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10FUEL CELL
Статья поступила в редакцию 08.08.12. Ред. рег. № 1398
The article has entered in publishing office 08.08.12. Ed. reg. No. 1399
УДК: 546.11; 621.352.6; 547.022
СИНТЕЗ ПРОТОН-ПРОВОДЯЩЕЙ МЕМБРАНЫ КИСЛОГО СУЛЬФАТА ЦЕЗИЯ ДЛЯ ИСПОЛЬЗОВАНИЯ В ВОДОРОДНОМ ТОПЛИВНОМ ЭЛЕМЕНТЕ
'Университет Кейптауна,
Факультет проектирования экологических и технологических систем, химического машиностроения п/я Х3, Рондебош, 7701, Кейптаун, ЮАР Тел.: +27216507727; e-mail: [email protected] 2Университет Кейптауна, Факультет химического машиностроения, Центр компетенций HySA/катализа п/я Х3, Рондебош, 7701, Кейптаун, ЮАР Тел.: +27216504988; e-mail: [email protected] 3Институт физики твердого тела при Латвийском университете Латвия, Рига, LV-1063, ул. Кенгарага, д. 8 Тел.: +371 26124062
Заключение совета рецензентов: 20.08.12 Заключение совета экспертов: 25.08.12 Принято к публикации: 30.08.12
Для стабилизации неорганического твердого электролита (кислого сульфата цезия) изготовлялся композит. Это необходимо, т.к. при контакте с топливом меняется растворимость и начинается деградация мембраны, что неблагоприятно влияет на протонную проводимость. Для этого использовались ПТФЭ, обладающие высокой химической стабильностью, и двуокись кремния для улучшения проводимости и увеличения площади поверхности. В статье рассмотрены различные композитные мембраны с различными объемными отношениями составляющих материалов и отличающиеся выбором прекурсоров.
Ключевые слова: топливный элемент, водород, протон-проводящая мембрана, сульфат цезия.
SYNTHESIS OF CESIUM HYDROGEN SULPHATE PROTON CONDUCTING MEMBRANE FOR HYDROGEN FUEL CELL APPLICATIONS
lUniversity of Cape Town, Environmental and Process Systems Engineering, Chemical Engineering Private Bag X3, Rondebosch, 7701, Cape Town, South Africa +27216507727; [email protected] 2HySA/Catalysis Competence Center, Department of Chemical Engineering, University of Cape Town, Private Bag X3, Rondebosch, 7701, Cape Town, South Africa +27216504988; [email protected] 3Institute of Solid State Physics, Department of Physics, University of Latvia 8 Kengaraga str., Riga, LV-1063, Latvia Tel.: +371 26124062
Composite formation was used to stabilize the inorganic solid acid electrolyte, cesium hydrogen sulphate membrane due to the solubility and degradation the membrane undergoes once the fuel is brought in contact with the membrane, which adversely affects the conductivity. The additive material of choice to prepare these composites will be PTFE, due to its resistance to chemical attack, and silica to enhance conductivity and surface area. Different types of composite membranes will be developed with varying volume fractions of the component concentration and using different precursors.
Keywords: fuel cell, hydrogen, proton-conducting membrane, cesium sulfate.
S. Naidoo1, Q. Naidoo2, G. Vaivars'
3
Referred: 20.08.12 Expertise: 25.08.12 Accepted: 30.08.12
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Sivapregasen Naidoo
Organization: University of Cape Town, Post-doctoral Research Fellow, PhD.
Education: PhD, University of the Western Cape, Faculty of Natural Science (2005-2008).
Experience: The author has more than 10 years' experience in large manufacturing concerns in the pharmaceutical industry including Fine Chemicals, Vital Health Foods and GlaxoSmithKline, where he worked in all science related departments from laboratory scientist, validation chemist, officer and related senior posts. He also was involved in global seniority decision making with regards to all areas of the pharmaceutical industry including regulatory and site auditing.
Main range of Scientific Interests: Dr Naidoo has vast experience in the application of Nano-science in synthesis methods. He has also performed research in the fields of Environmental Science, Catalysis, Fuel Cells, Bio-fuel Cells, Biogas Methane Decomposition, Pathology and Epidemiological Study.
Publications: Dr Naidoo has published regularly since (2008) in international journals.
Organization: University of Cape Town, Post-doctoral Research Fellow, PhD.
Education: PhD, University of the Western Cape, Faculty of Natural Science (2008-2011).
Experience: 2011-present Post-doctoral research fellow in Chemical Engineering, University of Cape Town. 1995-2003 Senior analytical chemist in Liu Zhou Steel Iron Company, China.
Main range of Scientific Interests: Design and synthesis functional catalysts, fuel cell research and electrochemistry study.
Publications: Dr Naidoo has published regularly since 2008 in international journals.
Qiling Naidoo
Organization: University of Latvia, Department of Chemistry, Institute of Solid State Physics, leading scientist and docent.
Education: PhD, Latvian Academy of Science, 1991.
Experience: From 2008: Leading Scientist, Hydrogen energy materials laboratory, Institute of Solid State Physics, Univesity of Latvia. From 2009: Docent in Physical and Electrochemistry, Department of Chemistry, University of Latvia. From 2011: the Latvia representative in The Fuel Cells and Hydrogen Joint Undertaking.
Main range of Scientific Interests: Fuel cell research and Physical chemistry, proton conducting polymer membranes, ionic liquids.
Publications: Dr. Vaivars has published more than 50 peer-reviewed articles, 80 conference abstracts and proceedings and 9 patents.
Guntars Vaivars
Introduction
Researchers throughout the world are seeking effective proton conducting membranes for fuel cell application. Proton conducting membranes form an integral part of the PEMFC. A major investment of resources has been made to obtain a proton conducting membrane, which could be incorporated in mobile power sources. Currently the major developments in fuel cell technology using hydrogen or methanol are based on perfluorinated and sulfonated solid polymer electrolytes. The fact that fuel cells run on hydrogen or hydrogen carriers points to the link between fuel cells and the so-called hydrogen economy. This is a significant observation since hydrogen fuel technology is a sustainable technology and may be one of the best long-term solutions to the energy problem. These solid electrolytes offer advantages over classical liquid electrolytes such as sulphuric acid. These advantages include reduced methanol fuel crossover and higher power densities. Sulphonic acid insertion into solid
polymer membrane is currently employed and is commercially referred to as Nafion. This membrane is very popular when used as a low temperature proton conducting polymer membrane with thickness between 25 and 175 ^m. This membrane is a good proton conductor and is the preferred membrane for both hydrogen and direct methanol fuel cells (DMFC). However, it has some problems that limit the manner in which it can be successfully used in fuel cells, such us the methanol fuel crossover reduces the cell potential difference.
CsHSO4 experiences two, phase transitions at temperatures of 318 and 417K. At 417K a high proton conductivity state is reached at a value of 10-2^-1cm-1. This super-ionic property makes CsHSO4 a favorable component in electrochemical devices such as batteries, fuel cells, displays and super-capacitors. CsHSO4 can also be used as a model to understand the mechanism of proton conductivity in "dry" proton conductors where "wet" and "ion jump" type conduction takes place [1-7].
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Proton movement contributing to ion conductivity proceeds by two interactions as follows: Firstly, an attractive interaction, where the diffusion of cations from their regular sites to the surface of the insulator (dispersoid) stabilized by sorption, thus, increasing the vacancy concentration at the interface. And secondly, repulsive interactions, which occur when the cation accumulates in the interstitial site area driven by the insulator. The two interactions contribute to defect formation at the interface, which can improve the interfacial ion conductivity of the ionic conductor/insulator composites [8]. Order-disorder phase transition responsible for ionic conduction is stimulated by the defects in the interface region. The critical temperature of the order-disorder phase transition is thought to be lowered by the defects occurring at the interface. Once the temperature drops below the super-ionic phase transition, there is an increase in ionic conductivity [9]. The conductivity of CsHSO4-SiO2, among other ionic salt composites, is governed by the ionic salt and oxide structural make up at the interface region. The stability of the interface region is dependent on the adhesion energy between ionic salt and oxide. Also the conductivity is dependent on the physical properties of the interface region between the ionic salts and the oxides, which essentially differ from the properties of individual compounds. It was also found that the structural reconstruction or formation of the meta-stable interface occurs due to strong adhesion between these particles. The resulting interface from the ionic salt spreading over the oxide particle will be structurally disordered accounting for the high conductivity. CsHSO4 has 3 phases. Phase I is the room temperature structure consisting of a hexagonal array of cesium ions stacked along a bent direction. Zigzag chains of hydrogen bonded HSO4- ions are perpendicular to the loose packed Cs+ ion plane and imprison the Cs+ ions. The structure is an intermediate between the non-protonic glasrite structure and the ferroelectric CsH2PO4 structure with two different types of hydrogen-bonded chains surrounding the Cs+ cations. X-ray diffraction shows the phase I structure, which is confirmed by infrared and Raman spectra, particularly the statistical distribution of protons between two potential minima of the OH...O hydrogen bonded chains. There are two first order phase transitions one observed at approximately 318 K (I—II) and the other at 417 K (II—III). Due to the weakening of the hydrogen bonds and a structural disorder of the anions, the I—^II transition has been interpreted on the basis of spectral vibration analysis where a conversion of infinite chains into cyclic dimers occurs. For the superionic phase III there is a further weakening of the hydrogen bonds and an increasing level of disorder. There is also evidence of rapid reorientation of HSO4- species and translation disorder of Cs+ ions. The high proton mobility was observed by H and D NMR, as well as quasi-elastic neutron scattering [6-14].
Upon cooling, metastable phases (III—II—I) occur, with long annealing below 280 K being required to reach the phase I stage. More importantly, the kinetics is largely influenced by the presence of water traces. Samples prepared by slowly evaporating solutions of Cs2SO4 and H2SO4 to produce small crystals (0.5 to 1 mm), do not undergo phase transitions when stored in a sealed container. However complex DSC traces are observed for the temperature range 315-380 K. This is supported by the measured enthalpy values (AH) for the I—II transition which varies reversibly between 4.8 and 14.4 kJ/mol being dependent on the "defect" concentration, while the II—III (II^III) phase transition AH value yields approximately 14.4 or 29.2 kJ/mol. The stretching intensity of the S-O and S-OH bands is directly influenced by the grinding force during the sample preparation and is not affected by the time factor. For the CsI pellet the pressure must be applied for a few hours before the 998 cm-1 band disappears but does not vary with time thereafter. Weak and strongly ground samples show different spectra. The Raman spectrum shows typical cyclic dimer occurrence as in phases II and I. Thus, CsHSO4 at room temperature can be structurally modified by mechanical treatment which converts the infinite chains of HSO4- ions into cyclic dimers. The applied pressure also referred to as the mechanical treatment, is responsible for the chain/dimer ratio.
Colourless crystals are formed depicting a twin structure, after equivalent molar amounts of Cs2SO4 and H2SO4 in an aqueous solution are heated to 333 K and allowed to cool to room temperature. Using a three dimensional Patterson map the positions of the Cs and S were found [13]. A map of residual electron density showed an anomaly reflected near the Cs atom of maximum of 1.6 eA-3. Cs and SO4 lie on the mirror planes y = V and %. There are one-dimensional chains of hydrogen bonds along the b axis. In the Para-electric phase CsHSO4 resembles CsH2PO4 the only difference is the extra H atom which combines O (1) and O (211) [14, 15]. There is also evidence of a disordered H atom. Below room temperature a phase transition could be responsible for the conversion to the ordered phase.
When using vibrational spectra, the external modes obey the selection rules where Z = 2 for the C2h factor group. Thus the average symmetry of the crystal is seen in the same way as by X-rays [16]. For the Raman spectra seven out of the nine Raman active lattice bands below 300 cm-1, were observed. In the low temperature powder spectrum seven Raman bands were also observed.
The lattice bands can be assigned approximately in terms of the rotational (R) and translational (T') vibrations of the HSO4- ion (250-100cm-1) and translational modes (T') of the Cs+ cation (below 80 cm-1). Internal vibrations do not indicate the center of symmetry or the SO4 mirror plane. In the skeletal bending region (650-350 cm-1) there are more bands than expected with coinciding frequencies [17, 18]. The
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spectral pattern for S-O stretching is similar to that of other alkali hydrogen sulphates which do not contain symmetrical hydrogen bonds [19]; the particularly strong Raman bands near 990 and 860 cm-1 could be due to an S-O proton acceptor and a S-OH proton donor group, respectively. For the OH—O hydrogen bond the vOH mode is responsible for broad ABC bands. These bands are characteristic of strong hydrogen bonds and their center of gravity at 2400 cm-1 is in agreement with O—O distances of 2.57A [18]. When CsHSO4 is heated to 350K, I—^II transition takes place. The vS-O frequency at 998 cm-1 shift to 1024 cm-1 and the vS-OH band broadens and decreases to 850 cm-1. As mentioned before, this also reflects the conversion of infinite chains to cyclic dimers containing a weaker hydrogen bond. The vOH band shift towards higher wave numbers also indicates the weakening of the hydrogen bonds. The Raman bands for R'HSO-4 and T'HSO-4 broaden considerably, in the lattice region, showing the anions to be in a structural disorder [17, 18].
Further broadening of the super-ionic phase III bands due to internal vibrations indicates an increasing structural disorder. More significantly, is the further loosening of the hydrogen bonds supported by the upward movement or shift of the vOH band. This is also indicated by the increased difference between vS-OH and vS-O frequencies. From an analytical point of view, the most obvious and striking change is observed in the lattice region where all the external mode bands collapse into a broad band near the Rayleigh line as for H3OClO4. This is also evident in a highly disordered (plastic) crystal, indicating a rapid re-orientation of the HSO4-ions on their sites and translational disorder of cesium with this being an identical characteristic in other cationic superconductors [20-22]. Thus the proton transfer in phase III is only partially involved in the total ionic motion but more importantly the Cs+ ions are major contributors to the conduction process [23]. This is in agreement with the super-ionic properties of parent nonproton conducting sulphates and work done using D NMR on CsDSO4.
This view is supported by deuterium studies, which shows the conductivity varying only in phase III where the activation energy (Ea) is 0.32 and 0.27 eV for CsHSO4 and CsDSO4 respectively. Due to a resulting 200 fold decrease in conductivity when hydrogen is substituted by deuterium, phases, II and I low temperature conductivity should essentially be protonic. The conductivity mechanism strongly supports a proton tunneling stage. A possible explanation for the Cs+ ion contribution to the conductivity is the transition from the super-ionic to the low temperature phase being accompanied by a 500 times drop in conductivity associated with a difference in the electrostatic repulsion between cations. The sulphate orientation is associated with the restricted motion of the larger cations preventing local motion of the protons. This property is evidence of the large pressure effect on super-ionic properties [24-27].
A compounds conductivity mechanism which depends on H3O+ ions, relies on the H2O content. This was the basis for the addition of a heterogeneous component with the intention to increase the proton conductivity of CsHSO4. Initial heating shows an increase in conductivity but this is due to the loss of water on drying. However the completion of three heating and cooling cycles heating to 485 K and cooling to 300 K, produced reproducible conductivity results. The composite investigated was (1-x) CsHSO4-xSiO where x, the mole ratio of heterogeneous component SiO2, varied from 0 to 0.8. The addition of the SiO2 does not increase the conductivity of the composite membrane if the initial powder was not heated near to its melting point temperature. Shown here, the conductivity is largely influenced by the mole ratio of SiO2. Increasing the SiO2 mole fraction in the proton conducting membrane composite increases the ion conductivity in the super-ionic phase at low temperatures. Increasing the SiO2 concentration of the composite to greater than 20% shows a significant increase in the conductivity. This concentration dependency supports the percolation theory that conducting pathways formed by highly conducting interfaces. This high conductivity at low temperatures can be ascribed to water adsorbed on the surfaces of the highly dispersed SiO2. To eliminate this contribution to the conductivity tests were carried out where the partial pressure of water was decreased to 10-1 Pa. There was no effect on the conductivity at low and high temperatures, thus supporting the conduction mechanism as structural protons rather than hydronium ions although it is a hydrated compound [28].
In the high temperature region (380-483 K) composites containing SiO2 with concentrations in the region of 10-60%, have higher conductivities than pure CsHSO4. This could be the stabilization of the molten CsHSO4 on the surface of the highly dispersed SiO2. Based on the conductor-insulator type percolation where the insulator does not possess protons or any conductive properties, high SiO2 concentrations in the region of 6080% result in a decrease in conductivity [29, 30].
Heterogeneous doping component addition of the composite with SiO2 affects not only the conductivity properties but also the temperature at which the super-ionic phase takes place. Differential thermal analysis (DTA) suggests two intensive peaks. The first DTA peak appears at 485 K for pure CsHSO4, and the other at 414 K representing the process of crystallization and phase transition respectively. Reported data supports these results where the phase transition enthalpy is equal to 5.5 kJ/mol and the crystallization enthalpy is equal to 9.5 kJ/mol [28, 31]. As the amount of SiO2 increases the temperature and enthalpy of crystallization (melting) and phase transition, decrease. The thermal effects relating to the phase transition of crystallization (melting) decrease one order of magnitude as compared to membrane at the SiO2 concentration greater than 50%, the XRD patterns show a weak halo appearing in addition to the crystalline phase. When the SiO2 concentration is greater than 30%
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in the composite, an additional reflection at 20 = 24.8 appears. The additional reflection corresponds to the super-ionic a-phase of CsHSO4 and lacks the X-ray diffraction patterns of the p-phase [32, 33]. This depicts the favourable influence heterogeneous doping has on the thermodynamic parameters of the CsHSO4 as an ionic conductor. XRD analysis of composite CsHSO-4SiO2 at varying concentrations of SiO2, after the composite was heated to 483 K, showed a decrease in intensity with the peaks becoming broader notably in composites containing more than 50% SiO2. All reflections correspond to the structure of low temperature p-phase of CsHSO4 in composites with SiO2 content greater than 50% in the CsHSO4. Possible explanations for the sharp increase in the composite conductivity include the diminishing of grain size and the formation of a spacecharge region at the interface. For low concentrations of SiO2; firstly, the presence of the amorphous CsHSO4 phase on the highly dispersed SiO2 surface relating to SiO2 concentrations less than 50% and secondly the above in combination with the possible stabilization of CsHSO4 with PTFE could be responsible for the sharp increase in the composite conductivity [28].
Due its low surface energy, PTFE is not wettable and does not directly adhere to any solid. PTFE can thus be characterized as both hydrophobic and oleophobic. The conductivity of PTFE is of the order of 10-10 Scm-1. The conductivity of the composite is clearly lower than that of the pure ion conducting substance as the composite is dominated by grain boundary contribution i.e. ion movement across boundaries of neighboring conducting substance particles. For PTFE in LiSn2P3Oi2 composite a percolation expression c = o0(v-vc)', for v > vc; where c0 is a pre-exponential factor, vc is the critical volume fraction of the conducting substance LiSn2P3O12, t is an exponent. For a 3-dimensional percolation system the t value is approximately 2, slightly lower than the exponent. The critical volume fraction is nearly the same as the conducting substance area fraction of approximately 0.3. For a random distribution of conducting particles (vc = 0.15-0.17) higher values are predicted for the critical volume fraction. This is explained in the aggregation process to give extended clusters formed in isolation the form of an intermediate stage. To lower the critical volume fraction the composite could be prepared by a different procedure that disperses of the ceramic particles.
The critical volume fraction varies between 0.1 and 0.6 in composite media [34]. Ionic resistance decreases as a function of the volume fraction of the higher conducting phase when the average number of contacts per grain is about 1 and stops decreasing when the average number of contacts per grain is about 2 [35].
Experimental Methods
SEM images were obtained on a Hitachi x650. The cross-sections of the composite membrane were obtained by breaking the membrane into small pieces under liquid
nitrogen. Hitachi x650, Accelerating voltage 25 kV, Aperture 0.4 mm, Tilt angle 0°, Resolution 6 nm and Working distance 15 mm.
Membrane fabrication was as follows, 25.0 g CsCO3 (99,9%, Sigma) was dissolved in 200 ml water in a reaction vessel. 15 ml (Kimix) 98% H2SO4 in 200 ml water was added to this solution, drop wise, by means of a burette. Total volume was 400 ml. The solution was stirred for 24 hours overnight. The solution was then heated in an oven at 378 K until it formed a thin flaky top layer and hard solid amorphous bottom layer. The solid mass adhered to the reaction vessel. The composite membrane was formed as follows, addition of PTFE and Silica to 7.0014 g of wet crystal precipitate was added 2.0091 g PTFE (Sigma) fine powder and 0.5012 g fumed silica 380 (Sigma), and all ground together. With the addition of water, a final volume of 400 ml was reached. It was heated to 573 K while continuously stirring, for half an hour, and then maintained at 383 K thereafter without stirring until all the solvent had evaporated.
Results and Discussion
The composition, Fig. 1, of the following membrane is 80% CsHSO4, 3% SiO2 and 17% PTFE pressed together to form a solid acid membrane composite. Membrane was 2 mm thick with 1 mm between the electrodes. It had a 46 cm diameter and weighed 7.2 g.
Рис. 1. Результаты СЭМ для композитной мембраны А, состоящей из 80% кристаллов CsHSO4, синтезированных
при помощи прекурсора Cs2CO3 Fig. 1. SEM images of composite membrane A containing 80% CsHSO4 crystals synthesized with Cs2CO3 as the precursor
Membrane B, Fig. 2, is composed of CsHSO4, PTFE and SiO2. From the SEM pictures, there is a visible difference in the appearance of the composite with a more granular structure when heated as opposed to the smooth amorphous appearance of the hydrated specimen.
From the SEM images, Fig. 3, there is a visible difference in the appearance of the composite with a more granular structure when heated above 405 K to the super-ionic conductive phase. Water loss and other solvents causing the void-like porous appearance could cause the formation of the porous structure.
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Рис. 2. Мембрана В прошла температурный цикл во время измерения сопротивления при АЧХ, максимальная температура достигала 444 К Fig. 2. Membrane B has undergone a thermal cycle while testing the resistance in FRA and the highest temperature reached was at 444 K
Рис. 3. Результаты СЭМ для мембраны С при 1000х увеличении Fig. 3. SEM image of membrane C at 1000x magnification
Рис. 4. Поперечный разрез мембраны D при 1000х увеличении Fig. 4. Membrane D shows a side view of the membrane at 1000x magnification
The membrane, Fig. 4, has also run a complete heating cycle in FRA and the highest temperature reached was a 444 K. The composition of both membranes is 57.5% CsHSO4, 36.5% PTFE and 6% SiO2.
Sulfates have strong bands in the region 1080 to 1150 cm-1 (Fig. 5), and also a medium strong band between 580 and 670 cm-1. There are very small or no differences in the structure after heating to 353 K. The sulphate band region includes peaks of similar resolution and strength to those analysed at room temperature. Hydroxyl bands typical in the region 3700-3100 cm-1 are absent. This indicates possible water or solvent residue loss on drying at low temperature. Slight peaks between 1410 and 11450 cm-1 suggest the incomplete conversion of Cs2CO3 in the presence of H2SO4 to Cs2SO4, as this band area is indicative of carbonates.
Рис. 5. Сравнительный анализ ИК спектра для различных составов Fig. 5. IR spectral analysis comparing different compositions
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The heating rate was gradual and no visible change to the membrane surface was observed. Membranes A, B, C and D with CsHSO4 were also viewed by the higher magnification 'zoom in' feature to observe the infrared peaks in greater detail. After heating membrane A to 378 K the peaks seem sharper and more defined. The structure has not changed but the loss of water and other solvents has enhanced the resolution compared to the infrared analysis at room temperature. Composites containing CsHSO4 and PTFE were prepared to investigate the mechanical and conductivity dependence on thermal influences and served as an isolation technique to identify the band regions for SiO2 peaks which would obviously be absent in the CsHSO4-PTFE composites and only present in the CsHSO4-PTFE-SiO2 composite. The thermal stability of the composites was thus investigated by observing the change in structure using infrared analysis. These were membranes E, F, G and H containing varying volume fractions of PTFE. For the tri-component composites namely membranes A, B, C, and D have a band region 1100-1300 cm-1 with two peaks indicative of polymer presence. This is further substantiated by membranes F and G where the 70 and 50% PTFE respective fractional percentage clearly shows distinct and defined peaks in the 1100-1300 cm-1 band width region. These peaks are not observed in the pure CsHSO4 spectrum. The presence of PTFE contributes to the diminished size of the water band. SiO has the opposite effect by increasing the size of the water band when the SiO fractional percentage is increased. The definition of the peak increases as the fractional composition of the SiO2 in the composite increases along with CsHSO4. Between 700 and 1250 cm-1, the silica species may be observed. Specifically Si-O bond stretching is found in the range 700-1100 cm-1 but this peak is also present in composites where SiO2 is absent. Therefore SO4 bonds may be represented by the identified peaks [49, 50]. The structure seems very stable with predictable changes in structure depicted in the transmittance infrared diagrams with varying compositions. The fractional composition is clearly seen with the trend in the size of the peaks, which are proportional to the concentration of the different components. The series of thermal cycles performed on the membranes have no significant effect on the structure of the individual components in the composite. Similar trends in the infrared spectroscopic analysis are observed in samples that have undergone thermal cycles as compared to those that have not been subjected to the FRA. FRA was performed simultaneously with thermal cycling. The pure CsHSO4, which is soluble in water and dilute methanol (1M), has a pH of 4. The pH was measured when the solution was at room temperature. A 0.1% solution sample was used. The process by which the pH was measured included varying the temperature and thereafter bringing the solution temperature back to room temperature once the reaction between the membrane and water has taken place. This showed the influence if any, of the thermal effect on the degradation
of the membrane with the subsequent release of protons into solution. Membrane composites with 0.8 volume fraction CsHSO4 and 17% PTFE, showed a greater dissolution in the dilute 1M methanol with up to 30% of the solid membrane dissolving at room temperature; whereas a membrane with a volume fraction of 34% PTFE content experienced 19% dissolution at similar temperatures. In both instances the stirring rate was 300 rpm. For the membrane with 17% volume fraction PTFE, the pH change was not significant within the first hour of continuous stirring.
Composites containing a higher volume fraction PTFE do not dissolve as easily as those with a lower PTFE content. This is evident in membrane A with 17 wt% PTFE (Fig. 6) and membrane C with 40 wt% PTFE. The composites were exposed to hydrogen gas for varying periods of time. The degradation rate was recorded over time by weighing the composites at different intervals.
Рис. 6. Соотношение pH и времени растворения Fig. 6. The pH and dissolution time relationship
Рис. 7. Соотношение между потерей веса мембраны
в масс% и объемной долей PTFE Fig. 7. The relationship between the membrane wt% lost and PTFE volume fraction
Once again, Fig. 7, the composites containing the higher concentrations of PTFE was not as easily degraded.
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The reaction 2CsHSO4 + 4H2 ^ Cs2SO4 + H2S + + 4H2O outlines the degradation of cesium hydrogen sulphate and formation of hydrogen sulphide gas. The evidence suggests the minimizing of the degradation process as the PTFE concentration is increased. Temperatures above 423 K where the super-ionic phase already reached could possibly encourage the formation of hydrogen sulphide and thus the instability of the membranes. The unstable membrane could enhance the conductivity by adopting the semi-crystalline ionic structure. The volume fraction of the insulator substance reduces the conductivity as previously recorded when the concentration thereof is increased. There was a significant difference in mass lost due to temperatures above 423 K.
The CsHSO4 membrane improved its water resistant properties with the addition of PTFE. This was achieved with insulator concentrations as low as 0.17 volume fraction of the composite and 0.13 silica volume fraction. The superionic phase transition was attained at approximately 405 K once silica previously preheated to 1173 K was included in the composition whereas without silica the phase transition temperature was recorded at approximately 413 K. Conductivity in the superionic phase recorded at 10-2 S/cm similar to that of the pure salt. Further improvement in the membrane stability with respect to water was noted at higher PTFE volume fractions where water was effectively prevented from being reabsorbed after heating. Stability testing showed that CsHSO4 membranes with higher PTFE content were able to maintain a constant mass for five days unprotected from exposure to the atmosphere. The pH of the test solution did not drop as rapidly in the composites when compared to the pure salt. The acidity of the test solution gives an indication of immiscibility and mechanical strength improvement of the membrane. Infrared analysis revealed the structural resistance to damage when heated to temperatures as high as 443 K. High temperature treated composites gave spectra with more defined peaks. This was possibly due to the water loss after 373 K and further removal of solvents that improved the resolution.
Conclusion
Future research and development will be conducted using different group I and II metals in place of cesium for the composite formations. IR and Raman spectroscopic studies will be used to determine the reversible phase transitions and corresponding temperatures. Conductivity analysis will be performed and compared to that of cesium containing composites. The use of different group 1 and 2 metals could possibly produce a stable solid electrolyte that is less soluble than previously manufactured membranes containing cesium. An in-depth study to compare the particle size distribution and uniformity of the composites with respect to conducting and insulating materials will be performed.
References
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