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Embrittlement induced fracture behavior and mechanisms of perfluorosulfonic-acid membranes after chemical degradation
Journal of Power Sources 453 (2020) 227893
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Perspective
Embrittlement induced fracture behavior and mechanisms of perfluorosulfonic-acid membranes after chemical degradation Xiaoyi Sun a, Shouwen Shi a, *, Yuanjie Fu a, Jian Chen b, Qiang Lin a, Jiaqi Hu a, Cong Li b, Jiayao Li a, Xu Chen a a
School of Chemical Engineering and Technology, Tianjin University, Tianjin, 300072, China Key Laboratory of Efficient & Clean Energy Utilization, The Education Department of Hunan Province, School of Energy and Power Engineering, Changsha University of Science and Technology, Changsha, 410114, China
b
H I G H L I G H T S
G R A P H I C A L A B S T R A C T
� Fracture behavior of Nafion membranes are explored after chemical degradation. � Chemical degradation causes predomi nant loss of side chains. � Crack transforms from ductile to brittle fracture with increasing degradation level. � Stretch zone size from cross-section observation is related to crack propagation. � Reduced crack propagation resistance is ascribed to decreased plastic zone size.
A R T I C L E I N F O
A B S T R A C T
Keywords: PFSA membrane Fracture Chemical degradation Mechanical property Durability
Chemical and mechanical degradations of perfluorosulfonic-acid membranes are two factors contributing to the reduced durability of fuel cells. While the mechanisms of isolated chemical or mechanical degradation are extensively investigated, the impact of chemical degradation on mechanical degradation is not fully understood. In this paper, the fracture behavior of Nafion 212 membranes with different chemical degradation levels are investigated. To characterize the degradation level and probe the molecular origins of chemical degradation, fluoride release, Fourier-transformer infrared spectra, conductivity and swelling behavior are measured. It is found that chemical degradation causes predominant loss of side-chains. A transition from ductile to brittle fracture is observed with increasing degradation level. In addition, the cross-sections are examined to link fracture behavior with microstructure changes. While the decreased fracture toughness is attributed to reduced mechanical properties such as Young’s modulus and break strain, the decreased crack propagation resistance is ascribed to reduced plastic zone size ahead of crack tip, which reduces the plastic energy dissipation. These findings not only provide new mechanical dataset of degraded membranes for performance and durability modelling to take into account of chemical degradation effect but also improve the understanding of membrane durability.
* Corresponding author. E-mail address: [email protected] (S. Shi). https://doi.org/10.1016/j.jpowsour.2020.227893 Received 1 November 2019; Received in revised form 9 February 2020; Accepted 11 February 2020 0378-7753/© 2020 Elsevier B.V. All rights reserved.
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1. Introduction
They observed cracks on the surface of degraded samples which they attributed to reduced ductility. Even though the tensile properties of chemical degraded membranes are investigated, only basic properties are obtained from stress-strain curves, such as elastic modulus and strain-to-failure, with the fracture behavior and underlying mechanisms left unexplored. The tensile properties alone are not sufficient enough to account for the reduced durability after chemical degradation. For instance, Rodgers et al. [2] found that even though 950 EW cell had the largest losses in mechanical strength, the crossover is not significantly impacted. And they concluded that there was no correlation between mechanical strength and cross over of PFSA membrane. As gas crossover is not only related to me chanical strength, but also related to crack initiation and propagation resistance. Therefore, it is imperative to investigate the fracture behavior of PFSA membrane after chemical degradation in order to alleviate the failure issue and develop effective mitigation strategies to enhance membrane durability. In addition, most numerical simulations on mechanical responses of PFSA membranes used mechanical proper ties of as-received (AsR) membranes without considering decay in me chanical properties with chemical degradation [32–35]. Thus, the mechanical properties after chemical degradation should be systemati cally investigated, which could be used as inputs in numerical modelling to update material parameters to account for chemical degradation is sues and accurately characterize the interaction between durability with performance. Hence, the objective of this paper is to investigate the mechanical properties of Nafion 212 membrane after chemical degradation, espe cially the fracture behavior, with the aim of elucidating how chemical degradation changes the mechanical properties as well as the underlying mechanisms. To link mechanical properties with degradation levels, the total fluoride release is measured and used as an indicator for degra dation level. In addition, molecular structure changes induced by chemical degradation are also explored to account for molecular origins of mechanical changes. As stresses are generated in-situ due to swelling and shrinkage of the membrane during humidity cyclings, the swelling behavior of the membrane is also investigated to understand the dimensional stability. Finally, the tensile and fracture behaviors of membranes at different degradation levels are compared, and the crosssections of fractured membranes are examined to correlate with fracture resistances. The findings in this paper are expected to provide more insights into the decay in mechanical properties after chemical degra dation as well as the understanding of the failure mechanisms of membranes under combined chemical and mechanical degradations.
Polymer-electrolyte fuel cells (PEFCs) are promising zero-emission power in many fields such as transportation and electricity. However, the performance and durability issues are two challenges impeding the commercialization of this promising technology. Among them, the performance and durability of polymer electrolyte membrane is a crit ical one. Perfluorosulfonic-acid (PFSA) membranes are frequently used as the ion-conductive membrane for their excellent ion conductivity and chemical-mechanical stability [1]. In PEFCs, the PFSA membrane acts as an electrolyte to transport proton and meanwhile as a separator to prevent reactant gases from crossing over. The multifunction roles of the PFSA membrane require membranes with high conductivity to ensure fuel cells performance and with high structural integrity to ensure fuel cells durability. Therefore, high conductivity and structural integrity are two desired properties for PFSA membranes. To obtain membranes with high conductivity, several attempts have been made. One way is to reduce the equivalent weight (EW) of the membrane through increasing the number of sulfonic acid groups. However, the sulfonic acid groups are vulnerable to radical attack, and membranes with low EW tend to display low chemical durability [2]. Another way is to reduce the thickness of the membrane in order to minimize ohmic losses. However, accompanying reduced thickness is the decreased gas transport resistance, leading to more crossover of reactant gases and consequently chemical degradation [3,4]. Therefore, chemical degradation of PFSA membranes is an inevitable problem encountered during the development of ion-conductive membranes to improve the performance of the membrane. The durability of mem branes is not only impacted by chemical degradation but also by me chanical degradation, in the form of creeps, pinholes, cracks and so on [5–8]. In most cases, these two degradation mechanisms couple, resulting in more severe membrane degradation [9]. Chemical degra dation decomposes membrane and produces defects which grow under cyclic mechanical loading. The enlarged defects result in more hydrogen crossover and aggravate chemical degradation. Till now, the mecha nisms of chemical degradation are well understood. Hydroxyl radical (�OH) resulting from the decomposition of hydrogen peroxide is thought as the main cause. The highly reactive radical species attack the mem brane through main chain unzipping and side chain scission mechanisms [5,10–13]. The extent of chemical degradation is more pronounced under higher temperature and lower humidity environment. With respect to mechanical degradation, creeps are more detrimental and more likely to occur at higher temperatures [8,14,15]. In addition, the humidity cycling promotes crack propagation and delamination under cyclic stress [16–20]. The combined chemical and mechanical degradations, as well as isolated chemical or mechanical degradation on the lifetime of fuel cells have been extensively investigated through accelerated stress testing (AST) [12,21–24]. To uncouple the interaction between chemical and mechanical degradations, it is desired to clarify how one degradation mode influences the other one. In terms of the effect of mechanical stress on chemical degradation, it was found that the chemical degradation rate was promoted by mechanical stress, irrespective of tensile or compressive stress [25,26]. Nevertheless, how chemical degradation affects mechanical properties is scarcely investigated with a few notable exceptions. Huang et al. [27] investigated the stress-strain responses of 48 h chemical degraded membrane and found that the strain-to-failure of the degraded membrane was significantly reduced. In addition, the failure mode of the degraded membrane was found to be fast brittle fracture. Similar observations were also reported by Alavijeh et al. [28, 29] and Lim et al. [30] accompanied with mild increase in elastic modulus of catalyst coated membrane after combine chemical and me chanical degradations. Combing with expansion tests, they concluded that in-situ hygrothermal cycles were sufficient to lead to membrane fracture. A reduction in modulus and ductility was observed by Patil et al. [31] after degrading under open circuit voltage (OCV) conditions.
2. Experimental 2.1. Materials and preparation The Nafion® 212 membranes manufactured by Dupont with the nominal thickness of 50 μm were used in as received state (Hþ form). Membranes were cut into rectangular shape with the length and width of 50 mm and 10 mm, respectively. FeSO4⋅7H2O, NaCl (99.5%, analytical reagent) and H2O2 (30 wt%) were purchased from Tianjin Yuanli Chemical Company (Tianjin, China). 2.2. Membrane morphological changes 2.2.1. Chemical degradations The membranes were chemically degraded using Fenton reagent solutions. Firstly, the Fenton reagent solution was prepared using a mixture of 0.001 mol L 1 Fe2þ solution and 30% H2O2 solution at a volume ratio of 1:3. Next, membranes were placed into the Fenton re agent solution and pretreated in a water bath environment at 80 � C for a certain period of time. In order to obtain membranes with different chemical degradation levels, membranes were pretreated in Fenton re agent solutions for 24, 36, 48 and 72 h. The as-received specimens were 2
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denoted as 0 h. Pretreated membranes were washed in deionized water for several times to remove excessive reagents and dried in oven at 60 � C. During chemical degradation process, the solutions were changed every 24 h where the fluoride concentration was measured using PXS270 ion meter (Leici Instrument, Shanghai, China) and PHS-25 pH meter (Leici Instrument, Shanghai, China). The fluoride ion meter was calibrated using 10 2 mol L 1, 10 3 mol L 1, 10 4 mol L 1, 10 5 mol L 1 standard solutions with the same amount of Total Ionic Strength Adjustment Buffer (TISAB). The total fluoride content was determined from the measured fluoride concentration and solution volume.
The double edge notch tension tests were performed in order to determine the fracture behavior of membranes after chemical degra dation. Rectangular specimens of the same size as uniaxial tensile tests were used. Prior to tests, two notches of the same size were cut using fresh razor blades at the edges of the specimen along the width direction, leaving the remaining part as the ligament area. For each type of chemical degraded specimen, multiple specimens with different liga ment lengths ranging from 1.5 mm to 7.5 mm were used. As a result, load-displacement curves at different ligament lengths were obtained, which can be used to determine the fracture resistance of chemical degraded membranes using the essential work of fracture (EWF) method [39]. The EWF method assumes that the crack-tip area consists of two parts: a plastic zone where plastic deformation occurs and a fracture process zone where necking/tearing takes place. Consequently, the en ergy for crack growth can be divided into two parts as follows:
2.2.2. Swelling Membranes with different chemical degradation levels were immersed in 1 M H2SO4 solution for 24 h to remove any possible remaining Fe ions, followed by placing into liquid water at room tem perature (23 � C), and the dimensional change of the membrane in length direction was recorded. The initial dimensions of the membrane were determined as soon as the membranes were taken out from the oven. The extent of swelling ΔL/L0 was calculated as the dimensional change of the membrane in liquid water ΔL, divided by its dry state L0. To ensure the reproducibility of the measurements, each measurement was carried out for three times and the average value was used. To set a baseline, the swelling of an as-received membrane that went through the same ther mal history in liquid water (preheated at 80 � C) was also measured, denoted as PH.
Wf ¼ We þ Wp
(1)
where Wf is the total work of fracture which can be obtained from the area encompassed by the load-displacement curve. We is the essential work of fracture which represents the energy consumed in the inner fracture process zone to create new crack surfaces and is proportional to the area of the fracture process zone. Whereas Wp is the non-essential work of fracture, corresponding to the energy dissipated in the outer plastic zone and is proportional to the volume of the outer plastic zone. Consequently, Equation (1) can be rewritten as: (2)
2.3. Characterization
wf ¼ we þ βwp ⋅L
2.3.1. FTIR Fourier transformer Infrared (FTIR) spectra were collected on an IRAffinity-1S spectrometer (Shimadzu Corp., Japan) within the range of 4000–400 cm 1 with a wavenumber resolution of 2 cm 1. The mem brane samples had a standard size of 5 cm � 1 cm, which covered the whole window to prevent interference. Samples were settled between two KBr plates and placed in the cell for measurement. The spectra were obtained against the air background spectrum.
where wf is the specific work of fracture and we is the specific essential work of fracture. β is a shape factor that reflects the shape of the outer plastic zone. By plotting specific work of fracture as a function of liga ment length L, the specific essential work of fracture, we, and βwp can be obtained from the intercept and slope of the curve, respectively. The loading control mode and strain rate of DENT were the same as uniaxial tensile tests. A more detailed description of the experimental procedure can be found in our previous study [40–42].
2.3.2. In-plane conductivity The in-plane conductivity of degraded membranes were measured using a four-electrode BT-110 conductivity clamp (Scribner Associates). A detailed description can be found in Refs. [36–38]. Prior to tests, samples were cut into rectangular stripes and immersed in 1 M H2SO4 solution for 24 h to remove any possible remaining Fe ions, followed by rinsing in deionized water for several times. The AC impedance was measured in liquid water at room temperature (23 � C) using a Princeton P4000 potentiostat from 1 Hz to 1 MHz. The in-plane conductivity was calculated as κ ¼ LIP/RIPAIP, where LIP is the distance between elec trodes, RIP is the measured resistance, and AIP is the cross-section area of the membrane. Each measurement was repeated at least twice to ensure the reproducibility.
2.3.4. Scanning electron microscopy (SEM) Physical changes at the surfaces and cross-sections of Nafion® 212 membranes before and after chemical degradation were observed using scanning electron microscopy (SU1510, Hitachi, Japan). The crosssections of Nafion® 212 membranes were prepared by breaking mem branes in liquid nitrogen to obtain a smooth cross-section. Samples were then stuck on conductive tapes and gold sputtered under vacuum before SEM observation. Different regions of DENT specimens after fracture tests were captured and compared to illustrate how chemical degrada tion changed the fracture properties of membranes. 3. Results and discussion 3.1. Morphologies of Nafion 212 membrane after chemical degradation
2.3.3. Tensile and double edge notch tension tests Two kinds of mechanical tests were conducted: uniaxial tensile tests and double edge notch tension (DENT) tests. All experiments were carried out on an in-situ tensile testing machine (IPBF-300, CARE Mea surement & Control Co., Ltd., China) at room condition (25 � C/50% RH). The uniaxial tensile tests were performed to determine the stressstrain curves of chemical degraded membranes and basic tensile me chanical properties, such as Young’s modulus, yield stress, strainhardening modulus and break strain. Rectangular specimens were used with the gauge length and width of 30 mm and 10 mm, respec tively. The uniaxial tensile tests were performed under strain control mode with the strain rate of 0.003 s 1. While the stress was recorded with a load sensor, the strain was measured with a non-contact displacement detecting system (NDDS) to avoid displacement by the machine itself.
The morphologies of Nafion 212 membrane after chemical degra dation are shown in Fig. 1. The morphologies of as-received specimens are smooth without any defects. After chemical degradation for 24 h, no discernible changes can be detected. However, increasing chemical degradation duration to 36 h results in bubbles appearing on the surface and pores appearing in the cross section. Similar observations were also made by Tang et al. [43] who ascribed the formation of bubbles to the decomposition of the repeating units of the membrane. Further increasing chemical degradation duration to 48 h increases the density of bubbles, in both surface and cross section areas. In addition, tears of the bubbles and bumps are also observed on the surface. However, with the prolong of chemical degradation time to 72 h, the bubbles become smaller but distribute more densely. Such a phenomenon can be ascribed to the emergence of new bubbles from the surfaces of old large bubbles. 3
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Fig. 1. Surface and cross-section morphologies of Nafion® 212 membrane after chemical degradation for (a) 0 h, (b) 24 h, (c) 36 h, (d) 48 h and (e) 72 h. The crosssection morphologies are shown in the inset at the upper right corner of each image. (f) is the percentage of bubble for membranes with different degradation levels.
It is easily to tell from the cross section in Fig. 1(e) that the membrane is more severely degraded as the surface becomes more roughened and uneven. In addition, with the increase in degradation time, the mem brane’s color changes from transparent (0 h) to opaque (36 h). To characterize the degree of chemical degradation, the percentage of bubbles for membranes with different degradation time is compared in Fig. 1(f). The density of bubbles generally increases with chemical degradation time except for 72 h.
To decide if the extent of ex-situ chemical degradation in this study is comparable to in-situ accelerated stress tests, the total fluoride release is also shown for comparison in the inset in Fig. 2. It can be found that the fluoride release rate in this study are similar with that tested under simulated fuel cell conditions. The fluoride emission rate (FER) of an operating fuel cell was measured intermittently by periodically col lecting effluent water from the anode and cathode gas outlets, which was then analyzed following standard test protocols to calculate the concentration of fluoride ions [23,29,44]. It should be noted that the total fluoride release alone does not always sufficiently enough to assess the health condition of fuel cells, as some degradation regions are highly localized [45–47]. Regardless of its shortcomings, the fluoride release could still be used for rough estimation due to its convenience, as the fluoride release can be easily measured from the products of fuel cells.
3.2. Fluoride release The extent of chemical degradation of membranes relies on many factors, such as hydrogen peroxide concentration, pretreated tempera ture and degradation time. To compare the mechanical properties of membranes with different degradation levels, a parameter is needed to represent the extent of degradation. As such, the fluoride release is used as an indicator. The total fluoride release as a function of degradation time is shown in Fig. 2. The total fluoride release increases almost lin early with degradation time, indicating a constant fluoride release rate.
3.3. Molecular structural changes of the membrane To characterize molecular structural changes induced by chemical degradation, FTIR investigation of the membrane was carried out. The obtained FTIR spectra is normalized to the CF2 stretching peak at 1145 cm 1, as shown in Fig. 3. According to the literature [48–50], different absorption peaks correspond to different assignments. The C–F
Fig. 2. Total amount of fluoride release as a function of degradation time. To compare the ex-situ measurement in this study with in-situ accelerated stress tests, data from Patil et al. [31] and Lim et al. [30] are also shown for comparison.
Fig. 3. FTIR spectra of Nafion® 212 membrane with different chemical degradation levels. 4
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absorption peaks at 1145 and 1200 cm 1 as well as the C–C absorption peak at 1300 cm 1 are characteristics of the PTFE backbone of the membrane. While the absorption peak at 970 cm 1 is assigned to the symmetric C–O–C stretching, the peak at 1057 cm 1 is assigned to the symmetric S–O stretching. These two peaks correspond to chemical bonds at the pendant side chain and terminated sulfonic groups of Nafion membrane, respectively. With the increase of chemical degra dation time, little variation in normalized C–F peak intensity is observed. By contrast, the heights of the C–O–C and S–O normalized peaks decrease significantly, suggesting that the sulfonic acid groups and the ether linkage at the pendant side chain are more vulnerable to radial attacks. Even though the molecular structure changes of Nafion membrane after chemical degradation have been investigated by many researchers [43,49,51–55], it is not trivial to confirm the molecular structure in this paper as different mechanisms are reported in the literature, and me chanical responses of the membrane is strongly related to its molecular structure. While some researchers concluded that chemical degradation took place more preferentially from the side chains [49,51,55], some researchers found that membrane decomposition started from the ends of main chain [43,54], both main chain and side chain [53], as well as not changed by degradation [52]. One possible reason might be due to the reduced –COOH end groups in new generation chemically stabilized PFSA membranes where the chain scission mechanism is more prevalent [56]. PFSA membranes have hydrophobic PTFE main chains and hy drophilic side chains terminated with sulfonic acid ionic groups, creating a nanophase-separated morphology. Because of the molecular structure changes, the morphology of the membrane also changes accordingly. According to the small-angle X-ray scattering (SAXS) re sults by Kusoglu et al. [26], the domain spacing of degraded membrane increased with increasing decomposition level. A relatively large non-ionic region which resembled void in the membrane was observed by Venkatesan et al. [57] after chemical degradation through trans mission electron microscopy (TEM) observation. Depending on the temperature and relative humidity, mechanical loads can be transmitted through hydrophobic or hydrophilic domains [58]. Therefore, it is anticipated that the mechanical properties will change, which will be investigated later.
uptake which is intimately related to the ion exchange capacity (IEC) of the membrane. As discussed in Fig. 3, the sulfonic group at the end of side chain is attacked by radicals, reducing the water uptake capacity of the membrane. Decreased water uptake after chemical degradation was reported [49,51], which was ascribed to morphological changes. The reduced swelling is beneficial from the mechanical perspective in that the in-plane stress can be reduced, improving the mechanical durability. Patil et al. [31] compared the contractile stress before and after chemical degradation, and it was found that the peak stress of degraded mem branes was lower than that of fresh samples. However, accompanying chemical degradation is the entanglement changes that might reduce the resistance for the membrane to withstand the resulting stresses. In their study, even though the stress of degraded membrane was lower, the degraded membranes broke into two pieces, while the fresh membrane maintained structural integrity. Thus, the mechanical properties, espe cially the fracture behavior of the membrane, need to be evaluated in order to have a full evaluation of membrane durability under combined chemical and mechanical degradations. To have a better understanding of the degradation level, especially how the performance of the membrane is impact by chemical degra dation, the in-plane conductivity of different degraded membranes are compared in Fig. 4(b). The conductivity decreases with increasing degradation time, similar with the trend in swelling. After chemical degradation for 24 h, the conductivity slightly decreases, indicating that short-term chemical degradation have minor impact on the conductiv ity. The similar conductivity value between 24 h-degraded membrane and 0 h-PT membrane suggests that the Fe ion is exchanged completely. However, further increasing degradation time to 36 h causes 40% decrease in conductivity, reducing from 0.083 S/cm for PH membrane to 0.0497 S/cm for 36 h degraded membrane. The significant reduction in conductivity after degradation for 36 h appears concurrently with the changes in surface morphology in Fig. 1, indicating that a certain amount of time is needed to induce decay in properties of the membrane. Further increasing degradation time from 36 h to 48 h and 72 h de creases the conductivity from 0.0497 S/cm to 0.0406 S/cm for 48 h degraded membrane and to 0.0378 S/cm for 72 h degraded membrane. Even though the conductivity decreases for more than 50% after degradation for 72 h, the membrane can still function well. The reduced conductivity is ascribed to the possible loss of sulfonic groups at the end of side chains, as indicated from the FTIR study in Fig. 3.
3.4. Swelling and conductivity The dimensional stability of the membrane is an important factor affecting the mechanical durability. Excessive swelling results in significantly in-situ residual stress, impairing the mechanical integrity of the membrane. Therefore, the extent of swelling is firstly evaluated before mechanical tests as shown in Fig. 4(a). With the increase of degradation level, the swelling decreases from 14.4% for 0 h-PH mem brane to 14.1% for 24 h degraded membrane, finally to 11.6% for 72 h degraded membrane. Decreased hygral expansion after in-situ acceler ated stress test degradation was also observed by Alavijeh et al. [28]. Such a reduction in swelling can be ascribed to the decreased water
3.5. Tensile properties The stress-strain curves of membranes degraded for different time are compared in Fig. 5 (a). The AsR membrane is ductile with high tensile strength. After chemical degradation for 24 h, the tensile strength does not change much while the proportional limit stress decreases significantly. For PFSA membranes, the proportional limit stress is traditionally used to represent the yield strength of the membrane [42, 59,60]. Increasing degradation time to 36 h results in significant reduced ductility. The break strain decreases from 265% for 24 h
Fig. 4. Comparison of (a) swelling and (b) conductivity of membranes with different degradation levels. 5
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Fig. 5. (a) Stress-strain curves of membranes degraded for different time; (b)–(e) are comparison of Young’s modulus, proportional limit stress, strain-hardening modulus and break strain, respectively.
degraded membrane to 82% for 36 h degraded membrane. The changes of break strain with degradation time is in accordance with the morphology changes in Fig. 1 where bubbles start to emerge on 36 h-degraded membrane. The obtained mechanical results could be due to morphological changes, as the presence of bubbles is a macroscopic reflection of microstructural and morphological changes. The reduction in ductility at an early stage was also observed during accelerated stress test where the ductility diminished after only two AST cycles [42]. Further increasing degradation time to 48 h and 72 h leads to gradually decrease of ductility. It seems that there is a critical point where the ductility starts to drop significantly. As discussed in section 3.1, non-ionic regions and voids will form after chemical degradation. These defects cause locally stress concentration and severe deformation to final failure. To better illustrate how mechanical properties changes with degra dation time, Young’s modulus, proportional limit stress, strainhardening modulus and break strain are extracted from stress-strain curves and compared in Fig. 5 (b)–(e). With increasing degradation level, the Young’s modulus, proportional limit stress and break strain decrease, while the strain-hardening modulus increases. In particular, significant changes are observed for Young’s modulus and break strain after chemical degradation for 36 h. Deformation at different stages have different morphology origins. While small strain deformation corresponds to chain rotation of bundles, large strain deformation cor responds to alignment of aggregates by disentangling from each other [61]. Based on these morphology assignments and molecular origins, the deformation mechanisms can be tentatively obtained. The degraded ionic groups at the terminal of side chains reduced the electrostatic interaction between side chains, thus lowering the restriction imposed on segmental motions and making it easier for chain rotation to occur. While changes in Young’s modulus after chemical degradation have been reported by some researchers [28,31], the strain-hardening prop erties have never been investigated. Patil et al. [31] thought that the molecular weight would decrease after chemical degradation and chain slippage would occur more easily as the shortened chains induced by chemical degradation would be less likely to form entanglements. However, the results in this study are not coincide with the above findings, which suggest that a more sophisticated mechanism might exist and more factors should be taken into consideration. One possible reason might be the morphological rearrangement. For instance, the domain spacing increases with increasing decomposition level [26]. The reduced ductility of degraded membranes increases the risk of me chanical failure as the membrane is unable to accommodate large deformation. Consequently, cracks and pinholes are more likely to form during humidity cyclings, leading to crossover of reactant gases. This
also explain why the life time under combine chemical/mechanical accelerated stress tests are much shorter than in isolated chemical or mechanical accelerated stress tests [9,21–23]. The reduced ductility alone is not sufficient to understand the mechanical failure mechanisms of the membrane. To address this issue, the fracture process and mechanisms also need to be considered. 3.6. Fracture behavior and mechanisms The above tensile properties describe the ability of the membrane to resist deformation and crack formation. In the presence of cracks, the fracture behavior is needed to assess the ability of the membrane to resist crack growth. To this end, the load-displacement curves of mem branes with different degradation levels are compared, as shown in Fig. 6. For AsR membranes in Fig. 6 (a), the load increases to the maximum point and gradually drops. The maximum load point corre sponds to crack initiation, after which crack starts to propagate [41,62]. The gradual decrease of load suggests that the crack propagates in a moderate way, showing a ductile fracture feature. However, with the increase of degradation level, such a feature diminishes. The load starts to drop abruptly for 36 h degraded membrane, showing a brittle fracture feature. Such a feature is more pronounced for 72 h degraded membrane where the load almost drops to zero as soon as reaching to the maximum load. Therefore, the membrane experiences a transition from ductile fracture to brittle fracture with increasing degradation level. To quantitatively assess how the crack initiation energy and propa gation energy change in response to chemical degradation, the essential work of fracture method is used as shown in Fig. 7. The intercept and slope are extracted and compared in Fig. 7(b)–(c), respectively. With the increase of degradation level, both the slope and intercept decrease. The intercept represents the energy needed for crack initiation, while the slope represents the ability of the membrane to resist crack propagation. The decreased intercept and slope with increasing degradation level suggests that the crack is not only easy to initiate but also more easily to propagate. The fracture initiation energy decreases by more than 40%, from 8.7 kJ m 2 for AsR membrane to 5 kJ m 2 for 48 h degraded membrane. A notable exception is the 72 h degraded membrane with the fracture initiation energy as high as 9 kJ m 2. The origin of such a discrepancy is not clear at the present stage, which could be related to microstructure and morphological changes. Nevertheless, this phe nomenon along with the increased strain-hardening modulus in Fig. 5 imply that the mechanical behavior are not as simple as we had spec ulated. With respect to the slope, significant decrease in slope (more than 70%) is observed, from 4.2 MJ m 3 for AsR membrane to 1 MJ m 3 for 72 h degraded membrane. As the slope is related to the outer plastic 6
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Fig. 6. Load-displacement curves at different ligament lengths for (a) AsR membrane; (b) 24 h degraded membrane; (c) 36 h degraded membrane; (d) 48 h degraded membrane; (e) 72 h degraded membrane.
Fig. 7. (a) Specific work of fracture as a function of ligament lengths for different degraded membranes; (b)–(c) are comparison of specific essential work of fracture, we and βwp, respectively.
dissipation zone where crazing is more likely to occur [39], the reduced non-ionic regions observed by Venkatesan et al. [57] are expected to promote craze formation. The findings in this study suggest that the detrimental effect of chemical degradation not only manifests as the reduced crack initiation resistance, but more importantly the crack propagation resistance is largely undermined. A fuel cell will not fail immediately after a crack forms. Instead, it will continue operating for a quiet long period of time, as water generated during operation will cover the surface of cracks, preventing crossover of reactants [63]. Thus, crack propagation resis tance is an important factor affecting the lifetime of fuel cells [16,17, 42]. Therefore, when designing new membranes, the crack propagation resistance as well as its sensitivity to chemical degradation should be
evaluated to achieve a high lifetime. To clarify the mechanisms of crack propagation process, the crosssections of cracked degraded specimens are examined, as shown in Fig. 8 (a). The AsR specimen is taken as an example for illustration. Several areas with different thicknesses are observed from which the crack propagation process can be identified. The region with the thickness of 50 μm is assigned to the pre-crack region which is cut by razor blades and the thickness is not affected. Right next to the pre-crack region is the stretch zone which represents the plastic zone size ahead of crack tip. The thickness of this region is 22 μm, which is less than 50% of the original thickness. Such a reduction in thickness is induced by severe plastic deformation. A similar reduction in thickness of the membrane is also observed during the fatigue crack propagation [19]. The shape of 7
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Fig. 8. (a) An illustration of the cross section of AsR Nafion membrane with notch zone, stretch zone and fracture zone. The crack propagates from left to right. (b) Comparison of stretch zone width of membranes with different degradation levels.
the stretch zone is semi-elliptical, and the region near the crack surface is severely deformed while that far from the crack surface is less affected. In addition, severely deformed strips can be also observed on the cross-section area. These features manifest that severe plastic deforma tion occurs in the stretch zone. At the right side of this image is the fracture region with the thickness reduced to 28 μm. This region is assigned to fast crack propagation where instability crack growth oc curs. Thus, the stretch zone represents the plastic zone ahead of crack tip. To quantify how the plastic zone ahead of crack tip evolves with degradation level, the stretch zone width which is the width between the notch region and fracture region, is compared for different degraded membranes in Fig. 8 (b). Similar with the trend of fracture resistance, the stretch zone width decreases with the degradation level. In other words, the size of the plastic zone ahead of crack tip decreases with degradation level. As a result, the energy dissipated through the plastic zone is reduced, accounting for the decreased slope in Fig. 7. In this way, chemical degradation affects crack propagation through reducing the plastic zone size of crack tip and exhibiting a ductile-to-brittle transition feature.
propagation resistance is evaluated according to the fracture behavior of the membrane based on fracture mechanics. After degradation, the loaddisplacement curves change from ductile fracture to brittle fracture behavior. Using the essential work of fracture method, both the crack initiation and propagation energies are found to decrease with degra dation level, with an exception of the 72 h degraded membrane. Based on the microstructure assignment of crack propagation, the crosssections of degraded membranes after double edge notch tension tests are examined. The stretch zone width, which represents the plastic zone size ahead of crack tip, is found to decrease with degradation level. Thus, the decreased crack propagation resistance after chemical degradation is ascribed to the reduced plastic zone size ahead of crack tip, which is induced as a result of membrane embrittlement. Therefore, chemical degradation not only impairs tensile mechanical properties, but also undermines crack propagation resistance through minimizing crack tip plastic zone size and thus reducing energy dissipation. The results in this study are helpful for the understanding of mechanical degradation and fracture failure induced by chemical degradation, which will provide more insights into the durability of membranes and develop effective mitigation strategies to enhance membrane durability.
4. Conclusions
Declaration of competing interest
In this study, the mechanical properties after chemical degradation, especially the fracture behavior, are investigated. To characterize the extent of degradation and link mechanical properties with degradation level, the amount of fluoride release is measured as a function of degradation time. It is found that the total amount of fluoride release increases almost linearly with degradation time. More importantly, the fluoride release value is similar with that obtained from in-situ acceler ated stress tests, indicating that the mechanical properties obtained here could, to some extent, represent the actual material state of membrane in operating fuel cells. The FTIR is also carried out to probe the mo lecular origin of chemical degradation in order to account for changes in mechanical properties induced by chemical degradation from a molec ular perspective. The changes in FTIR spectra suggest a predominant loss of sulfonic acid groups and the ether linkage from side chains. As the mechanical stress is generated due to sorption and desorption of water, the swelling properties of degraded membranes are also measured to characterize the potential induced stresses during humidity cyclings. It is found that the swelling behavior decreases with degradation time, which is ascribed to the reduced sulfonic acid groups and therefore reduced water uptake. Accompanying with reduced swelling behavior is the decreased conductivity which shows a 50% reduction after degra dation for 72 h. With respect to mechanical properties, increasing degradation level reduces the Young’s modulus, proportional limit stress and break strain, while the strain-hardening modulus is increased, implying the embrit tlement of the membrane after chemical degradation. These properties only accounts for factors leading to crack initiation. The crack
The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement Xiaoyi Sun: Formal analysis, Investigation, Methodology, Visuali zation, Writing - original draft. Shouwen Shi: Conceptualization, Funding acquisition, Methodology, Project administration, Supervision, Visualization, Writing - review & editing. Yuanjie Fu: Investigation, Methodology. Jian Chen: Methodology, Project administration. Qiang Lin: Investigation, Methodology, Project administration. Jiaqi Hu: Investigation, Methodology. Cong Li: Methodology, Project adminis tration. Jiayao Li: Investigation, Methodology. Xu Chen: Conceptuali zation, Supervision, Writing - review & editing. Acknowledgements The authors gratefully acknowledge the financial support from the National Natural Science Foundation of China (No.51805364, No.11902216), China Postdoctoral Science Foundation (No. 2019TQ0225, No.2019M661017). The research is also partial finan cially supported by the Open Research Fund of Science and Technology Innovation Platform of Key Laboratory of Efficient & Clean Energy Utilization, Changsha University of Science & Technology (No. 2018NGQ002). 8
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Perspective
Embrittlement induced fracture behavior and mechanisms of perfluorosulfonic-acid membranes after chemical degradation Xiaoyi Sun a, Shouwen Shi a, *, Yuanjie Fu a, Jian Chen b, Qiang Lin a, Jiaqi Hu a, Cong Li b, Jiayao Li a, Xu Chen a a
School of Chemical Engineering and Technology, Tianjin University, Tianjin, 300072, China Key Laboratory of Efficient & Clean Energy Utilization, The Education Department of Hunan Province, School of Energy and Power Engineering, Changsha University of Science and Technology, Changsha, 410114, China
b
H I G H L I G H T S
G R A P H I C A L A B S T R A C T
� Fracture behavior of Nafion membranes are explored after chemical degradation. � Chemical degradation causes predomi nant loss of side chains. � Crack transforms from ductile to brittle fracture with increasing degradation level. � Stretch zone size from cross-section observation is related to crack propagation. � Reduced crack propagation resistance is ascribed to decreased plastic zone size.
A R T I C L E I N F O
A B S T R A C T
Keywords: PFSA membrane Fracture Chemical degradation Mechanical property Durability
Chemical and mechanical degradations of perfluorosulfonic-acid membranes are two factors contributing to the reduced durability of fuel cells. While the mechanisms of isolated chemical or mechanical degradation are extensively investigated, the impact of chemical degradation on mechanical degradation is not fully understood. In this paper, the fracture behavior of Nafion 212 membranes with different chemical degradation levels are investigated. To characterize the degradation level and probe the molecular origins of chemical degradation, fluoride release, Fourier-transformer infrared spectra, conductivity and swelling behavior are measured. It is found that chemical degradation causes predominant loss of side-chains. A transition from ductile to brittle fracture is observed with increasing degradation level. In addition, the cross-sections are examined to link fracture behavior with microstructure changes. While the decreased fracture toughness is attributed to reduced mechanical properties such as Young’s modulus and break strain, the decreased crack propagation resistance is ascribed to reduced plastic zone size ahead of crack tip, which reduces the plastic energy dissipation. These findings not only provide new mechanical dataset of degraded membranes for performance and durability modelling to take into account of chemical degradation effect but also improve the understanding of membrane durability.
* Corresponding author. E-mail address: [email protected] (S. Shi). https://doi.org/10.1016/j.jpowsour.2020.227893 Received 1 November 2019; Received in revised form 9 February 2020; Accepted 11 February 2020 0378-7753/© 2020 Elsevier B.V. All rights reserved.
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1. Introduction
They observed cracks on the surface of degraded samples which they attributed to reduced ductility. Even though the tensile properties of chemical degraded membranes are investigated, only basic properties are obtained from stress-strain curves, such as elastic modulus and strain-to-failure, with the fracture behavior and underlying mechanisms left unexplored. The tensile properties alone are not sufficient enough to account for the reduced durability after chemical degradation. For instance, Rodgers et al. [2] found that even though 950 EW cell had the largest losses in mechanical strength, the crossover is not significantly impacted. And they concluded that there was no correlation between mechanical strength and cross over of PFSA membrane. As gas crossover is not only related to me chanical strength, but also related to crack initiation and propagation resistance. Therefore, it is imperative to investigate the fracture behavior of PFSA membrane after chemical degradation in order to alleviate the failure issue and develop effective mitigation strategies to enhance membrane durability. In addition, most numerical simulations on mechanical responses of PFSA membranes used mechanical proper ties of as-received (AsR) membranes without considering decay in me chanical properties with chemical degradation [32–35]. Thus, the mechanical properties after chemical degradation should be systemati cally investigated, which could be used as inputs in numerical modelling to update material parameters to account for chemical degradation is sues and accurately characterize the interaction between durability with performance. Hence, the objective of this paper is to investigate the mechanical properties of Nafion 212 membrane after chemical degradation, espe cially the fracture behavior, with the aim of elucidating how chemical degradation changes the mechanical properties as well as the underlying mechanisms. To link mechanical properties with degradation levels, the total fluoride release is measured and used as an indicator for degra dation level. In addition, molecular structure changes induced by chemical degradation are also explored to account for molecular origins of mechanical changes. As stresses are generated in-situ due to swelling and shrinkage of the membrane during humidity cyclings, the swelling behavior of the membrane is also investigated to understand the dimensional stability. Finally, the tensile and fracture behaviors of membranes at different degradation levels are compared, and the crosssections of fractured membranes are examined to correlate with fracture resistances. The findings in this paper are expected to provide more insights into the decay in mechanical properties after chemical degra dation as well as the understanding of the failure mechanisms of membranes under combined chemical and mechanical degradations.
Polymer-electrolyte fuel cells (PEFCs) are promising zero-emission power in many fields such as transportation and electricity. However, the performance and durability issues are two challenges impeding the commercialization of this promising technology. Among them, the performance and durability of polymer electrolyte membrane is a crit ical one. Perfluorosulfonic-acid (PFSA) membranes are frequently used as the ion-conductive membrane for their excellent ion conductivity and chemical-mechanical stability [1]. In PEFCs, the PFSA membrane acts as an electrolyte to transport proton and meanwhile as a separator to prevent reactant gases from crossing over. The multifunction roles of the PFSA membrane require membranes with high conductivity to ensure fuel cells performance and with high structural integrity to ensure fuel cells durability. Therefore, high conductivity and structural integrity are two desired properties for PFSA membranes. To obtain membranes with high conductivity, several attempts have been made. One way is to reduce the equivalent weight (EW) of the membrane through increasing the number of sulfonic acid groups. However, the sulfonic acid groups are vulnerable to radical attack, and membranes with low EW tend to display low chemical durability [2]. Another way is to reduce the thickness of the membrane in order to minimize ohmic losses. However, accompanying reduced thickness is the decreased gas transport resistance, leading to more crossover of reactant gases and consequently chemical degradation [3,4]. Therefore, chemical degradation of PFSA membranes is an inevitable problem encountered during the development of ion-conductive membranes to improve the performance of the membrane. The durability of mem branes is not only impacted by chemical degradation but also by me chanical degradation, in the form of creeps, pinholes, cracks and so on [5–8]. In most cases, these two degradation mechanisms couple, resulting in more severe membrane degradation [9]. Chemical degra dation decomposes membrane and produces defects which grow under cyclic mechanical loading. The enlarged defects result in more hydrogen crossover and aggravate chemical degradation. Till now, the mecha nisms of chemical degradation are well understood. Hydroxyl radical (�OH) resulting from the decomposition of hydrogen peroxide is thought as the main cause. The highly reactive radical species attack the mem brane through main chain unzipping and side chain scission mechanisms [5,10–13]. The extent of chemical degradation is more pronounced under higher temperature and lower humidity environment. With respect to mechanical degradation, creeps are more detrimental and more likely to occur at higher temperatures [8,14,15]. In addition, the humidity cycling promotes crack propagation and delamination under cyclic stress [16–20]. The combined chemical and mechanical degradations, as well as isolated chemical or mechanical degradation on the lifetime of fuel cells have been extensively investigated through accelerated stress testing (AST) [12,21–24]. To uncouple the interaction between chemical and mechanical degradations, it is desired to clarify how one degradation mode influences the other one. In terms of the effect of mechanical stress on chemical degradation, it was found that the chemical degradation rate was promoted by mechanical stress, irrespective of tensile or compressive stress [25,26]. Nevertheless, how chemical degradation affects mechanical properties is scarcely investigated with a few notable exceptions. Huang et al. [27] investigated the stress-strain responses of 48 h chemical degraded membrane and found that the strain-to-failure of the degraded membrane was significantly reduced. In addition, the failure mode of the degraded membrane was found to be fast brittle fracture. Similar observations were also reported by Alavijeh et al. [28, 29] and Lim et al. [30] accompanied with mild increase in elastic modulus of catalyst coated membrane after combine chemical and me chanical degradations. Combing with expansion tests, they concluded that in-situ hygrothermal cycles were sufficient to lead to membrane fracture. A reduction in modulus and ductility was observed by Patil et al. [31] after degrading under open circuit voltage (OCV) conditions.
2. Experimental 2.1. Materials and preparation The Nafion® 212 membranes manufactured by Dupont with the nominal thickness of 50 μm were used in as received state (Hþ form). Membranes were cut into rectangular shape with the length and width of 50 mm and 10 mm, respectively. FeSO4⋅7H2O, NaCl (99.5%, analytical reagent) and H2O2 (30 wt%) were purchased from Tianjin Yuanli Chemical Company (Tianjin, China). 2.2. Membrane morphological changes 2.2.1. Chemical degradations The membranes were chemically degraded using Fenton reagent solutions. Firstly, the Fenton reagent solution was prepared using a mixture of 0.001 mol L 1 Fe2þ solution and 30% H2O2 solution at a volume ratio of 1:3. Next, membranes were placed into the Fenton re agent solution and pretreated in a water bath environment at 80 � C for a certain period of time. In order to obtain membranes with different chemical degradation levels, membranes were pretreated in Fenton re agent solutions for 24, 36, 48 and 72 h. The as-received specimens were 2
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denoted as 0 h. Pretreated membranes were washed in deionized water for several times to remove excessive reagents and dried in oven at 60 � C. During chemical degradation process, the solutions were changed every 24 h where the fluoride concentration was measured using PXS270 ion meter (Leici Instrument, Shanghai, China) and PHS-25 pH meter (Leici Instrument, Shanghai, China). The fluoride ion meter was calibrated using 10 2 mol L 1, 10 3 mol L 1, 10 4 mol L 1, 10 5 mol L 1 standard solutions with the same amount of Total Ionic Strength Adjustment Buffer (TISAB). The total fluoride content was determined from the measured fluoride concentration and solution volume.
The double edge notch tension tests were performed in order to determine the fracture behavior of membranes after chemical degra dation. Rectangular specimens of the same size as uniaxial tensile tests were used. Prior to tests, two notches of the same size were cut using fresh razor blades at the edges of the specimen along the width direction, leaving the remaining part as the ligament area. For each type of chemical degraded specimen, multiple specimens with different liga ment lengths ranging from 1.5 mm to 7.5 mm were used. As a result, load-displacement curves at different ligament lengths were obtained, which can be used to determine the fracture resistance of chemical degraded membranes using the essential work of fracture (EWF) method [39]. The EWF method assumes that the crack-tip area consists of two parts: a plastic zone where plastic deformation occurs and a fracture process zone where necking/tearing takes place. Consequently, the en ergy for crack growth can be divided into two parts as follows:
2.2.2. Swelling Membranes with different chemical degradation levels were immersed in 1 M H2SO4 solution for 24 h to remove any possible remaining Fe ions, followed by placing into liquid water at room tem perature (23 � C), and the dimensional change of the membrane in length direction was recorded. The initial dimensions of the membrane were determined as soon as the membranes were taken out from the oven. The extent of swelling ΔL/L0 was calculated as the dimensional change of the membrane in liquid water ΔL, divided by its dry state L0. To ensure the reproducibility of the measurements, each measurement was carried out for three times and the average value was used. To set a baseline, the swelling of an as-received membrane that went through the same ther mal history in liquid water (preheated at 80 � C) was also measured, denoted as PH.
Wf ¼ We þ Wp
(1)
where Wf is the total work of fracture which can be obtained from the area encompassed by the load-displacement curve. We is the essential work of fracture which represents the energy consumed in the inner fracture process zone to create new crack surfaces and is proportional to the area of the fracture process zone. Whereas Wp is the non-essential work of fracture, corresponding to the energy dissipated in the outer plastic zone and is proportional to the volume of the outer plastic zone. Consequently, Equation (1) can be rewritten as: (2)
2.3. Characterization
wf ¼ we þ βwp ⋅L
2.3.1. FTIR Fourier transformer Infrared (FTIR) spectra were collected on an IRAffinity-1S spectrometer (Shimadzu Corp., Japan) within the range of 4000–400 cm 1 with a wavenumber resolution of 2 cm 1. The mem brane samples had a standard size of 5 cm � 1 cm, which covered the whole window to prevent interference. Samples were settled between two KBr plates and placed in the cell for measurement. The spectra were obtained against the air background spectrum.
where wf is the specific work of fracture and we is the specific essential work of fracture. β is a shape factor that reflects the shape of the outer plastic zone. By plotting specific work of fracture as a function of liga ment length L, the specific essential work of fracture, we, and βwp can be obtained from the intercept and slope of the curve, respectively. The loading control mode and strain rate of DENT were the same as uniaxial tensile tests. A more detailed description of the experimental procedure can be found in our previous study [40–42].
2.3.2. In-plane conductivity The in-plane conductivity of degraded membranes were measured using a four-electrode BT-110 conductivity clamp (Scribner Associates). A detailed description can be found in Refs. [36–38]. Prior to tests, samples were cut into rectangular stripes and immersed in 1 M H2SO4 solution for 24 h to remove any possible remaining Fe ions, followed by rinsing in deionized water for several times. The AC impedance was measured in liquid water at room temperature (23 � C) using a Princeton P4000 potentiostat from 1 Hz to 1 MHz. The in-plane conductivity was calculated as κ ¼ LIP/RIPAIP, where LIP is the distance between elec trodes, RIP is the measured resistance, and AIP is the cross-section area of the membrane. Each measurement was repeated at least twice to ensure the reproducibility.
2.3.4. Scanning electron microscopy (SEM) Physical changes at the surfaces and cross-sections of Nafion® 212 membranes before and after chemical degradation were observed using scanning electron microscopy (SU1510, Hitachi, Japan). The crosssections of Nafion® 212 membranes were prepared by breaking mem branes in liquid nitrogen to obtain a smooth cross-section. Samples were then stuck on conductive tapes and gold sputtered under vacuum before SEM observation. Different regions of DENT specimens after fracture tests were captured and compared to illustrate how chemical degrada tion changed the fracture properties of membranes. 3. Results and discussion 3.1. Morphologies of Nafion 212 membrane after chemical degradation
2.3.3. Tensile and double edge notch tension tests Two kinds of mechanical tests were conducted: uniaxial tensile tests and double edge notch tension (DENT) tests. All experiments were carried out on an in-situ tensile testing machine (IPBF-300, CARE Mea surement & Control Co., Ltd., China) at room condition (25 � C/50% RH). The uniaxial tensile tests were performed to determine the stressstrain curves of chemical degraded membranes and basic tensile me chanical properties, such as Young’s modulus, yield stress, strainhardening modulus and break strain. Rectangular specimens were used with the gauge length and width of 30 mm and 10 mm, respec tively. The uniaxial tensile tests were performed under strain control mode with the strain rate of 0.003 s 1. While the stress was recorded with a load sensor, the strain was measured with a non-contact displacement detecting system (NDDS) to avoid displacement by the machine itself.
The morphologies of Nafion 212 membrane after chemical degra dation are shown in Fig. 1. The morphologies of as-received specimens are smooth without any defects. After chemical degradation for 24 h, no discernible changes can be detected. However, increasing chemical degradation duration to 36 h results in bubbles appearing on the surface and pores appearing in the cross section. Similar observations were also made by Tang et al. [43] who ascribed the formation of bubbles to the decomposition of the repeating units of the membrane. Further increasing chemical degradation duration to 48 h increases the density of bubbles, in both surface and cross section areas. In addition, tears of the bubbles and bumps are also observed on the surface. However, with the prolong of chemical degradation time to 72 h, the bubbles become smaller but distribute more densely. Such a phenomenon can be ascribed to the emergence of new bubbles from the surfaces of old large bubbles. 3
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Fig. 1. Surface and cross-section morphologies of Nafion® 212 membrane after chemical degradation for (a) 0 h, (b) 24 h, (c) 36 h, (d) 48 h and (e) 72 h. The crosssection morphologies are shown in the inset at the upper right corner of each image. (f) is the percentage of bubble for membranes with different degradation levels.
It is easily to tell from the cross section in Fig. 1(e) that the membrane is more severely degraded as the surface becomes more roughened and uneven. In addition, with the increase in degradation time, the mem brane’s color changes from transparent (0 h) to opaque (36 h). To characterize the degree of chemical degradation, the percentage of bubbles for membranes with different degradation time is compared in Fig. 1(f). The density of bubbles generally increases with chemical degradation time except for 72 h.
To decide if the extent of ex-situ chemical degradation in this study is comparable to in-situ accelerated stress tests, the total fluoride release is also shown for comparison in the inset in Fig. 2. It can be found that the fluoride release rate in this study are similar with that tested under simulated fuel cell conditions. The fluoride emission rate (FER) of an operating fuel cell was measured intermittently by periodically col lecting effluent water from the anode and cathode gas outlets, which was then analyzed following standard test protocols to calculate the concentration of fluoride ions [23,29,44]. It should be noted that the total fluoride release alone does not always sufficiently enough to assess the health condition of fuel cells, as some degradation regions are highly localized [45–47]. Regardless of its shortcomings, the fluoride release could still be used for rough estimation due to its convenience, as the fluoride release can be easily measured from the products of fuel cells.
3.2. Fluoride release The extent of chemical degradation of membranes relies on many factors, such as hydrogen peroxide concentration, pretreated tempera ture and degradation time. To compare the mechanical properties of membranes with different degradation levels, a parameter is needed to represent the extent of degradation. As such, the fluoride release is used as an indicator. The total fluoride release as a function of degradation time is shown in Fig. 2. The total fluoride release increases almost lin early with degradation time, indicating a constant fluoride release rate.
3.3. Molecular structural changes of the membrane To characterize molecular structural changes induced by chemical degradation, FTIR investigation of the membrane was carried out. The obtained FTIR spectra is normalized to the CF2 stretching peak at 1145 cm 1, as shown in Fig. 3. According to the literature [48–50], different absorption peaks correspond to different assignments. The C–F
Fig. 2. Total amount of fluoride release as a function of degradation time. To compare the ex-situ measurement in this study with in-situ accelerated stress tests, data from Patil et al. [31] and Lim et al. [30] are also shown for comparison.
Fig. 3. FTIR spectra of Nafion® 212 membrane with different chemical degradation levels. 4
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absorption peaks at 1145 and 1200 cm 1 as well as the C–C absorption peak at 1300 cm 1 are characteristics of the PTFE backbone of the membrane. While the absorption peak at 970 cm 1 is assigned to the symmetric C–O–C stretching, the peak at 1057 cm 1 is assigned to the symmetric S–O stretching. These two peaks correspond to chemical bonds at the pendant side chain and terminated sulfonic groups of Nafion membrane, respectively. With the increase of chemical degra dation time, little variation in normalized C–F peak intensity is observed. By contrast, the heights of the C–O–C and S–O normalized peaks decrease significantly, suggesting that the sulfonic acid groups and the ether linkage at the pendant side chain are more vulnerable to radial attacks. Even though the molecular structure changes of Nafion membrane after chemical degradation have been investigated by many researchers [43,49,51–55], it is not trivial to confirm the molecular structure in this paper as different mechanisms are reported in the literature, and me chanical responses of the membrane is strongly related to its molecular structure. While some researchers concluded that chemical degradation took place more preferentially from the side chains [49,51,55], some researchers found that membrane decomposition started from the ends of main chain [43,54], both main chain and side chain [53], as well as not changed by degradation [52]. One possible reason might be due to the reduced –COOH end groups in new generation chemically stabilized PFSA membranes where the chain scission mechanism is more prevalent [56]. PFSA membranes have hydrophobic PTFE main chains and hy drophilic side chains terminated with sulfonic acid ionic groups, creating a nanophase-separated morphology. Because of the molecular structure changes, the morphology of the membrane also changes accordingly. According to the small-angle X-ray scattering (SAXS) re sults by Kusoglu et al. [26], the domain spacing of degraded membrane increased with increasing decomposition level. A relatively large non-ionic region which resembled void in the membrane was observed by Venkatesan et al. [57] after chemical degradation through trans mission electron microscopy (TEM) observation. Depending on the temperature and relative humidity, mechanical loads can be transmitted through hydrophobic or hydrophilic domains [58]. Therefore, it is anticipated that the mechanical properties will change, which will be investigated later.
uptake which is intimately related to the ion exchange capacity (IEC) of the membrane. As discussed in Fig. 3, the sulfonic group at the end of side chain is attacked by radicals, reducing the water uptake capacity of the membrane. Decreased water uptake after chemical degradation was reported [49,51], which was ascribed to morphological changes. The reduced swelling is beneficial from the mechanical perspective in that the in-plane stress can be reduced, improving the mechanical durability. Patil et al. [31] compared the contractile stress before and after chemical degradation, and it was found that the peak stress of degraded mem branes was lower than that of fresh samples. However, accompanying chemical degradation is the entanglement changes that might reduce the resistance for the membrane to withstand the resulting stresses. In their study, even though the stress of degraded membrane was lower, the degraded membranes broke into two pieces, while the fresh membrane maintained structural integrity. Thus, the mechanical properties, espe cially the fracture behavior of the membrane, need to be evaluated in order to have a full evaluation of membrane durability under combined chemical and mechanical degradations. To have a better understanding of the degradation level, especially how the performance of the membrane is impact by chemical degra dation, the in-plane conductivity of different degraded membranes are compared in Fig. 4(b). The conductivity decreases with increasing degradation time, similar with the trend in swelling. After chemical degradation for 24 h, the conductivity slightly decreases, indicating that short-term chemical degradation have minor impact on the conductiv ity. The similar conductivity value between 24 h-degraded membrane and 0 h-PT membrane suggests that the Fe ion is exchanged completely. However, further increasing degradation time to 36 h causes 40% decrease in conductivity, reducing from 0.083 S/cm for PH membrane to 0.0497 S/cm for 36 h degraded membrane. The significant reduction in conductivity after degradation for 36 h appears concurrently with the changes in surface morphology in Fig. 1, indicating that a certain amount of time is needed to induce decay in properties of the membrane. Further increasing degradation time from 36 h to 48 h and 72 h de creases the conductivity from 0.0497 S/cm to 0.0406 S/cm for 48 h degraded membrane and to 0.0378 S/cm for 72 h degraded membrane. Even though the conductivity decreases for more than 50% after degradation for 72 h, the membrane can still function well. The reduced conductivity is ascribed to the possible loss of sulfonic groups at the end of side chains, as indicated from the FTIR study in Fig. 3.
3.4. Swelling and conductivity The dimensional stability of the membrane is an important factor affecting the mechanical durability. Excessive swelling results in significantly in-situ residual stress, impairing the mechanical integrity of the membrane. Therefore, the extent of swelling is firstly evaluated before mechanical tests as shown in Fig. 4(a). With the increase of degradation level, the swelling decreases from 14.4% for 0 h-PH mem brane to 14.1% for 24 h degraded membrane, finally to 11.6% for 72 h degraded membrane. Decreased hygral expansion after in-situ acceler ated stress test degradation was also observed by Alavijeh et al. [28]. Such a reduction in swelling can be ascribed to the decreased water
3.5. Tensile properties The stress-strain curves of membranes degraded for different time are compared in Fig. 5 (a). The AsR membrane is ductile with high tensile strength. After chemical degradation for 24 h, the tensile strength does not change much while the proportional limit stress decreases significantly. For PFSA membranes, the proportional limit stress is traditionally used to represent the yield strength of the membrane [42, 59,60]. Increasing degradation time to 36 h results in significant reduced ductility. The break strain decreases from 265% for 24 h
Fig. 4. Comparison of (a) swelling and (b) conductivity of membranes with different degradation levels. 5
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Fig. 5. (a) Stress-strain curves of membranes degraded for different time; (b)–(e) are comparison of Young’s modulus, proportional limit stress, strain-hardening modulus and break strain, respectively.
degraded membrane to 82% for 36 h degraded membrane. The changes of break strain with degradation time is in accordance with the morphology changes in Fig. 1 where bubbles start to emerge on 36 h-degraded membrane. The obtained mechanical results could be due to morphological changes, as the presence of bubbles is a macroscopic reflection of microstructural and morphological changes. The reduction in ductility at an early stage was also observed during accelerated stress test where the ductility diminished after only two AST cycles [42]. Further increasing degradation time to 48 h and 72 h leads to gradually decrease of ductility. It seems that there is a critical point where the ductility starts to drop significantly. As discussed in section 3.1, non-ionic regions and voids will form after chemical degradation. These defects cause locally stress concentration and severe deformation to final failure. To better illustrate how mechanical properties changes with degra dation time, Young’s modulus, proportional limit stress, strainhardening modulus and break strain are extracted from stress-strain curves and compared in Fig. 5 (b)–(e). With increasing degradation level, the Young’s modulus, proportional limit stress and break strain decrease, while the strain-hardening modulus increases. In particular, significant changes are observed for Young’s modulus and break strain after chemical degradation for 36 h. Deformation at different stages have different morphology origins. While small strain deformation corresponds to chain rotation of bundles, large strain deformation cor responds to alignment of aggregates by disentangling from each other [61]. Based on these morphology assignments and molecular origins, the deformation mechanisms can be tentatively obtained. The degraded ionic groups at the terminal of side chains reduced the electrostatic interaction between side chains, thus lowering the restriction imposed on segmental motions and making it easier for chain rotation to occur. While changes in Young’s modulus after chemical degradation have been reported by some researchers [28,31], the strain-hardening prop erties have never been investigated. Patil et al. [31] thought that the molecular weight would decrease after chemical degradation and chain slippage would occur more easily as the shortened chains induced by chemical degradation would be less likely to form entanglements. However, the results in this study are not coincide with the above findings, which suggest that a more sophisticated mechanism might exist and more factors should be taken into consideration. One possible reason might be the morphological rearrangement. For instance, the domain spacing increases with increasing decomposition level [26]. The reduced ductility of degraded membranes increases the risk of me chanical failure as the membrane is unable to accommodate large deformation. Consequently, cracks and pinholes are more likely to form during humidity cyclings, leading to crossover of reactant gases. This
also explain why the life time under combine chemical/mechanical accelerated stress tests are much shorter than in isolated chemical or mechanical accelerated stress tests [9,21–23]. The reduced ductility alone is not sufficient to understand the mechanical failure mechanisms of the membrane. To address this issue, the fracture process and mechanisms also need to be considered. 3.6. Fracture behavior and mechanisms The above tensile properties describe the ability of the membrane to resist deformation and crack formation. In the presence of cracks, the fracture behavior is needed to assess the ability of the membrane to resist crack growth. To this end, the load-displacement curves of mem branes with different degradation levels are compared, as shown in Fig. 6. For AsR membranes in Fig. 6 (a), the load increases to the maximum point and gradually drops. The maximum load point corre sponds to crack initiation, after which crack starts to propagate [41,62]. The gradual decrease of load suggests that the crack propagates in a moderate way, showing a ductile fracture feature. However, with the increase of degradation level, such a feature diminishes. The load starts to drop abruptly for 36 h degraded membrane, showing a brittle fracture feature. Such a feature is more pronounced for 72 h degraded membrane where the load almost drops to zero as soon as reaching to the maximum load. Therefore, the membrane experiences a transition from ductile fracture to brittle fracture with increasing degradation level. To quantitatively assess how the crack initiation energy and propa gation energy change in response to chemical degradation, the essential work of fracture method is used as shown in Fig. 7. The intercept and slope are extracted and compared in Fig. 7(b)–(c), respectively. With the increase of degradation level, both the slope and intercept decrease. The intercept represents the energy needed for crack initiation, while the slope represents the ability of the membrane to resist crack propagation. The decreased intercept and slope with increasing degradation level suggests that the crack is not only easy to initiate but also more easily to propagate. The fracture initiation energy decreases by more than 40%, from 8.7 kJ m 2 for AsR membrane to 5 kJ m 2 for 48 h degraded membrane. A notable exception is the 72 h degraded membrane with the fracture initiation energy as high as 9 kJ m 2. The origin of such a discrepancy is not clear at the present stage, which could be related to microstructure and morphological changes. Nevertheless, this phe nomenon along with the increased strain-hardening modulus in Fig. 5 imply that the mechanical behavior are not as simple as we had spec ulated. With respect to the slope, significant decrease in slope (more than 70%) is observed, from 4.2 MJ m 3 for AsR membrane to 1 MJ m 3 for 72 h degraded membrane. As the slope is related to the outer plastic 6
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Fig. 6. Load-displacement curves at different ligament lengths for (a) AsR membrane; (b) 24 h degraded membrane; (c) 36 h degraded membrane; (d) 48 h degraded membrane; (e) 72 h degraded membrane.
Fig. 7. (a) Specific work of fracture as a function of ligament lengths for different degraded membranes; (b)–(c) are comparison of specific essential work of fracture, we and βwp, respectively.
dissipation zone where crazing is more likely to occur [39], the reduced non-ionic regions observed by Venkatesan et al. [57] are expected to promote craze formation. The findings in this study suggest that the detrimental effect of chemical degradation not only manifests as the reduced crack initiation resistance, but more importantly the crack propagation resistance is largely undermined. A fuel cell will not fail immediately after a crack forms. Instead, it will continue operating for a quiet long period of time, as water generated during operation will cover the surface of cracks, preventing crossover of reactants [63]. Thus, crack propagation resis tance is an important factor affecting the lifetime of fuel cells [16,17, 42]. Therefore, when designing new membranes, the crack propagation resistance as well as its sensitivity to chemical degradation should be
evaluated to achieve a high lifetime. To clarify the mechanisms of crack propagation process, the crosssections of cracked degraded specimens are examined, as shown in Fig. 8 (a). The AsR specimen is taken as an example for illustration. Several areas with different thicknesses are observed from which the crack propagation process can be identified. The region with the thickness of 50 μm is assigned to the pre-crack region which is cut by razor blades and the thickness is not affected. Right next to the pre-crack region is the stretch zone which represents the plastic zone size ahead of crack tip. The thickness of this region is 22 μm, which is less than 50% of the original thickness. Such a reduction in thickness is induced by severe plastic deformation. A similar reduction in thickness of the membrane is also observed during the fatigue crack propagation [19]. The shape of 7
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Fig. 8. (a) An illustration of the cross section of AsR Nafion membrane with notch zone, stretch zone and fracture zone. The crack propagates from left to right. (b) Comparison of stretch zone width of membranes with different degradation levels.
the stretch zone is semi-elliptical, and the region near the crack surface is severely deformed while that far from the crack surface is less affected. In addition, severely deformed strips can be also observed on the cross-section area. These features manifest that severe plastic deforma tion occurs in the stretch zone. At the right side of this image is the fracture region with the thickness reduced to 28 μm. This region is assigned to fast crack propagation where instability crack growth oc curs. Thus, the stretch zone represents the plastic zone ahead of crack tip. To quantify how the plastic zone ahead of crack tip evolves with degradation level, the stretch zone width which is the width between the notch region and fracture region, is compared for different degraded membranes in Fig. 8 (b). Similar with the trend of fracture resistance, the stretch zone width decreases with the degradation level. In other words, the size of the plastic zone ahead of crack tip decreases with degradation level. As a result, the energy dissipated through the plastic zone is reduced, accounting for the decreased slope in Fig. 7. In this way, chemical degradation affects crack propagation through reducing the plastic zone size of crack tip and exhibiting a ductile-to-brittle transition feature.
propagation resistance is evaluated according to the fracture behavior of the membrane based on fracture mechanics. After degradation, the loaddisplacement curves change from ductile fracture to brittle fracture behavior. Using the essential work of fracture method, both the crack initiation and propagation energies are found to decrease with degra dation level, with an exception of the 72 h degraded membrane. Based on the microstructure assignment of crack propagation, the crosssections of degraded membranes after double edge notch tension tests are examined. The stretch zone width, which represents the plastic zone size ahead of crack tip, is found to decrease with degradation level. Thus, the decreased crack propagation resistance after chemical degradation is ascribed to the reduced plastic zone size ahead of crack tip, which is induced as a result of membrane embrittlement. Therefore, chemical degradation not only impairs tensile mechanical properties, but also undermines crack propagation resistance through minimizing crack tip plastic zone size and thus reducing energy dissipation. The results in this study are helpful for the understanding of mechanical degradation and fracture failure induced by chemical degradation, which will provide more insights into the durability of membranes and develop effective mitigation strategies to enhance membrane durability.
4. Conclusions
Declaration of competing interest
In this study, the mechanical properties after chemical degradation, especially the fracture behavior, are investigated. To characterize the extent of degradation and link mechanical properties with degradation level, the amount of fluoride release is measured as a function of degradation time. It is found that the total amount of fluoride release increases almost linearly with degradation time. More importantly, the fluoride release value is similar with that obtained from in-situ acceler ated stress tests, indicating that the mechanical properties obtained here could, to some extent, represent the actual material state of membrane in operating fuel cells. The FTIR is also carried out to probe the mo lecular origin of chemical degradation in order to account for changes in mechanical properties induced by chemical degradation from a molec ular perspective. The changes in FTIR spectra suggest a predominant loss of sulfonic acid groups and the ether linkage from side chains. As the mechanical stress is generated due to sorption and desorption of water, the swelling properties of degraded membranes are also measured to characterize the potential induced stresses during humidity cyclings. It is found that the swelling behavior decreases with degradation time, which is ascribed to the reduced sulfonic acid groups and therefore reduced water uptake. Accompanying with reduced swelling behavior is the decreased conductivity which shows a 50% reduction after degra dation for 72 h. With respect to mechanical properties, increasing degradation level reduces the Young’s modulus, proportional limit stress and break strain, while the strain-hardening modulus is increased, implying the embrit tlement of the membrane after chemical degradation. These properties only accounts for factors leading to crack initiation. The crack
The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement Xiaoyi Sun: Formal analysis, Investigation, Methodology, Visuali zation, Writing - original draft. Shouwen Shi: Conceptualization, Funding acquisition, Methodology, Project administration, Supervision, Visualization, Writing - review & editing. Yuanjie Fu: Investigation, Methodology. Jian Chen: Methodology, Project administration. Qiang Lin: Investigation, Methodology, Project administration. Jiaqi Hu: Investigation, Methodology. Cong Li: Methodology, Project adminis tration. Jiayao Li: Investigation, Methodology. Xu Chen: Conceptuali zation, Supervision, Writing - review & editing. Acknowledgements The authors gratefully acknowledge the financial support from the National Natural Science Foundation of China (No.51805364, No.11902216), China Postdoctoral Science Foundation (No. 2019TQ0225, No.2019M661017). The research is also partial finan cially supported by the Open Research Fund of Science and Technology Innovation Platform of Key Laboratory of Efficient & Clean Energy Utilization, Changsha University of Science & Technology (No. 2018NGQ002). 8
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