the stretching of the sulphonate and hydroxyl groups, respectively, in addition to the adsorption peak at 1700 cm⁻¹ belonging to the C=O group. However, no new peaks appeared as a result of the chemical reaction; this proves that only a physical combination occurred between the two polymeric substances.

3.1.2. Raman Scattering Spectra

Raman spectral analysis is a complementary technique to infrared spectroscopy that enables the identification of characteristic vibrational bands of molecular structures. The Raman scattering spectra of PAN, SPAN, PGMA, SPGMA, P(AN-co-GMA), and SP(AN-co-GMA) are presented in Figure 4A. The Raman spectrum of PAN showed an absorption peak at 2240 cm⁻¹ attributed to the stretching of C≡N [27]. The SPAN spectrum exhibited a characteristic band at 1648 cm⁻¹ due to stretching of the C=N group [28]. Furthermore, an absorption peak was observed at 1182 cm⁻¹ due to the stretching of O=S=O [29]. The Raman scattering spectrum of PGMA revealed absorption bands at 1405–1455 cm⁻¹ due to the stretching of the methylene groups. In addition, another absorption band was observed at 1710–1745 cm⁻¹ due to the vibration of the ester groups. However, the typical absorption band of the epoxy ring appeared at 800–950 cm⁻¹ [30]. In the SPGMA spectrum, the characteristic absorption band of the sulphonate group was observed at 1044 and 1166 cm⁻¹ [31]. The spectrum of poly(AN-co-GMA) revealed absorption peaks at 2238 and 1722 cm⁻¹ corresponding to the vibration of the C≡N and carbonyl groups, respectively [32]. In addition, a characteristic band was observed at 800–950 cm⁻¹ due to the stretching of C-O-C.

The Raman spectra of the PVC and sulfonated polymers/PVC composites are shown in Figure 4B. The Raman spectrum of the PVC exhibited a strong peak at 637 cm⁻¹ corresponding to the C-Cl bonds [33,34]. In addition, the Raman spectra of the composites of sulfonated polymers with PVC showed a slight shift of the C-CL bond in addition to the appearance of absorption peaks at 1700 cm⁻¹ due to the C=O group present in SPAN, SPGMA, and SP(AN-co-GMA), which prove the combination of the two polymeric substances.

3.1.3. X-ray Diffractograms

The X-ray diffraction patterns of PAN, SPAN, PGMA, SPGMA, P(AN-co-GMA), and SP (AN-co-GMA) are shown in Figure 5A. The X-ray diffraction patterns of the PAN and SPAN revealed two narrow peaks at 2θ = 17° and 29° corresponding to basal distances of 5.07 Å and 2.99 Å, respectively, which are typical fingerprints of the hexagonal lattice of PAN [35]. However, in the pattern of SPAN, the number of peaks and their intensity were increased because of the alignment of polymer chains and, subsequently, the formation of a higher crystalline structure as a result of introducing sulfonated groups onto PAN chains. The XRD pattern of PGMA revealed a crystalline structure with peaks appearing at 2θ = 18°. In addition, a broad peak centered at 2θ (19°) was assigned to the amorphous structure of the epoxy group [36]. On the contrary to SPAN, the XRD pattern of SPGMA exhibited substantial increases in the scattering intensities recorded, especially at scattering angles smaller than 50° (2θ = 32°, 49°, and 19°), which is most likely due to the destruction of the crystalline structure of PGMA as a result of the introduction of the sulfonic group to the polymer chains. This behavior can be explained by the elimination of the hydrogen bonding network between the epoxy rings of PGMA [37]. The XRD patterns of P(AN-co-GMA) exhibited a high crystallite size, which appears at 89.9 Å. In addition, the presence of SO₃ in SP(AN-co-GMA) also presented a high value of basal distances at 49 Å.

The X-ray diffraction patterns of the pristine PVC–SP(AN-co-GMA)/PVC composite are shown in Figure 5B. The XRD pattern of PVC exhibited diffraction peaks at 2θ = 17°, 23°, 29°, and 35°. This can be attributed to the amorphous structure of PVC [38]. On the other hand, the XRD pattern of the SP(AN-co-GMA)/PVC composite membrane showed broad characteristic peaks at 2θ = 23° and 25°, which is due to the alignment of the sulfonated polymers chains, consequently increasing the crystallinity of composite matrix. This can be confirmed by the crystallinity degree of these composites, which is presented in Figure 5C. From this figure, it can be seen that PVC is amorphous, while the sulfonated co-polymers were semi-crystalline. Hence, an incorporation of sulfonated co-polymers into the PVC matrix led to an increase of the crystallinity degree of the resulting composites. However, this increase was more pronounced in the case of SPGMA/PVC than in SPAN/PVC. Therefore, the increase of the crystallinity degree in the SP(AN-co-GMA)/PVC composite was associated with the increase of the GMA proportion in SP(AN-co-GMA).

3.2. Thermal Stability

To assess the thermal stability of the polymeric substances, thermogravimetric analysis (TGA) was carried out in the range of ambient temperature to 700 °C in a static atmosphere of nitrogen. The thermograms of the homo-and co-polymers and their sulfonated derivatives are shown in Figure 6 and the thermal decomposition stages are tabulated in

Table 1. TGA data of PAN, PGMA, and P(AN-co-GMA) and their sulfonated derivatives, and pristine PVC and its composites with sulfonated polymers.

Polymer	Stage	Temperature Range °C	Weight Loss%	Thermal Decomposition
PGMA	First	0–145	1.5	About 120 °C, due to water evaporation. About 300 °C is attributed to polymer surface decomposition and random chain scission.
	Second	145–300	7
	Third	300–428	80
	Fourth	428–700	10
SPGMA	First	0–132	7	About 120 °C, due to loss of water. At 250 °C correspond to the decomposition of the sulfonic groups.
	Second	132–250	11
	Third	250–500	25
	Fourth	500–700	10
PAN	First	0–120	2.58	Above 300 °C, due to the backbone degradation of PAN.
	Second	120–350	20
	Third	350–575	15.2
	Fourth	575–700	9.6
SPAN	First	0–164	5	About 250 °C corresponds to the decomposition of the sulfonic groups.
	Second	164–364	5
	Third	364–471	3
	Fourth	471–700	20
P(AN-co-GMA)	First	0–134	5.2	About 400 °C corresponds to the degradation and deacetylation of the polymer.
	Second	134–334	3.3
	Third	334–483	80.33
	Fourth	483–700	4.75
SP(AN-co-GMA)	First	0–136	6.5	Around 250 °C is due to the thermal degradation of the sulphonic acid groups.
	Second	136–236	7
	Third	236–473	35
	Fourth	473–700	4
PVC	First	0–150	10	Below 100 °C can be attributed to the evaporation of trapped THF. At 384.4 °C, which results from the breaking of the C-Cl bond. Above 400 °C, due to the decomposition of the polymer backbones.
	Second	150–280	38
	Third	280–469	32
	Fourth	469–800	14.5
SPAN/PVC	First	0–166	10	The TGA data represent that the SPAN/PVC electrolyte membranes are thermally stable up to around 200 °C, which demonstrates sufficient thermal properties for application in DMFC.
	Second	166–264	10
	Third	264–449	32
	Fourth	449–800	20
SPGMA/PVC	First	0–170	12.58	The TGA data represent that the SPGMA/PVC electrolyte membranes are thermally stable up to around 200 °C, which demonstrates sufficient thermal properties for application in DMFC.
	Second	170–270	10
	Third	270–475	15.2
	Fourth	475–800	9.6
SP(AN-co-GMA)/PVC	First	0–136	5.5	The TGA data represent that the SP(AN-co-GMA)/PVC electrolyte membranes are thermally stable up to around 200 °C, which demonstrates sufficient thermal properties for application in DMFC.
	Second	136–250	7
	Third	250–473	35
	Fourth	473–800	4

. As shown in Figure 6A, the thermal stability of the sulfonated PAN was higher than PAN. The overall weight loss was 45% of the initial weight of PAN; however, its corresponding weight loss was 33% in the case of SPAN [39]. The TGA thermogram of PGMA and its sulfonated form shown in Figure 6B exhibited that at 300 °C, the weight loss of PGMA reached 55% of its initial weight, while it was 18% for the sulfonated form at the same temperature. About 97% of the initial weight of PGMA was lost during heating to 700 °C, while the sulfonated form lost only 51% of its initial weight at this temperature. In the same context, the T₅₀ recorded for PGMA and SPGMA were 335 and 527 °C, respectively. These findings confirm the increasing thermal stability of PGMA as a result of the sulfonation process. This behavior can be explained by the increase in the alignment of polymer chains and crystallinity of the PGMA via sulfonation [40]. The thermograms of P(AN-co-GMA) and its sulfonated form presented in Figure 6C exhibited that about 86% of the initial weight of the copolymer was lost when it was heated to 700 °C. Meanwhile, at the same temperature, the weight loss of the sulfonated form was only 65%. Interestingly, the T₅₀ of P(AN-co-GMA), 455 °C, lies between the T₅₀ of PAN and PGMA; therefore, the high thermal stability of the sulfonated copolymer can be attributed to the hydrogen bonding network forming as a result of the sulfonic groups.

The thermograms of the pristine PVC and SP(AN-co-GMA)/PVC composite are presented in Figure 7 and the thermal decomposition stages are tabulated in

Table 2. Mechanical properties of pristine PVC and sulfonated homo- and co-polymer/PVC composite membranes.

Polyelectrolyte Membrane	Thickness (mm)	Tensile Strength (MPa)	Elongation at Break (%)
Nafion® 117	0.183	12.2 ± 0.1	18.2
PVC	0.14 ± 0.02	6.22 ± 0.50	8.66 ± 0.60
SPAN/PVC	0.18 ± 0.08	1.73 ± 0.10	54.10 ± 2.20
SPGMA/PVC	0.18 ± 0.09	1.12 ± 0.20	114.51 ± 5.80
SP(AN-co-GMA) 2:1/PVC	0.14 ± 0.08	0.94 ± 0.08	23.65 ± 1.30
SP(AN-co-GMA) 1:1/PVC	0.10 ± 0.07	1.01 ± 0.09	25.42 ± 2.20
SP(AN-co-GMA) 1:2/PVC	0.15 ± 0.07	0.96 ± 0.05	20.21 ± 1.50

. These data show that the weight loss percentages at different temperatures for the composite are less than that of pristine PVC, confirming its higher thermal stability. Generally, the thermogram of both the pristine PVC and the composite exhibited a minor transition at 100 °C corresponding to a weight loss of absorbed water. However, sulphonic groups decompose at about 250 °C. The main thermal decomposition around 380 °C is caused by polymer chain scission. As mentioned above, the higher thermal stability of composite membrane can be attributed to increase of its crystallinity degree, owing to incorporation of a partially crystalline co-polymer into the PVC matrix. From these findings, it was concluded that the composite membranes have thermal stability appropriate for use in DMFCs, as thermal breakdown for the functional groups and the polymer backbone was found to occur above 200° [41,42].

3.3. Mechanical Properties

The mechanical properties of PEMs have a significant impact on the consistent performance of fuel cells. The mechanical properties of pristine PVC and sulfonated homo- and co-polymers composite membranes are listed in Table 2. These findings demonstrate that an incorporation of each of the sulfonated homo- or co-polymers into the PVC matrix significantly reduced the tensile strength of the obtained composite membrane. Meanwhile, their elongation percentages were higher than that of the pristine PVC membrane. The composite membranes based on sulfonated polyacrylonitrile exhibited higher tensile strength than those of the composite membranes based on SPGMA or SP(AN-co-GMA) preprepared from copolymers with varying AN:GMA monomeric ratio. On the other hand, a combination of PVC with sulfonated copolymers resulting from P(AN-co-GMA) prepared at varying AN:GMA monomer ratios of 2:1, 1:1, or 1:2 can be utilized to obtain a membrane with a tensile strength comparable to those recorded in the case of sulfonated polyglycidyl methacrylate [43]. Meanwhile, the elongation percentage for the SPGMA/PVC composite-based membrane (115%) was almost double that of SPAN/PVC composite-based membranes (54%) and was much higher than that of the SP(AN-co-GMA)/PVC composite-based membrane (20–25%). This drastic drop in the mechanical strength of PVC when incorporating sulfonated homo- or co-polymers can be explained on the basis that the sulfonated polymers’ particulates might act as stress concentrators, resulting in early failure of the membrane. On the contrary, the more pronounced decline in the tensile strength of the SPGMA/PVC and SP(AN-co-GMA)/PVC composite-based membranes compared with those based on the SPAN/PVC composite may be attributed to changes in the structural and conformational properties of these sulfonated polymers that define the extent of compatibility between the sulfonated polymers and PVC and, subsequently, the dispersion of these modified polymers within the PVC matrix. Since SPGMA and SP (AN-co-GMA) have higher crystallinity and thermal stability, in addition to more irregularities in their shape compared to SPAN (as indicated from the XRD, TGA, and SEM data), they could result in a weak interfacial interaction with the PVC matrix and poor dispersion, thereby reducing the interfacial shear strength and the strength of composite membranes. On the other hand, the positive effect of incorporating SPGMA into the PVC matrix on elongation at the break can be attributed to the microstructural features of SPGMA and their distribution into the SPGMA/PVC composite matrix, where the longitudinal orientation of the SPGMA fibers that were noted on the SEM examination (Figure 8a) can lead to fiber-related energy dissipation mechanisms, such as fiber–matrix debonding, fiber pullout, and fiber breakage restricting the deformation of the matrix during tensile loading, which results in a loss of elongation.

3.4. Morphological Features

SEM was used to examine the morphological features of the synthesized homo- and copolymers, as well as their sulfonated forms. The micrographs are presented in Figure 8a. The SEM micrographs of PAN and its sulfonated form exhibit that both have a coral reef shape as a result of the formation of the aggregated micro-particulates from the polymer chains. However, slight swelling in the polymer micro-particulates was observed as a result of the sulfonation process [44]. The SEM micrograph of PGMA shows aggregated irregular micro-globules with shapes that resemble bunches of grapes [45,46], while the sulfonated PGMA shows noticeable morphological changes, which appeared in the form of irregular flakes composed of bundles of strips (longitudinal-oriented fibers) with a rough surface [47]. The SEM micrographs of P(AN-co-GMA) exhibit a compact and dense structure with a wrinkled surface similar to dough. On the other hand, the sulfonation of P(AN-co-GMA) led to a noticeable change in the structure of the copolymer, which appears looser with irregular structures.

The SEM micrographs of the pristine PVC–SP(AN-co-GMA)/PVC composite membranes are presented in Figure 8b. The SEM micrograph of the surface and cross-section of the pristine PVC indicates that the PVC has a network structure forming non-uniform hexagonal compartments. However, the SEM micrograph of the surface of the SP(AN-co-GMA)/PVC composite membrane reveals a solid structure with a non-even surface. Similarly, the micrograph of the cross-section of the composite membrane (Figure 8c) exhibits a more dense and compact structure when compared to the corresponding micrograph of the pristine PVC membrane, which is characterized with a loose structure. The changes can be explained on the basis of the inclusion of the sulfonated polymer particulates as a filler in the porous PVC matrix. This explanation can be also confirmed from the SEM micrograph for the composite membrane surface, which reveals a non-porous structure with an even surface for this membrane, contrary to the corresponding micrograph of the pristine PVC membrane [48].

3.5. Surface Roughness

The roughness of the surface of the pristine PVC and polyelectrolyte membranes based on the PVC and sulfonated homo- and co-polymer composites are recorded in

Table 3. Surface roughness of pristine PVC and composite membranes.

Polyelectrolyte Membrane	Roughness (µm)
Nafion® 117	0.09
PVC	1.10 ± 0.02
SPAN/PVC	1.10 ± 0.05
SPGMA/PVC	3.52 ± 0.10
SP(AN-co-GMA) 2:1/PVC	3.03 ± 0.11
SP(AN-co-GMA) 1:1/PVC	2.74 ± 0.10
SP(AN-co-GMA) 1:2/PVC	3.08 ± 0.10

. The plain PVC membrane exhibited the lowest roughness value. Meanwhile, the roughness values of the SPGMA/PVC and SP(AN-co-GMA)/PVC composite-based membrane were almost triple that of the plain PVC membrane. On the other hand, the SPAN/PVC composite-based membrane had a surface roughness comparable to that of the PVC membrane [49,50]. The increase of the surface roughness for the SPGMA/PVC and SP(AN-co-GMA)/PVC composite-based membranes was ascribed to the irregular shape and greater size of the hydrophilic sulfonated homo- or co-polymers filler within the PVC matrix, resulting in their poor dispersion within the membrane matrix. On the contrary, SPAN particulates can be highly dispersed within the PVC matrix, filling its pores and resulting in a dense structure and less surface pores. This could prevent methanol diffusion within the composite membrane matrix. Despite the findings, no positive correlation was noticed between the membrane roughness and the transmission rate, as indicated by the methanol permeability test (discussed below).

3.6. Contact Angle

The water contact angle is frequently used as an indicator of the hydrophilic/hydrophobic characteristics of the membrane surface. The water contact angle values of the plain PVC and composite membranes are listed in

Table 4. Contact angle on the surface of pristine PVC and composite membranes.

Polyelectrolyte Membrane	Mean Theta (θ)
Nafion® 117	110
PVC	46.31 ± 1.20
SPAN/PVC	21.07 ± 2.08
SPGMA/PVC	23.04 ± 3.09
SP(AN-co-GMA) 2:1/PVC	30.07 ± 3.08
SP(AN-co-GMA) 1:1/PVC	22.95 ± 3.07
SP(AN-co-GMA) 1:2/PVC	20.25 ± 3.17

. In general, these results indicate that a combination of sulfonated homo- or copolymers with PVC leads to a lower contact angle of the PVC membranes. The lowest contact angle value was recorded for PEM based on SP(AN-co-GMA)/PVC at an AN:GMA monomer ratio of 1:2.

3.7. Ion Exchange Capacity

At room temperature, the IEC of the sulfonated homo- or co-polymers/PVC composite-based membranes was measured. The ion exchange capacity denotes the number of exchangeable protons or sulfonic groups within a polymer matrix, which is responsible for proton conduction and, thus, can be considered an indirect and reliable approximation of proton conductivity. The IEC of the prepared PVC and its composites with sulfonated homo-or co-polymers (SPAN, SPGMA, and SP(AN-co-GMA)) membranes are listed in

Table 5. The ion exchange capacity of pristine PVC and composite membranes.

Polyelectrolyte Membrane	IEC (meq/g)
Nafion® 117	0.91
PVC	0.01 ± 0.01
SPAN/PVC	0.20 ± 0.05
SPGMA/PVC	0.22 ± 0.05
SP(AN-co-GMA) 2:1/PVC	0.37 ± 0.02
SP(AN-co-GMA) 1:1/PVC	0.38 ± 0.02
SP(AN-co-GMA) 1:2/PVC	0.48 ± 0.05

. As is well known, the data show that the plain PVC membrane has no IEC because it has no ionized functional groups present. Meanwhile, the composite membranes recorded IEC values ranging from 0.20 to 0.48 meq/g. The composite membranes based on SP(AN-co-GMA) recorded higher IEC values than those of the composite membranes based on sulfonated homo-polymers. On the other hand, the IEC value for the composite membranes based on SPGMA was higher than that of the corresponding composite membranes based on SPAN. This behavior can be explained by the fact that more sulfonic groups can be substituted onto copolymer chains; therefore, a higher GMA content is predicted to be more reactive than the nitrile towards sodium sulfite [27]. Furthermore, the activation process has no noticeable impact on the ion exchange capacity of the resulting composite membranes.

3.8. Water and Methanol Uptake

Polyelectrolyte membrane (PEM) is a key part of the PEMFCs’ arrangement, and their proton conductivity is heavily dependent on their water content, wherein the hydration of PEM is an influential factor to improve proton conductivity and, therefore, fuel cell efficiency. On the other hand, in the case of DMFC, in which methanol is employed as a source of hydrogen, the polyelectrolyte membrane’s methanol uptake must be reduced to avoid fuel cross–over. The water and methanol uptake for the PEM prepared in this study are listed in

Table 6. Water and methanol uptake on the surface of pristine PVC and composite membranes.

Polyelectrolyte Membrane	Water Uptake (%)	Methanol Uptake (%)
Nafion® 117	65.44	22
PVC	00.05 ± 0.01	14.51 ± 0.08
SPAN/PVC	25.16 ± 1.10	16.26 ± 0.50
SPGMA/PVC	16.41 ± 1.05	07.77 ± 0.08
SP(AN-co-GMA) 2:1/PVC	20.10 ± 1.50	12.06 ± 0.09
SP(AN-co-GMA) 1:1/PVC	18.10 ± 1.05	14.14 ± 0.10
SP(AN-co-GMA) 1:2/PVC	17.65 ± 1.04	13.92 ± 0.20

. These results reveal that a combination of PVC with both sulfonated homo-or co-polymers leads to an increase in the water-holding capacity of the resultant composite-based PEMs. Generally, the percentage of water uptake was higher than the methanol uptake; therefore, the composite membranes based on either SPGMA or SP(AN-co-GMA) had a lower capacity for holding water compared to their counterparts based on SPAN. In the same context, the composite membranes based on copolymers with a higher monomer ratio of acrylonitrile exhibited a higher percentage of water uptake. This behavior can be explained by the fact that nitriles can be converted to amides by acid catalysis and become protonated. Protonation increases the electrophilicity of the nitrile so that it will accept water.

3.9. Methanol Permeability

Proton-exchange membranes in DMFCs must have low methanol permeability while maintaining good ionic conductivity. The methanol transmission of the plain PVC and composite membranes based on SP(AN-co-GMA) prepared with a copolymer with an AN:GMA monomeric ratio of 1:2 was determined using a diffusion cell, and the results are shown in Figure 9. The methanol permeability was estimated by determining the slope of the transmission curve, and the results are presented in

Table 7. Methanol permeability of pristine PVC and SP(AN-co-GMA)/PVC composite membranes.

Polyelectrolyte Membrane	Methanol Permeability (cm2 s−1)	Efficiency Factor
Nafion® 117	3.39 × 10−6	2.6 × 105
PVC	9.36 × 10−7	-
SPAN/PVC	9 × 10−7	6 × 105
SPGMA/PVC	9.16 × 10−7	5.7 × 105
SP(AN-co-GMA)/PVC	8.7 × 10−7	6.6 × 105

. These data show that the methanol permeability of the SP(AN-co-GMA)/PVC composite-based membrane is 8.7 × 10⁻⁷ cm²/s, while it was 9.36 × 10⁻⁷ and 3.39 × 10⁻⁶ cm²/s for the pristine PVC and Nafion^®, respectively. This decline in the methanol permeability of the SPAN/PVC composite-based membrane can be attributed to the dense and compact structure of the composite matrix that was previously proven from the SEM examination, in addition to the high hydrophilicity of the composite membrane’s surface, which hindered the adsorption of methanol, therefore reducing its cross-over. Compared to the SPAN/PVC composite-based membranes with a dense structure, the SP(AN-co-GMA)/PVC and SGMA/PVC composite-based membranes, in spite of their higher surface roughness and pore density, exhibited lower methanol permeability, as they led to transport of methanol molecules through a more complicated and torturous path, owing to the presence of the sulfonated polymers’ filler particulates [8,42].

The efficiency factor, a ratio of IEC/methanol permeability, is used to evaluate the performance of the polyelectrolyte membranes in DMFCs [9,48]. This factor was calculated for the SP(AN-co-GMA)/PVC composite-based membranes and is listed in Table 7. The efficiency factor of the composite membrane prepared in this study was around triple that of Nafion^®. In addition to the higher efficiency of these developed composite-based membranes compared to Nafion, they are also cost-effective. Therefore, they can be considered promising candidates for DMFC applications.

4. Conclusions

In this research work, novel proton exchange membranes based on sulfonated poly-acrylonitrile, sulfonated polyglycidyl methacrylate, and sulfonated poly(acrylonitrile-co-glycidyl methacrylate)/poly(vinyl chloride) composites were successfully developed using physical combining and casting with a solvent evaporation technique for DMFCs application. The structural and functional properties of the developed composites-based membranes were characterized using FTIR, XRD, SEM, TGA, hydrophilicity, water and methanol uptake, methanol permeability, and IEC measurements. Overall, the most desirable attributes of PEMs, such as IEC, water uptake, methanol permeability, and thermal stability, were adequate, especially when incorporating sulfonated copolymers into the PVC matrix. The IEC was 0.48 meq/g for the SP(AN-co-GMA) (1:2)/PVC composite-based membranes. Moreover, it was found that these composite-based PEMs outperform Nafion^® with their lower methanol permeability (8.7 × 10⁻⁷ cm²/s) and lower cost. However, their low mechanical strength and IEC in comparison to Nafion^®117 should be improved in future research by optimizing the design of the composite materials through precise tuning of the sulfonation and incorporation process to achieve better mechanical, thermal, and chemical stability and, therefore, increased PEM fuel cell performance.