A review of advanced proton-conducting materials for hydrogen separation

Progress in Materials Science 74 (2015) 1–50

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A review of advanced proton-conducting materials for hydrogen separation Zetian Tao a,⇑, Litao Yan d, Jinli Qiao a,b,⇑, Baolin Wang a, Lei Zhang c, Jiujun Zhang a,c,⇑ a Key Laboratory for Advanced Technology in Environmental Protection of Jiangsu Province, Yancheng Institute of Technology, Yancheng, Jiangsu, China b College of Environmental Science and Engineering, Donghua University, Shanghai 201620, China c Department of Chemical & Biochemical Engineering, University of British Columbia, Vancouver, BC V6T 1W5, Canada d Department of Chemical Engineering, University of Missouri–Columbia, Columbia 65211, USA

a r t i c l e

i n f o

Article history: Received 3 March 2014 Accepted 6 March 2015 Available online 29 April 2015 Keywords: Hydrogen separation Proton and electron conduction Ceramic materials and membranes

a b s t r a c t This paper provides a comprehensive overview of developments and recent trends in H2 separation technology that uses dense proton– electron conducting ceramic materials and their associated membranes. Various proton–electron conducting materials and their associated membranes are summarized and classified into several important categories, such as Ni-composite proton-conducting materials, as well as tungstate-based, BaPrO3-based, LaGaO3-based, and niobate/tantalite composite metal oxide-based ceramic materials/membranes. Various membrane designs, including asymmetric ceramic membranes (supported and self-supported) and surface-modified membranes, are also reviewed. Several important properties of ceramic materials and membranes, such as proton and electron conductivity and performance (i.e., H2 transport flux and lifetime stability), are also discussed. To highlight the technical progress in this area, all possible ceramic materials and associated membranes are summarized, along with their properties and performance, to help readers quickly locate the information they are looking for. Based on this review, several challenges hindering the maturation of this technology are analyzed in depth, and possible research directions for overcoming these challenges are suggested. Ó 2015 Elsevier Ltd. All rights reserved.

⇑ Corresponding authors at: Key Laboratory for Advanced Technology in Environmental Protection of Jiangsu Province, Yancheng Institute of Technology, Yancheng, Jiangsu, China (J. Qiao, J. Zhang). Tel.: +86 515 88298926; fax: +86 515 88298927. E-mail addresses: [email protected] (Z. Tao), [email protected] (J. Qiao), [email protected] (J. Zhang). http://dx.doi.org/10.1016/j.pmatsci.2015.04.002 0079-6425/Ó 2015 Elsevier Ltd. All rights reserved.

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Contents 1.

2.

3.

4. 5. 6.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 1.1. Importance of hydrogen fuel in practical applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 1.2. Hydrogen separation membrane technologies and their associated materials . . . . . . . . . . . . . . . . 3 1.3. Advanced proton-conducting ceramic membranes and their associated materials . . . . . . . . . . . . 5 1.4. Necessity of and perspectives on developing advanced proton–electron conducting ceramic materials and their associated membranes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6 Fundamental understanding of the H2 separation process and material performance through both theoretical and experimental analysis. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 2.1. Hydrogen separation process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 2.2. Theoretical understanding of H2 transport mechanisms . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9 2.3. Theoretical and experimental analyses to investigate the hydrogen transport mechanisms in ceramic membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 10 Development of dense proton–electron conducting ceramic materials and their associated membranes for hydrogen separation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13 3.1. Synthesis, characterization, and performance of single-phase ceramic mixed materials and their associated membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 14 3.2. Synthesis, characterization, and performance of Ni and proton-conducting material composed materials and their associated membranes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16 3.3. Synthesis, characterization, and performance of fluorite-structure materials and their associated membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17 3.4. Synthesis, characterization and performance of tungstate-based materials and their associated membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20 3.5. Synthesis, characterization, and performance of BaPrO3-based materials and their associated membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 22 3.6. Synthesis, characterization, and performance of LaGaO3-based materials and their associated membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24 3.7. Synthesis, characterization, and performance of niobate/tantalite composite metal oxide-based materials and their associated membranes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24 3.8. Synthesis, characterization, and performance of ceramic asymmetric membranes . . . . . . . . . . . 25 3.9. Synthesis, characterization, and performance of ceramic hollow fibers . . . . . . . . . . . . . . . . . . . . 30 3.10. Modifying the surface of a hydrogen separation membrane . . . . . . . . . . . . . . . . . . . . . . . . . . . . 32 3.11. Stability of hydrogen separation materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 32 Ceramic membrane-based H2 separation systems design and fabrication . . . . . . . . . . . . . . . . . . . . . . . . 36 Challenges of H2 separation using dense proton–electron conducting ceramic materials/membranes . 38 Summary and proposed research directions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 40 Acknowledgments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 44 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 44

1. Introduction In today’s world, clean-energy science and technology, which include energy storage and conversion, are the most important R&D topics for the sustainable development of human society, and are also becoming the most critical elements in overcoming fossil fuel depletion and global pollution. The rate of fossil fuel depletion is accelerating due to increased energy consumption in people’s daily lives and in industry processes. For example, it is expected that our present energy consumption rate will double by 2050 [1]. In terms of environmental impacts, fossil fuel consumption contributes air emissions of 65 trillion tons of carbon (in the form of CO2), as well as considerable quantities of other polluting gases such as SOX, NOX, VOCs, and sulfur-containing substances [2–4]. Thus, to maintain human life and development on earth, clean and sustainable energy sources must be urgently explored and utilized without delay. Among various potential clean and sustainable energy sources, hydrogen (H2) has been recognized as one of the most appealing because it can be produced mainly

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through the reforming of sustainable fuels, such as methanol and biomass, and through water electrolysis using clean and more sustainable energy sources, such as hydroelectric, solar, wind, and geothermal. 1.1. Importance of hydrogen fuel in practical applications It has been recognized that the energy density of H2 is higher than any other fuel’s. This energy can be released through direct combustion and/or fuel cell conversion to produce power without any emission at the source. For example, during the direct combustion of H2 in air or the chemical reaction of H2 with air, the only emission products are water and heat. However, because of both technical difficulties and high costs, a ‘‘Hydrogen Economy’’ has begun to look practicable only in recent years, despite having been proposed in the 1970s. To realize this Hydrogen Economy, major industrialized countries have initiated a wide range of fundamental and applied research to develop the commercial use of hydrogen as a primary energy carrier [5]. For example, in the United States, the Hydrogen Technical Advisory Panel (HTAP) has advised the implementation of the Department of Energy’s (DoE) Hydrogen Program to investigate the economic, technical, and environmental consequences of developing hydrogen energy systems. Their research progress indicates that the electricity produced from hydrogen will become one of the nation’s primary energy sources in the 21st century [6]. Numerous studies have shown that it is highly feasible to use H2 as a fuel, including by using H2 conversion devices to replace conventional power generators—such as combustion engines or even large batteries—in cars [7–9], buses [10], forklift trucks (FLTs) [11], submarines [12], backup power units, and power plants [13,14]. The major application areas of H2 can be summarized as follows: (1) Hydrogen-powered cars: Numerous industry partnerships have evolved to advance automotive electrification and reduce carbon emissions. (2) Hydrogen-powered buses: In various cities, including San Francisco and Berlin, H2-powered public bus services have been initiated. (3) Hydrogen-powered forklift trucks: Market leaders such as Walmart and BMW rely on fuelling technologies from Linde to decarbonize internal logistics and materials handling services. (4) Hydrogen-based backup and off-grid power: Hydrogen can be stored as a backup power supply. (5) Hydrogen-powered submarines: Hydrogen-powered submarines can remain under water much longer than conventional submarines and also are much quieter. Various navies have used hydrogen gas and fuelling stations for submarines. (6) Hydrogen-powered ships: Hydrogen-powered ferry boats can be used to reduce the high emissions they presently generate. 1.2. Hydrogen separation membrane technologies and their associated materials As mentioned above, H2 can be produced from both water electrolysis using clean/sustainable energy sources [15,16] and from fossil fuels [17–19]. However, due to the technological challenges still involved in H2 production using clean/sustainable energy sources, such as insufficient maturity/reliability and high cost, production of H2 in large quantities is mainly done through the reforming of natural gas and the gasification of coal, which remain economically favorable methods at this point. Unfortunately, the H2 produced from fossil fuels is normally in the form of gas mixtures containing H2, CO2, CO, and other minor impurities, such as O2, H2O, N2, SOX, NOX, VOCs, and sulfur-containing substances, making the gas not ready for direct use. Therefore, it is necessary to separate and purify H2 from these mixtures using low-cost, simple, highly energy efficient technologies. In addition, other gas components that are separated out, such as CO2 and CO, can be collected, stored, and utilized for other purposes. In this way, the potential emissions from these processes are controllable and manageable, which is not the case in the direct combustion of fossil fuels, during which environmentally detrimental substances are freely emitted into the air, causing multiple pollution issues.

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To produce H2 through fossil fuel reforming and gasification, H2 in the reformed gas mixtures has to be separated to obtain purified H2. The current commercially available technology is the pressure swing adsorption (PSA) process [20–22]. Although PSA technology has a fast cycle and rapid adsorbent regeneration during the continuous production of purified hydrogen, the disadvantages seem to be high energy consumption during the cryogenic distillation step and the need for high pressure during additional compression procedures. To overcome the challenges of low energy efficiency and high cost in H2 separation, several membrane separation technologies have been explored and technically verified. Normally, membranes are selective barriers that allow the passage of certain components of a gas mixture and block others. In H2 separation, a membrane selectively allows H2 to pass through from the side containing the feed mixture to the other side, which contains only the separated H2. In this way, H2 is purified. The membranes explored in the literature can be summarized as follows: (1) Non-porous polymeric membranes [23,24]: The majority of industrial membrane processes for gas separation utilize non-porous polymeric membranes because of their reasonable gas selectivity, good mechanical properties, and low-temperature operation (below 110 °C). Gases are selected based on the membrane’s permeability (the product of diffusivity and solubility coefficient) using a solution–diffusion mechanism. The gases from the inlet feed flow dissolve in the polymer, diffuse to the other side, and desorb to the outlet flow. A common parameter characterizing the ability of a polymer to separate two gases (e.g., a and b) is called the ideal selectivity [25]:

aa=b ¼

Pa Pb

ð1Þ

where Pa and P b are the permeability of gas components a and b, respectively. Although these polymer membranes have been commercially viable, their drawbacks—such as insufficient selectivity, low flux, low tolerance to SOX, CO2, and HCl, as well as high mechanical deformation sensitivity—pose challenges in terms of membrane lifetime. (2) Dense metallic membranes [26,27]: Normally, dense metallic membranes have extremely high selectivity for H2 separation. For example, Pd-based and Pd alloy-based membranes allow only H2 to permeate, producing hydrogen with an impurity of less than 1 part per billion [28]. The separation is based on a solution–diffusion mechanism, in which the H2 molecule dissociates into atoms and diffuses through the metal, recombining into molecules on the other side of the membrane. These metallic membranes normally work at 300–600 °C and produce high-purity H2. However, during thermal cycling, irreversible change can take place in the Pd lattice structure as a result of the repeated transitioning between phases, resulting in membrane degradation. Furthermore, the high cost of palladium limits these membranes’ application in large-scale H2 production. Research needs to be directed toward reducing Pd usage to make this technology more cost-effective and sustainable. (3) Carbon membranes [29,30]: Carbon membranes are porous, and their H2 separation mechanisms are based on surface diffusion or molecular sieving. Surface diffusion refers to the process by which molecules migrate along the porous surface; this may occur in parallel with Knudsen diffusion and thereby increase the selectivity by different adsorbing abilities. Molecular sieving happens when the pore size is so small that the larger molecules are sieved out. Although using carbon membranes to achieve molecular sieving is a promising approach, membrane brittleness is a major drawback. Furthermore, the performance can degrade in the presence of strongly adsorbing vapors, such as H2S, NH3, or chlorofluorocarbons [31]. Performance degradation can also occur with decreasing temperature, as the pore size shrinks. So far, carbon membranes are still not economically feasible or sufficiently commercially available. (4) Dense ceramic membranes [32–35]: In general, dense ceramic membranes have much higher H2 selectivity than microporous membranes. These ceramic materials can be made so as to have proton and electron conductivities that allow H2 to permeate. The working process can be summarized thus: H2 adsorbs onto the membrane surface and dissociates into protons and electrons, both of which then diffuse to the other side of the membrane surface, where they recombine to form the H2 molecule. Theoretically, the selectivity is comparable to that of Pd

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or Pd alloy membranes, since only hydrogen permeates. However, these ceramics are much less expensive than Pd-based membranes. These ceramic membranes also have the advantage of relatively high durability and stability in steam containing CO, CO2, and H2S [36]. Furthermore, when a ceramic membrane is operated at high temperatures, the H2 migration kinetics is fast and able to produce industrially significant flux. Compared to traditional membrane separation technology, dense ceramic membranes consisting of high-temperature proton and electron conductors can achieve excellent stability as well as high hydrogen selectivity, and at a suitable cost [37–40]. Using dense proton–electron conducting ceramic membranes, H2 is easily obtained in situ from steam reforming, partial oxidation of carbon fuels, and gasification of coal and biomass at high temperatures. These processes can be combined with the water–gas shift reaction to convert the CO in a reforming gas, such as syngas, into more hydrogen and carbon oxide, which can potentially decrease the cost of H2 production. With all the above advantages, dense proton–electron conducting ceramic membrane technology appears to be the most promising approach for hydrogen separation in real-world applications. In summary, to achieve H2 separation for industrial applications, membrane technologies have the advantages of low cost and low energy consumption [16]. For example, H2 separation using membranes can operate continuously using potentially less than half the energy required for the PSA process [41]. Furthermore, the membrane separation process is also more spatially flexible, offering H2 purification at different scales. Among the four kinds of hydrogen separation membranes (polymeric, metallic, carbon, and ceramic [42]), only polymer membranes have been in commercial use to any considerable extent at the current state of technology. The other three advanced inorganic materials-based membranes are still under development but continue to attract significant engineering interest; notably, some have been used in pre-commercial applications. Of these three categories, dense proton–electron conducting ceramic membranes are the most promising due to their many advantages. This review paper will focus upon these membranes and their associated ceramic materials.

1.3. Advanced proton-conducting ceramic membranes and their associated materials As discussed above, advances in H2 separation technologies, particularly those based on proton– electron conducting ceramic membranes, have shown great promise for improving energy efficiency and reducing cost. Indeed, it has been demonstrated that these membranes, consisting of high-temperature proton-conducting and electron-conducting phases, seem to have reasonable stability high H2 selectivity and low cost. Normally, many such membranes are operated from 300 to 900 °C. At 900 °C, a hydrogen flux as high as 20 cm3 min1 cm2 was achieved in ambient hydrogen pressure, running for 190 h, with no noticeable degradation in the syngas atmosphere [43]. However, the energy efficiency and cost of these membranes are strongly dependent on the materials’ properties, such as morphology/nanostructure and composition. Therefore, materials development is a critical approach to achieve technological breakthroughs in this area. In general, ceramic-based membranes are often made of polycrystalline ceramic materials, particularly perovskites and/or metals (such as palladium and nickel), which allow specific gas species to permeate or sorb in the dense material. Current technology includes two types of ceramic H2 separation membranes: (1) the metal–ceramic composite membrane, in which the metal is used as the electron-conducting phase and the ceramic oxide serves as the proton-conducting phase—e.g., membranes based on Ni–BaZr0.1Ce0.7Y0.2O3d [39,44,45] or Ni–La0.5Ce0.5O2d [40,46] and (2) the single-phase ceramic oxide, which simultaneously transports protons and electrons—e.g., membranes based on doped-ACeO3 (A = Ba, Sr). Doped SrCeO3 and BaCeO3 perovskite membranes have been extensively researched because of their excellent mixed conductivity at high temperatures, as electron conductivity can be obtained by doping the B site with a multivalent cation [47–49]. Unfortunately, BaCeO3 and SrCeO3 are chemically unstable, easily reacting with CO2 and H2O [50,51]. Many elements have been explored as doping agents to improve the chemical stability of barium cerates. In addition, some novel proton-conducting materials have been developed, such as LaNbO4, BaCa1.18Nb1.82O9d,

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and Ln6WO12 [52–54]. However, these novel materials are hampered by low conductivity, precluding their further development. With respect to membrane design, it is worth mentioning that besides membranes with symmetrical structures, those with asymmetrical structures have also been explored to separate out many gases, including H2 and O2, by improving permeability [55]. For example, asymmetrically structured ceramic H2 separation membranes have been studied to reduce their thickness and resistance [56,57]. Regarding material requirements for dense proton–electron conducting ceramic membranes, stability is the first consideration. Syngas includes hydrogen, carbon dioxide, and moisture, and some high-temperature proton conductors (HTPCs) are not stable in such an acidic, moist atmosphere [58–61]. Therefore, the membrane materials must have: (1) very high chemical and thermal stability (enough durability for real operating environments), (2) high hydrogen permeation flux rates, (3) low parasitic power requirements, (4) low materials cost, and (5) low membrane fabrication cost. Generally, pure proton-conducting materials cannot be used for hydrogen separation because they have insufficient electron conduction [62,63]; an auxiliary external electrical circuit has to be applied across the membrane to create a hydrogen chemical potential gradient so as to drive the flow of hydrogen. Fortunately, a dense ceramic membrane consisting of mixed proton and electron conduction channels (MPECs) offers fast and selective H2 transport, and provides a simple, low-cost, efficient way to separate hydrogen from gas streams at high temperatures. Because the MPEC membrane transports both protons and electrons at the same time, the H2 separation process can be carried out in a non-galvanic mode without using an external electrical circuit. Therefore, developing MPEC materials with such unique transport properties is essential to H2 separation technology. In summary, there are three ways to fabricate MPECs: (1) prepare membranes containing an electron-conducting metal phase and a proton-conducting ceramic phase, (2) dope a special element into the pure conducting phase to improve the electron conductivity, and (3) develop novel MPECs, such as La2Ce2O7 [40,46] and LnWO6 series oxides [64–70].

1.4. Necessity of and perspectives on developing advanced proton–electron conducting ceramic materials and their associated membranes Although recent years have produced extensive efforts to develop proton–electron conducting ceramic materials and their associated membranes for H2 separation, several challenges must be overcome before commercialization. These technical challenges include: (1) insufficient chemical stability of the materials in acidic gas atmospheres at high temperatures, (2) insufficient proton and electron conductivities of the materials for high H2 flux production, (3) insufficient H2 dissociation/reassociation reaction rates at the membrane surfaces, (4) insufficient mechanical strength and durability for long-term operation, (5) insufficient thermal stability for long-term operation, and (6) lack of fundamental understanding of material/membrane performance and degradation. Although these challenges currently hinder the commercialization of this membrane technology, it is still considered the most promising for H2 separation because of its relatively high stability, excellent H2 selectivity, high energy efficiency, and low cost. Ongoing effort is definitely needed to achieve breakthroughs and push this technology forward to commercialization. The past two decades have seen the publication of a huge number of research and development papers on H2 separation technology, signaling both its importance and the rapid progress being made in this area. Although it is impossible here to adequately address all aspects of this research as communicated in the literature, a review paper covering key features of new and advanced proton- and electron-conducting ceramic materials and their associated membranes for H2 separation would be of high value to the various interested research communities. This paper reviews recent progress in the research and development of H2 separation technology, with a focus on the most widely explored new materials and their associated membranes, specifically emphasizing certain novel proton-conducting materials. We strongly hope that researchers, students, industry professionals, academics, and hydrogen manufacturers, as well as related professionals concerned with energy conversion technologies, will be very interested in, and benefit from, this review paper.

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2. Fundamental understanding of the H2 separation process and material performance through both theoretical and experimental analysis In the study of hydrogen separation, it is generally recognized that in addition to new material synthesis, characterization, and performance validation, it is critical to gain a fundamental understanding of material/membrane structural effects and functional mechanisms in order to achieve performance optimization and develop new designs for proton–electron conducting materials and their associated membranes. This fundamental understanding can be attained through theoretical approaches—such as modeling and associated calculations—as well as performance measurements and failure mode diagnosis. In this section, we summarize the progress that has been made to date in gaining a fundamental understanding of H2 separation materials/membranes through theoretical analysis and experimental diagnosis. 2.1. Hydrogen separation process The H2 separation process using a membrane (specifically, a dense proton–electron conducting ceramic membrane (DPECCM)) under a concentration gradient is schematically illustrated in Fig. 1. In this MPEC membrane, H2 moves from the high partial pressure side, P0H2 , to the low partial pressure side, P00H2 . Overall, the process of H2 separation through a DPECCM membrane mainly involves two steps: (1) H2 is first absorbed into the surface of the DPECCM, then dissociates into protons and electrons; (2) protons and electrons diffuse together to the other side of the DPECCM surface, where they reassociate to form molecular H2 again [71]. To obtain a practical H2 separation rate, both surfaces of the membrane should have catalytic effects for H2 dissociation and reassociation. As shown in Fig. 1, the overall process can be summarized in the following reactions:

H2 ! H2;ads

ðAdsorption on the membrane surfaceÞ

H2;ads ! 2Hads

ðIÞ

ðDissociation into adsorbed atomic hydrogen on the membrane surfaceÞ ðIIÞ

Hads ! Hþads þ e ðFurther dissociation into proton and electron on the membrane surfaceÞ ðIIIÞ Inlet feed gas

Outlet H2 gas

H2 + CO2 + CO + H2O + Others 2H+

H2 => 2H+ + 2e-

PH' 2

2H+ + 2e- => H2 2e-

PH" 2

H2 + CO2 + CO + H2O + Others

Dense proton-electron conducng ceramic membrane (DPECCM) Fig. 1. Scheme of H2 separation process using a dense proton–electron conducting ceramic membrane (DPECCM). Note that to increase the dissociation/reassociation processes at the membrane surfaces, the surface in contact with the inlet feed gas can be coated with a thin layer of catalyst (the oxidation catalyst layer), as can the surface in contact with the outlet H2 gas (the reduction catalyst layer).

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Diffusion

Reassociation

Diffusion

ðHþads þ eÞS0 ! ðHþads þ eÞBM ! ðHþads þ eÞS00 ! ðH2 ÞS00 ! ðH2 ÞG ðDiffusion of the dissociated proton and electron through the membrane to become molecular H2 ; where S0 is the membrane surface at the gas inlet side; BM is the bulk membrane; S00 is the membrane surface at the H2 outlet side; and G is the gas; respectivelyÞ:

ðIVÞ

As mentioned above, the process of hydrogen permeation is primarily controlled by two processes: the surface process and bulk diffusion. If the permeation process is controlled by bulk membrane transport, decreasing the thickness of the membrane can boost the transport rate; if it is controlled by the surface process, modifying the membrane’s surface may increase the hydrogen permeation flux. For example, it was reported that the hydrogen permeation flux increased significantly when a Pd-modified surface was used on a Ni–BCY membrane, due to the excellent ability of Pd to catalyze H2 dissociation [72]. For the bulk control process, the transport of protons through the bulk membrane is the rate-determining step if the electron conductivity is much higher than the proton conductivity. According to the chemical/electrochemical potentials and conductivities of both protons and electrons within the membrane, the H2 flux ðjH2 Þ during the separation process can be written as Eq. (2) [73]:

jH 2 ¼

P0H2 rHþ rel ln P00H2 4F 2 L rHþ þ rel RT

!

ð2Þ

This is the Wagner equation, where R is the universal gas constant, T is the temperature, F is the Faraday constant, L is the thickness of the ceramic membrane, rHþ is the proton conductivity within the membrane, rel is the sum of the electron and hole conductivities, P 0H2 is the H2 pressure at the inlet gas chamber as indicated in Fig. 1, and P 00H2 is the H2 pressure at the outlet H2 gas chamber. Note that rHþ and rel are both independent of hydrogen pressure in this equation. Eq. (2) shows that the hydrogen flux is inversely proportional to the thickness of the membrane; so, the thinner the membrane, the faster the H2 flux will be if the selectivity of the membrane permeability for H2 is high. An effective method to obtain high hydrogen permeation flux is to reduce the thickness of the membrane, thereby lowering the bulk diffusion resistance. This can be done by supporting a thin proton–electron conducting ceramic membrane onto an inert porous support using various technical methods, such as co-pressing [45,57], spin coating [74,75], dip coating, spray coating, sputtering [76], sol–gel methods [77], pulsed laser deposition (PLD), and chemical vapor deposition (CVD) [78–80]. From Eq. (2) it can also been seen that the H2 flux is controlled by the proton and electron conductivities, the temperature, and the hydrogen pressure on both sides of the membrane. Increasing the temperature and pressure at the inlet gas side, or decreasing the membrane thickness and pressure at the outlet gas side, can effectively improve the H2 flux. Of course, increasing the dissociation/reassociation rates at the membrane surfaces can also improve the H2 flux. If the proton conductivity within the membrane is much smaller than the electron conductivity, that is, rHþ  rel , Eq. (2) can be rewritten as:

jH 2 ¼

RT 4F 2 L

rHþ ln

P0H2

!

P00H2

ð3Þ

In this case, developing highly conductive proton-conducting ceramic membranes is fairly important for achieving high H2 flux, although doing so is challenging in the current state of technology. If the proton conductivity within the membrane is much larger than the electron conductivity ðrHþ  rel Þ, Eq. (2) can be rewritten as:

jH 2 ¼

RT 4F 2 L

rel ln

P0H2 P00H2

!

ð4Þ

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In this case, developing ceramic membranes with high electron conductivity becomes critical in achieving high H2 flux during the separation process. Normally, it is hard to obtain both high proton and high electron conductivities, although some attempts have been reported [81]. For example, metal-based Ni–BaCe0.8Y0.2O3d membranes were explored to achieve mixed conductivities, and ceramic–ceramic materials such as doped ceria-based BaCe0.8Y0.2O3d were also reported [82,83]. It may be that the surface diffusion and dissociation processes are the control steps in H2 transport. Surface modification has proven to be an effective way of speeding up these surface processes. 2.2. Theoretical understanding of H2 transport mechanisms The major driving force behind H2 transport within a DPECCM is the pressure difference across the membrane. This difference can induce a Nernstian potential (E) across the membrane, according to the following equation [84]:

E¼

P00H2 RT ln nF P0H2

! ð5Þ

where R, T, P0H2 , and P 00H2 have the same meanings as in Eq. (2), and n is the electron number in the H2 dissociation/reassociation reaction (H2 $ 2Hþ þ 2e ), equal to 2. The value of this potential difference is in the range of 50–500 mV, depending on the temperature and H2 pressure difference across the membrane. Under this potential difference, the dissociated protons and electrons move from the inlet gas side to the outlet gas side, and the transport rate is determined by the magnitude of the potential difference. During H2 dissociation/reassociation at the ceramic membrane surfaces, the oxygen vacancy and the oxygen anion sites inside the surface oxide materials (i.e., the ceramics) play a critical role [71]. Normally, the feed gas mixture contains moisture; the oxygen vacancies, expressed in the literature as Vo , will react with water to fill the lattice positions with oxide ions (expressed as Oox ), forming proton-retaining oxide ions (also called proton charge carriers, expressed as OHo ) at normal lattice sites:

H2 O þ Vo þ Oox $ 2OHo

ðVÞ

Reaction (V) indicates that more oxygen vacancies in the crystal structure of oxides will induce more proton-retaining oxide ions, leading to higher proton conductivity. This equation also indicates the importance of water in proton conductivity. Furthermore, H2 in the inlet feed gas can also directly react with the oxide ions to produce proton-retaining oxide ions:

H2 þ 2Oox $ 2OHo þ 2e OHo

ðVIÞ

where is the major intermediate for proton conduction within the membrane, under the driving force of potential difference. It is generally accepted that the proton inside the OHo species can transfer between oxygen ions in the normal position, producing proton conductivity [85]. In this case, the distance between oxygen ions in the crystal lattice of the ceramic material should play a significant role in proton conductivity—that is, when the vibration distance between the oxygen ions increases, the proton conductivity will increase. This has been observed by measuring the activation energy for proton conductivity as a function of oxygen ion distance [86]. The magnitude of the oxygen ion distance in a crystal structure is strongly dependent on the type of materials and their associated crystal structures. For example, the most popular material and structure for H2 separation membranes is ABO3 perovskite structure, shown in Fig. 2. B represents a transition metal cation element (Mn, Cr, Fe, etc.) or a rare earth element such as Ce or Zr, which has a high positive valence, forming a six-coordinate octahedron with neighboring oxygen ions and occupying the center of the perovskite structure. The larger ionic radius of the A cation results in lower positive valences, and its coordination number can be as high as 12. The distance between the oxygen ions can be changed by doping a trivalent element into the structure, as shown in Fig. 2, leading to significant changes in the structure’s proton conductivity [87]. After a trivalent element M is doped in, the formula of these high-temperature, proton-conducting perovskite oxides can be written as

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A

O

B

Fig. 2. ABO3 perovskite structure.

AB1xMxO3. Doping a trivalent M cation into the B site can increase the vibration distance between the oxygen ions and also create more oxygen vacancies, leading to higher proton conductivity [88–90]. As indicated in Reaction (V), oxygen vacancy and the distance between the oxygen ions play important roles in proton conductivity. However, the mixed conductivities of protons and electrons may make the transport mechanisms more complicated. Generally speaking, the oxygen vacancy concentration, the distance between the oxygen ions at the point of the crystal structure, hydroxide defects, electron holes, water vapor in the feed gas, and H2 pressures will reach equilibrium at a certain temperature, co-affecting the proton and electron conductivities of the ceramic membrane. To identify the roles of the major factors, extensive theoretical and experimental analyses have been carried out, as reported in the literature [91]. Progress in these areas will be briefly discussed in the following subsections. 2.3. Theoretical and experimental analyses to investigate the hydrogen transport mechanisms in ceramic membranes As discussed above, after protons and electrons are formed through H2 dissociation at the membrane surface, the protons are incorporated into the bulk membrane, then move to the other side under the driving force of potential difference. The question is: what are the mechanisms of proton transport in the crystal oxide structure of the ceramic membrane? Fabbri et al. [92] thought that a proton would interact with the electron cloud of a neighboring oxygen ion without occupying the regular interstitial position in the lattice; protons would then bind to oxygen ions to form hydroxide defects. Although there is still no agreement on the proton transfer mechanisms of ceramic materials such as perovskite, two have been proposed in the literature: the vehicle mechanism, and the Grotthuss—or hopping—mechanism. According to the vehicle mechanism, as schematically represented in Fig. 3 [93], proton transport is conducted by the migration of OH, and the oxygen in the hydroxyl ion functions as the vehicle. This

Fig. 3. Schematic diagram of the vehicle mechanism. (Reprinted with permission from Ref. [93]. Copyright 2005 ACS.)

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mechanism, which can be thought of as the diffusion of H3O+, seems very simple and quite clear, and is also supported by the fact that proton conductivity is strongly dependent on the dopant concentration, which can produce oxygen vacancies. For example, Maekawa et al. [94] used 1H NMR spectroscopy to measure the relationship between proton concentration and dopant concentration in BaCe1xYxO3d-based materials; their results are shown in Table 1. At the same time, they also measured the conductivity of BaCe1xYxO3d (x < 0.1) under 1 kPa water. As shown in Fig. 4, the conductivity increases as the dopant concentration increases. However, the proton conductivity decreases when the dopant concentration reaches a certain point, due to defects. Künstler [96] studied the relationship between proton concentration and dopant concentration in BaCe1xInxO3d (x < 0.25) material and found that the conductivity increased with the dopant concentration up to x 6 0.2. All these results support the proposed vehicle mechanism of proton conduction. Furthermore, it is believed that this mechanism may play a leading role at high temperatures due to high activation energy. Unfortunately, this model has difficulty explaining the H/D isotope effects on proton conductivity, and the high diffusivities of protonic defects in oxides that are closely packed with respect to oxygen and only allow oxygen and hydroxyl ions to migrate via vacancies. Therefore, the vehicle mechanism may be mainly applicable to compounds with loosely bonded, small molecules, especially acidic hydrates, such as NafionÒ , HCl, and Sb2O3nH2O. For the Grotthuss mechanism, proton conduction occurs by proton hopping between adjacent oxygen ions, which show pronounced dynamics on their own sites at normal lattice sites [97–99]. Fig. 5 schematically represents the Grotthuss proton conduction mechanism [100]. It can be seen that the

Table 1 Proton concentration of BaCe1xYxO3d, determined from room temperature permission from Ref. [94]. Copyright 2004 Elsevier.)

1

H NMR absolute intensities. (Reprinted with

Composition, x, in BaCe1xYxO3d

Proton concentration, mol%

0.01 0.03 0.05 0.10

2.0 3.3 6.3 7.4

Fig. 4. Measured dc conductivities for BaCe1xYxO3d (BCY) and SrCe0.95Y0.05O3d (SCY). (h) BCY x = 0:10; (d) BCY x = 0:05; (s) BCY x = 0:03; (4) BCY x = 0:01, and (j) SCY. (Reprinted with permission from Refs. [94,95]. Copyright 2000, 2004 Elsevier.)

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Fig. 5. Schematic representation of the Grotthuss mechanism for proton conduction in ceramic material. The solid dots is the proton; the open dots are oxygen. (Reprinted with permission from Ref. [100]. Copyright 2010 SCIELO.)

adsorption of the proton on the surface of the ceramic forms a weak O–H bond with one oxygen ion; this bonded proton then rotates and hops, forming a new weak O–H bond with the new neighboring oxygen ion. This process repeats again and again, until proton conduction is finished. Note that in this process, two rates are relevant: the rate of proton transfer and the rate of molecular reorientation. However, it is still not clear which is the rate-determining step in proton conduction. It has been suggested that proton transfer between neighboring oxygen ions is a rate-determining step for proton conduction, since the rotation of the proton around the oxygen is very fast with low activation energy [101]. Recently, quantum MD simulations suggested that the two rates of the Grotthuss conduction mechanism had similar probabilities of occurring. Although the rate-determining step in proton conduction is not clear, the rotation and hopping mechanism has been demonstrated by many experiments. For example, results obtained by FTIR showed that adsorption obviously occurred at the protonated oxide, indicating a bending vibration of the O–H bond. Furthermore, molecular dynamic simulation showed that the O–H bond consisted of a proton and its neighboring oxygen ion, and the frequency of the bending vibration of O–H was strongly dependent on the position of the oxygen ion in the crystal [102]. The same researchers also found proton diffusion between two closed oxygen ions after the O–H bond was broken, and diffusion of the O–H bond around the oxygen ion, which corresponded to the vehicle mechanism. To understand the proton transport mechanism, the influence of H/D isotope upon proton conductivity was also investigated [103]. In the vehicle mechanism, H replaced by D should have an insignificant effect on the proton conductivity because OH is the main conducting medium. But in the Grotthuss mechanism, the proton moves by hopping. The replacement of H by D will double the mass and increase the difficulty of hopping, leading to a reduction in the proton conductivity rate. For example, Stevenson et al. [103] studied the influence of H/D isotope on Yb- and Nd-doped BaCeO3. As shown in Fig. 6, the proton conductivity of BaCe0.85Gd0.15O3d was obviously reduced when H2O was replaced by D2O at 500 °C. They also observed the effect of temperature and found that the influence of the isotope upon conductivity was notable at low temperatures but unremarkable at high temperatures. At 900 °C, H/D isotope had no influence on proton conductivity, which signaled a change in the transport mechanism. The same phenomenon was also found by Iwahara et al. [104] in their research on SrZr1.95Y0.05O3d. Considering the weak influence of H/D isotope at high temperatures and the increase in oxygen ion transport number in these systems, the vehicle mechanism was believed to be dominant in proton conduction at high temperatures. With respect to electron conductivity, to achieve electroneutrality, the electrons produced by Reaction (VI) will transport together with the protons to the other side of the membrane. The electron conductivity inside the ceramic membrane can also be affected by the type, structure, and doping situation of the materials. The electron transport mechanism should be similar to what occurs inside a metal or semiconductor phase, although the transport rate in those cases may be faster or slower than that of a proton inside a ceramic membrane. When the electron transport rate is higher than the proton transport rate, the rate-determining step should be the proton conduction process; when the proton transport rate is higher, the rate-determining step should be the electron transfer process. Under high oxygen and low water partial pressure, these materials show mixed p-type electron and oxygen ion conductivity [105]. Conversely, electron conductivity is n-type in a reducing

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Fig. 6. Time-dependent conductivity variation in BaCeO3: 15% Gd at 500 °C due to alternate exposures to air saturated with water and heavy water vapor, indicating reversible H/D isotope effect. (Reprinted with permission from Ref. [103]. Copyright 1993 Elsevier.)

Fig. 7. Total conductivity of SrCe0.95Eu0.05O3d as a function of P H2 . (Reproduced by permission of The Electrochemical Society from Ref. [106]. Copyright 2003 ECS.)

atmosphere with hydrogen permeation [106]. As shown in Fig. 7, the total conductivity of SrCe0.95Eu0.05O3d was measured under a reducing atmosphere and found to be linearly dependent 1=4

on P H2 . Kosacki et al. [107] studied Yb–SrCeO3 with almost the same results. The H2 pressure dependence of the conductivity can be explained if we assume that electron conduction is n-type in a reducing atmosphere. In this model, a charge transfer reaction could be realized through the transfer of an electron from a dopant ion in a low oxidation state to a neighboring ion in a high oxidation state. 3. Development of dense proton–electron conducting ceramic materials and their associated membranes for hydrogen separation As discussed above, developing dense proton–electron conducting ceramic materials (DPECCMs) and their associated membranes is the key to obtaining high-performance H2 separation technology.

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To achieve both high-performance DPECCMs and their associated membranes, it is necessary to address synthesis, characterization, and performance validation in terms of improving and optimizing proton–electron conductivity, chemical stability, and mechanical stability, while maintaining low costs; each of these areas has been extensively reported on in the literature. This section will present an overview of these aspects. 3.1. Synthesis, characterization, and performance of single-phase ceramic mixed materials and their associated membranes Regarding single-phase ceramic mixed materials, Iwahara et al. [108] reported the proton conductivity of SrCeO3-based perovskite ceramics at high temperatures. In later research, Bonanos et al. [90] studied BaCeO3-based materials extensively because of their relatively higher proton conductivities [90]. For example, the hydrogen permeation performance of a self-short-circuited SrCe0.95Yb0.05O3d (SCYB)-based membrane had previously been reported by Hamakawa et al. [109]. Later, Guan et al. [110] reported that a hydrogen permeation flux of 2.7  108 mol cm2 s1 at 900 °C could be obtained using a self-short-circuited SrCe0.95Y0.05O3d (SCY)-based membrane, which was lower than the value estimated by proton conductivity measurements. Their impedance spectroscopy showed that the difference between the calculated and the observed permeation rates could be attributed to interfacial polarization. Therefore, to obtain the maximum hydrogen permeation rate through an SCY-based membrane, both ambipolar conductivity and interfacial properties need to be improved [110]. The hydrogen permeation performance of proton-conducting BaZr0.3Ce0.6Y0.1O3 (BZCY) and BaZr0.3Ce0.6Y0.1Zn0.05O3 (BZCYZn) perovskite membranes were investigated by Zhang et al. [111]. They found that H2 permeation fluxes through BZCY and BZCYZn membranes were slightly higher than those through baseline membranes, and adding water vapor to the sweep gas led to a slight increase in H2 permeation fluxes for both membranes, most likely because the water vapor facilitated hydrogen adsorption and/or an association process. In addition, hydrogen pumps using high-temperature proton conductors, including SrCe0.95Yb0.05O3 (SCYb), were examined as candidates for hydrogen separation from syngas [112,50]. All of these materials need to be self-short-circuited to provide electron conduction, since they have low electron conductivity and are not suitable for non-galvanic hydrogen separation directly. Thus, the hydrogen fluxes of SCYb and SCY single membranes are limited by their low electron conductivity. It is well known that proton materials with high proton conductivity do not always have sufficient electron conductivity, making them unsuitable for mixed ionic and electron conduction. H2 separation membrane materials must have high proton and electron conductivities to avoid having to apply external electrical power to the membrane. To address the issue of low electron conductivity, many researchers subsequently focused on increasing the electron conductivities of proton-conducting materials by developing mixed proton– electron conducting materials [113]. The first method developed to increase electron conductivity was doping the materials with aliovalent cations in a proton conductor [114]; the second method was mixing the materials with a metal or other electron-conducting material to enhance the electron conduction [115–117]. For example, initial attempts to improve the electron conductivity of SrCe1xMxO3 proton-conducting perovskites, where M is a rare earth metal, were done by replacing the usual univalent dopant with a multivalent dopant cation, such as Eu, Tm, or Sm [105,118]. As observed, introducing a multivalent element into a material can effectively increase its electron conductivity, with the degree of increase being dependent on the type and level of doped multivalent elements. In search of the enhancement mechanism of the doping, Lin et al. [119] tried some strontium cerate oxides and found that the electron hopping mechanism played the dominant role, and the charge transfer between two neighboring ions of different valences could result in enhanced electron conduction. In pure SrCeO3 or BaCeO3 materials, the transfer between Ce ions can be represented by:

Ce3þ þ h $ Ce4þ

ðVIIÞ

where h is the electron hole. In addition, oxygen vacancies and electron holes can be created when B-site Ce4+ is replaced by a lower valence element, such as Tm3+; some Tm3+ may be converted to Tm2+ to maintain charge neutrality. The charge transfer between thulium ions can be written as:

Z. Tao et al. / Progress in Materials Science 74 (2015) 1–50

Tm3þ þ h $ Tm4þ

15

ðVIIIÞ

Wachsman et al. [106] introduced a polaron model to understand electron transport in SrCe1xMxO3 materials. In this model, a charge transfer reaction was proposed—that is, an electron could be transferred from a dopant ion in a low-oxidation state to a neighboring ion in a high oxidation state. They found that the ionized dopant concentration was strongly dependent on the ionization potential of each dopant, under given thermodynamic conditions [120]. Based on the results of previous research on doped SrCeO3, it was found that the electron conductivity exhibited a relationship with the ionization potentials of the doping elements. In another of their studies [121], they also compared the electron conductivities obtained on several doped SrCeO3 oxides under pure oxygen, with the ionization potentials of their dopant ions. The total conductivity in pure oxygen was considered to be the electron conductivity. Their data table showed that except for Eu-doped SrCeO3, the electron conductivity increased with decreasing ionization potential. The third ionization potential of Tm was 23.7 eV, which was smaller than the fourth ionization potential of Tb (39.8 eV), Eu (24.8 eV), or Yb (25.0 eV). Therefore, Tm-doped SrCeO3 yielded the highest electron conductivity: 1.29  102 S cm1 at 800 °C. As has been recognized, the electron hopping distance is another important parameter in the hopping mechanism. Although the third ionization of Eu was lower than that of Yb, the electron conductivity of Eu-doped SrCeO3 was also lower than that of Yb-doped SrCeO3. The ionic radius of Eu (III) is 108.9 pm, which is larger than that of Yb (116 pm), probably suggesting that the hopping distance of Yb is smaller. However, the electron conduction behavior in a reducing atmosphere will be different than in an oxygen atmosphere. It was found that n-type electron conduction was the dominant process in the total electron conduction. The third ionization of Eu (24.8 eV) was higher than that of Sm (23.3 eV), suggesting that the reaction Eu2+ to Eu3+ might require more energy than Sm2+ to Sm3+, and Eu2+ might be more thermodynamically stable than Sm2+. Therefore, the n-type electron conduction in Eu-doped SrCeO3 was observed to be greater than that in Sm-doped SrCeO3, leading to Eu-doped SrCeO3 having a higher hydrogen flux than Sm-doped SrCeO3 under the same conditions, as shown in Fig. 8. The hydrogen permeability of SrCeTmO3 was tested by Lin et al. [119], who obtained a hydrogen flux of 3  108 mol cm2 s1 at 900 °C with a 1.6 mm-thick SrCeTmO3 membrane when 10% H2/He and air were used as the feed and sweep gases, respectively. The research also found that proton and electron conduction via the SrCeTmO3-based membrane occurred when it was exposed to water vapor or

Fig. 8. Hydrogen flux as a function of temperature for Eu- and Sm-doped SrCeO3d. Closed symbols = 100% H2; open symbols: P H2 = 0.972 atm, P H2 O = 0.028 atm. (Reprinted with permission from Ref. [47]. Copyright 2004 Elsevier.)

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hydrogen-containing gases, indicating the mixed proton–electron conducting characteristics of SrCeTmO3 membrane [119]. Wachsman et al. [106] also investigated the hydrogen permeation flux of Eu-doped SrCeO3 under different water pressures. Although there was an increase in proton concentration with increasing water pressure, the results indicated that hydrogen permeability decreased (as shown in Fig. 9), because an increase in oxygen pressure led to a decrease in electron concentration. As discussed above, the combined proton and electron conductivities (ambipolar conductivity) determine the hydrogen permeation flux. Oh et al. [121] investigated the effect of Eu dopant concentration in SrCeO3 on ambipolar conductivity and found that the highest ambipolar conductivity could be obtained over the compositional range from SrCe0.85Eu0.15O3d to SrCe0.8Eu0.2O3d, depending on the temperature. Matsumoto et al. [122] introduced a transition metal (e.g., Ru) as a partial substitute for Ce in Y-doped BaCeO3 to improve the electron conductivity of perovskite oxide. A significant increase in the hydrogen permeation flux was observed because of the increase in ambipolar diffusion associated with hole conduction. In addition, a high hydrogen flux of 0.026 mL min1 cm2 has been obtained with a BaCeNdO3 membrane using 15 vol.% hydrogen as the feed gas [123]. Wei et al. [124] investigated the mixed proton–electron conductivities of SrCe1xTbxO3d (x = 0.025, 0.05, and 0.1) under hydrogen-containing gases and found that an increase in downstream CO partial pressure could lead to a slight increase in the hydrogen flux. It is worth pointing out that more work is required to maximize electron conductivity by optimizing the doping level and element. 3.2. Synthesis, characterization, and performance of Ni and proton-conducting material composed materials and their associated membranes Although doping aliovalent cations into proton conductors can improve their electron conductivities, these conductivity levels are nonetheless limited when compared to high electron conductors. It is still necessary to add a second electron-conducting phase to form a composite with mixed electron conductions to function as a useful H2 separation membrane. The resulting proton-conducting ceramic material is called a cermet. The significant criteria in designing a cermet include chemical and thermal compatibility between the metal (and oxide) and ceramic phases, as well as possible thermal

Fig. 9. Hydrogen flux as a function of applied hydrogen chemical potential gradients under various conditions of P H2 O at 850 °C. (Reprinted with permission from Ref. [47]. Copyright 2004 Elsevier.)

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expansion matching between the various phases during fabrication of the membrane and its subsequent use. Typically, a metal (e.g., Pd or Ni) is added in the form of a powder or oxide [71]. A further possibility for improving the electron conductivity is to add a secondary ceramic phase (e.g., an n-type semiconductor) into the proton-conducting perovskite oxide. It is known that a (Ba,Sr)FeO3-based phase can yield high electron conductivity through forming an electron pathway. For example, the hydrogen permeation and electrical properties of composite oxide membranes composed of SrZrO3 and SrFeO3 were investigated by Unemoto et al. [125]. It was observed that hydrogen could permeate through the SrZrO3 and SrFeO3 composite oxides with platinum as a surface catalyst. However, the hydrogen fluxes were still insufficient compared to expectations. Wachsman and Jiang [126] used a chemical vapor deposition method to obtain a mixed proton–electron conducting material using perovskite oxide as the proton-conducting phase and Pd as the electron-conducting phase. However, Pd is an expensive metal and hence not cost-effective for use in separation systems. Thus, economical Ni has become the focus of development. In general, forming a dual-phase cermet via the addition of Ni can not only increase the electron flow in cerate- or zirconate-based oxides but also improve their mechanical stability [71]. More importantly, in the presence of Ni, the hydrogen permeability also increases, due to enhanced H2 ionization and adsorption at the membrane surface arising from the membrane’s improved endothermic hydrogen solubility [127]. Considering the interfacial overpotential of the H2 ionization process, the rate of non-galvanic hydrogen permeation (N) through a mixed-conducting membrane can be described by Eq. (6) [128]:

N ¼ ramb

Eg 2FL

ð6Þ

where L is the thickness of the membrane, F is the Faraday constant, g is the interfacial overpotential, E is the Nernst potential across the membrane as described in Eq. (5), and ramb is the membrane’s ambipolar conductivity. According to Eq. (6), if Ni is introduced into the membrane, the ambipolar conductivity will increase due to the increase in electron conductivity, leading to greater H2 permeation through the membrane. Aside from this, incorporation of a metal phase can also decrease the interfacial polarization resistance (or interfacial overpotential, g in Eq. (6)), resulting in a higher permeation rate. In experiments, after a high-temperature sintering process, Zhang et al. [129] prepared a Ni–BCY composite membrane that showed predominantly electron-conductive behavior. Using this membrane, they studied the interfacial and bulk resistances during H2 transport and found that temperature, membrane thickness, and hydrogen partial pressure were three major factors influencing the ratio of interfacial to bulk resistance. Song et al. [130] added Ni into SrCe0.8Y0.2O3 to form a dual-phase cermet for H2 separation. After comparing the hydrogen permeation fluxes with and without Ni, they found that adding Ni to the ceramic phase could increase its hydrogen permeability through increasing its electron conductivity; the obtained hydrogen permeation rate was 0.105 mL min1 cm2 for 0.25 mm-thick Ni/SCYb membrane with 20% H2/balance He as the feed gas. The results showed that Ni had good compatibility with SCYb during a high-temperature sintering process. Meng et al. [131] investigated Ni–BaCeTbO3 (BCTb) cermets for use in H2 separation membranes. Their results showed that these Ni–BCTb membranes were not sufficiently stable during hydrogen permeation, due to phase decomposition of the BCTb perovskite. Hence, more effort should be directed towards stability improvement of these membranes. 3.3. Synthesis, characterization, and performance of fluorite-structure materials and their associated membranes Normally, fluorite-structure oxides are good oxygen ion conductors, with a conductivity as high as 0.1 S cm1 [132]. This kind of membrane has been used extensively for oxygen separation and also for intermediate-temperature solid oxide fuel cells [133–136]. However, compared with the high oxygen ion conductivity of dopant fluorite-structure oxides, the proton conductivity of these oxides is generally low. For example, Nigara et al. [137–143] investigated the hydrogen permeation flux of doped ZrO2 and CeO2 under a humidified atmosphere and found the proton conductivity to be insufficient for H2 separation.

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Another material, (CeO2)0.8(YbO1.5)0.2, had a proton conductivity of only about 106 S cm1 at 1000 K [144], far lower than oxygen ion conductivity, but higher than the proton conductivity of single-crystal YSZ (108 S cm1) [142]. It seems that this class of pure fluorite-structure materials might not yield sufficient proton conductivity for H2 separation. To utilize such materials for H2 separation, some structural modifications may be necessary. Recently, Xie et al. [145] found that La1.95Ca0.05Ce2O6.975 fluorite could yield excellent proton conductivity, as well as electron conductivity of about 10–1.5 S cm1, as shown in Fig. 10. Sun et al. [146] found that a considerable hydrogen permeation flux could be obtained with Ce0.8Sm0.2O2d (SDC) when it was used as the hydrogen separation membrane. Recently a high concentration of elemental La was doped into CeO2, and the resulting ceramic membranes showed good hydrogen permeation fluxes [46,57]. All of these results again demonstrate the effectiveness of the doping strategy. Results reported by Iwahara et al. [147] indicate that good proton conductivity may require the proton-conducting ceramic material to have enough alkaline properties. As mentioned above, oxide ion vacancy plays an important role in the proton conduction process; oxide vacancies can be formed by doping La into the CeO2 crystal structure: CeO2

La2 O3 ! 2La0Ce þ 3Oox þ Vo

ðIXÞ

According to Reaction (IX), doping a high concentration of elemental La into the CeO2 structure will increase the oxygen vacancy, resulting in defect association between two oxygen vacancies and their closed La0Ce . This will decrease the mobility of oxygen ions in the structure and decrease the oxygen conductivity. However, the experimental results reported by Liu et al. [46,57] showed that proton conductivity increased with increasing lanthanum doping when the atomic doping level was less than 0.5 but decreased when the level was higher than that. This was probably due to the blocking effect of La2O3. As a result, they proposed a plausible mechanism for proton conduction in high-lanthanum-doped ceria on the basis of defect chemistry, as schematically shown in Figs. 11 and 12 [46,57]. Quantum mechanical calculations indicated that the energy of oxygen vacancy formation in the next nearest neighbor position (NNN) of La3+ was lower than in the nearest neighbor (NN) position of ceria [132]. Therefore, the oxygen vacancy generated by La doping could occupy the NNN

Fig. 10. Electron conductivity of La1.95Ca0.05Ce2O6.975 and La1.95Ca0.05Zr2O6.975 under different atmospheres. (Reprinted with permission from Ref. [145]. Copyright 2005 Elsevier.)

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Fig. 11. Plausible mechanism of proton conduction in La-doped ceria. (Reprinted with permission from Ref. [46]. Copyright 2010 Elsevier.)

Fig. 12. Schematic description of defect configuration in CeO2. (Reprinted with permission from Ref. [46]. Copyright 2010 Elsevier.)

position of the La ion, and the formed protonic defects after water was incorporated into the oxygen vacancy would also be located in this position. The electrostatic repulsion between the central cation (Ce4+ or La3+) and the proton would lead to the formation of O–H–O bonds, and the proton could hop between two adjacent oxygen ions at high temperatures. According to the Coulomb equation, the coulombic force increases with higher valence but decreases with increasing length between two electrical charges. Thus, the repulsion between La3+ and the proton should be lower than that between Ce4+ and the proton because Ce4+ has a higher valence and a smaller radius than La3+. Therefore, these protons would have a higher affinity for La ions, so they could move toward the oxygen ions in NN positions. At high temperatures and under the chemical potential of hydrogen, these protons would hop between two oxygen ions close to La3+, leading to a proton conduction process. In addition, oxygen ions with negative valence have a strong affinity for hydrogen and so would be the catalytic reaction centers after the protons jumped into the crystal structure, thereby enhancing hydrogen adsorption and facilitating surface diffusion. The oxygen vacancy concentration would be increased by increasing the La content in the ceria, enhancing water incorporation and proton conduction, and resulting in the high proton conductivity of La-doped ceria materials. In a further study, a trivalent element was co-doped into ceria to increase the proton conductivity of ceria-based materials. It was shown that doping a small amount of Sm into ceria increased the

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proton conductivity. The quantum mechanical calculations in Andersson et al. [148] also suggested that the oxygen vacancy caused by this Sm ion doping occupied the nearest neighbor (NN) position of the Sm ion. Based on that finding, the increased proton conductivity caused by doping a small amount of Sm into LDC could have a destructive effect on oxygen vacancy. Put simply, it is still not fully understood how increased proton conductivity is induced by trivalent cation co-doping in ceria, and more research is needed to achieve high-performance ceria materials for H2 separation. 3.4. Synthesis, characterization and performance of tungstate-based materials and their associated membranes Recently, lanthanide tungstate-based membranes have attracted increased attention because of their mixed proton and electron conductivity at high temperatures in hydrogen-containing atmospheres, as well as their high chemical stability [66,149,150]. For example, phase equilibriums in the La2O3WO3 system were studied in the 1970s [151,152]. Ivanova et al. [152] studied the two-phase structure of the La2O3WO3 system and pointed out that a La6WO12 phase existed from room temperature to 2150 °C. Yoshimura et al. [153] also indicated the existence of a La10W2O21 phase up to 1740 ± 30 °C, along with the formation of a La6WO12 phase above this temperature and up to 1960 °C. Some recent studies have also focused on phase formation. For example, Magraso´ et al. [154] prepared lanthanum tungstates with atomic La/W ratios between 4.8 and 6.0 using a freeze-drying method, and demonstrated that a single-phase material could only be obtained when the La/W ratios were between 5.3 and 5.7 at a sintering temperature of 1500 °C. Subsequently, many complex compounds have been studied, including La27(W1xNbx)5O55.55d and La28yW4+yO54+d [69,155]. These systems are considered to have a F43m space group structure, shown in Fig. 13. It can be seen that the lanthanum ions are coordinated with eight oxygen atoms: La (4a) forms relatively regular LaO8 cubes, while La (24 g) has a more distorted environment. W (4b) sites are fully occupied in octahedral coordination, and the extra W (24 g) is located on the lanthanum sites. The proton conductivity of La/W-based ceramic materials has also been studied. Ln6WO12 (Ln = Nd, Eu, or Er) was used as a H2 permeation membrane material due to its high proton and electron conductivities. For example, Escolástico et al. [66] studied Nd6WO12–Eu6WO12–Er6WO12 and found that these compounds had cubic fluorite symmetry. Their conductivities were measured in different atmospheres, Eu6WO12 yielding the highest conductivity. Hydrogen permeation measurements were performed on a 15 mm diameter Nd6WO12 disk with a thickness of 510 lm, sintered at 1550 °C; the

Fig. 13. Schematic structure of lanthanum tungstate material. (Reprinted with permission from Ref. [66]. Copyright 2013 Elsevier.)

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results are shown in Fig. 14. It can be seen that the hydrogen flux is about 0.023 mL min1 cm2 at 1000 °C with a mixture H2–He in the feed stream (2/8 mol/mol and 150 mL min1). Escolástico et al. [54,67,70] also investigated the effects of partially substituting Nd with La and the A sites deficiency using a La6WO12 material. Fig. 15 shows that (Nd5/6La1/6)5.5WO12d has higher hydrogen permeability at 1000 °C than Nd5.5WO12d because of the former’s higher total conductivity, caused by La having a higher ionic radius than Nd. At 1000 °C, the H2 flux for a (Nd5/6La1/6)5.5WO12d membrane (disk, 900 lm thick) made by sintering at 1550 °C reached 0.038 mL min1 cm2. It can be seen that the hydrogen flux is proportional to ln(p (H2, feed)/p (H2, perm)). In addition, the effect of H2O pressure on (La5/6Nd1/6)5.5WO12d was measured; these results are shown in Fig. 16. Notably, adding H2O promoted membrane hydration, and a water splitting reaction occurring at the permeate side. It can be seen that when water was introduced to both sides of a 900 lm-thick membrane, the hydrogen flux reached 0.044 mL min1 cm2 at 1000 °C. However, Lanthanum tungsten oxide showed a significant hydrogen permeation flux above 800 °C in a hydrogen potential gradient, induced by ambipolar proton–electron transport. Partial substitution of W with Mo was thought to be an efficient strategy to improve the n-type conductivity in LWO while it seemed not affecting the ionic conductivity, resulting in improving hydrogen permeation flux below 800 °C [156–160]. The hydrogen permeation flux of 6  104 mL min1 cm1 for Mo doped LWO was observed by Haugsrud et al. [160]. Serra et al. [161] investigated the hydrogen permeation performance and chemical stability of Mo doped NdWO2, and claimed that Nd5.5W0.5Mo0.5O11.25d could be promoting candidate for H2 separation at high temperatures. Furthermore, Hydrogen separation membranes based on dense ceramic composites in the La27W5O55.5–LaCrO3 system was investigated by Polfus and Serra et al. [162,163]. The advantages of the composited system were mainly observed at lower operation temperatures. Lanthanide tungstate-based membranes were found to have high chemical stability under hydrogen permeation conditions when 15% CO2 in Ar was used as the sweep gas [67]. Membranes treated under CO2 atmosphere for different lengths of time showed no degradation in either hydrogen flux or membrane crystal structure. Hydrogen flux as a function of time at 800 °C is shown in Fig. 17. The membrane’s high CO2 tolerance might have been induced by the system’s fluorite structure.

Fig. 14. Hydrogen permeation flux of a Nd6WO12 membrane as a function of temperature. Feed composition was moist H2 (20%) in helium, with argon as the sweep gas (both sides at atmospheric pressure). (Reprinted with permission from Ref. [66]. Copyright 2009 ACS.)

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Fig. 15. (a) Hydrogen flux of Nd5.5WO12 and (Nd5/6La1/6)5.5WO12 at 1000 °C as a function of hydrogen partial pressure at the feed side; (b) hydrogen flux of Nd5.5WO12 and (Nd5/6La1/6)5.5WO12 at 1000 °C as a function of the logarithm of the ratio of hydrogen partial pressure at the hydrogen feed side, p (H2, feed), to that at the permeating side, p (H2, perm). (Reprinted with permission from Ref. [54]. Copyright 2011 Elsevier.)

As discussed above, lanthanide tungstate-based membranes have high mixed electron–proton conductivity and stability in reducing, CO2-rich, and H2S-containing atmospheres. Therefore, Ln6WO12, (Nd5/6La1/6)5.5WO12d, and others in this category are promising membrane materials for industrial H2 separation.

3.5. Synthesis, characterization, and performance of BaPrO3-based materials and their associated membranes BaPrO3-based materials and membranes have been employed for H2 separation due to their proton conductivities. For example, Fang et al. [115] observed that the proton conductivity of BaPrO3-based materials increased when the electronegativity of A ions (such as Ba, Sr, and Ca) decreased, in the order Ba > Sr > Ca, while the reducibility of B ions (Pr, Ce, and Zr) increased in the order Pr > Ce > Zr. Considering Pr is more easily reduced than Ce and Zr, BaPrO3 might have high proton conductivity.

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Fig. 16. Hydrogen flux for (La5/6Nd1/6)5.5WO12d at different hydration configurations in a reactor. 50% H2 in the feed stream. (Reprinted with permission from Ref. [70]. Copyright 2012 Elsevier.)

Fig. 17. Hydrogen flux as a function of time, using 15% CO2–85% Ar as the sweep gas and 50% He–50% H2 as the feed gas at 800 °C. Both sides of the membrane were humidified. (Reprinted with permission from Ref. [67]. Copyright 2013 Elsevier.)

Fukui et al. [164] reported that the proton conductivity of Gd–BaPrO3 was about 0.1 S cm1 between 500 and 700 °C, while Bhide and Virkar [165] found that sintered Gd–BaPrO3 could be decomposed to Pr6O11 and Ba(OH)2 in boiling water. Magraso et al. [166] also found that the proton-conductive

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BaPrO3 system easily reacted with H2O and CO2, indicating that this kind of material might not be suitable for hydrogen separation. However, Fabbri et al. [167] recently found that BaZrPrYO3 could be used as an electrolyte for solid oxide fuel cells where proton conductivity was an important consideration. Whether or not BaZrO3 and BaPrO3 solid solutions can be used for hydrogen separation requires further study. 3.6. Synthesis, characterization, and performance of LaGaO3-based materials and their associated membranes Doped perovskite LaGaO3 materials have been explored as membrane materials for H2 separation and as solid electrolytes for solid oxide fuel cells [168]. Recently, Ma et al. [169–171] found that La0.9Sr0.1Ga0.8Mg0.2O3a (LSGM) yielded high proton conductivity in the presence of wet hydrogen: approximately 1.4  102–1.4  101 S cm1, with a proton transport number of 0.99 from 600 to 1000 °C—close to that of doped BaCeO3. This result indicated that LSGM was a pure proton conductor under wet hydrogen. However, LSGM has several drawbacks that limit its application: (1) It is not easy to obtain in a pure phase; consequently, it has low proton conductivity. (2) Its chemical stability at high temperatures is insufficient, as it easily decomposes into impurity phases, such as La2O3, LaSrGaO4, and La4SrO7. If the Sr doping level is high, Ga volatilization becomes more serious. It was also observed that La2O3 decomposition products appeared on the LSGM surface after extended operation at 800 °C, indicating that LSGM would only be suitable for applications below 800 °C. (3) It has low mechanical strength. (4) Gallium is expensive. (5) When fabricating LSGM, it is not easy to control its composition precisely. (6) Its low electron conductivity makes it difficult to use in hydrogen separation. (7) It is very difficult to find a suitable electron conductor to prepare the mixed proton–electron conducting materials. Based on these disadvantages, LSGM materials are of limited use in H2 separation, although some attempts have been made. 3.7. Synthesis, characterization, and performance of niobate/tantalite composite metal oxide-based materials and their associated membranes Shimara et al. [172] found that La3xSrxNbO7 yielded proton conductivity under wet hydrogen, with a maximum proton conductivity of about 103 S cm1 when x was 0.4 at 1000 °C. Traditionally, these materials could not be considered for use in hydrogen separation membranes because they did not satisfy the requirements for hydrogen separation [172]. Recently, however, two composites—consisting of proton-conducting, Ca-doped LaNbO4 and electron-conducting LaNb3O9—were explored for hydrogen separation, and yielded a hydrogen flux of 1.75  1010 mol cm2 s1 [173]. Obviously, though, this is still too low for H2 separation. The low hydrogen flux value was mainly due to the material’s low proton conductivity. A Ba3Ca1.18Nb1.82O9 complex-perovskite structure—representative of A2B0 1+xB00 MxO6d (where A = Ba or Sr, B0 = trivalent ion, and B00 = pentavalent ion) or A3B0 1+xB00 2xMxO9 (where A = Ba or Sr, B0 = divalent ion, and B00 = pentavalent ion)—was reported to have acceptable proton conductivity and excellent chemical stability in the presence of CO2 and H2O [173–176]. The off-stoichiometric perovskite Ba3CaNb2O9 exhibited a higher proton conductivity and better chemical stability than Nd-doped BaCeO3 [175]. Recently, Wang et al. [52] discovered that dopants in Nb sites could improve proton conductivity without sacrificing chemical stability. However, no reports describe using Ba3Ca1.18Nb1.82O9.8 (BCN18) as a hydrogen separation membrane material, probably because of its low sintering ability. To address this issue, Wang et al. [177] developed a two-step method to fabricate dense BCN18 material on NiO–BaZr0.1Ce0.7Y0.1Yb0.1O3.1 at a relatively low sintering temperature, as shown in Fig. 18. This work indicated that using BCN18-based membranes for H2 separation is feasible. One consideration, however, is that the electron conductivity of BCN18 is

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Fig. 18. SEM images of a dense BCN18 electrolyte on NiO–BZCYYb (Reprinted with permission from Ref. [177]. Copyright 2012 Elsevier.)

relatively low; another phase should be introduced into the BCN18-based membrane to enhance electron conductivity—for example, nickel. In addition, the chemical stability of BCN18 in atmospheres containing H2S needs to be improved.

3.8. Synthesis, characterization, and performance of ceramic asymmetric membranes As discussed previously, the overall hydrogen permeation process consists of several consecutive kinetic steps: gas/solid interfacial dissociation, solid-state diffusion, and solid/gas interfacial reassociation. The overall kinetics is often expressed as Lc = D/K, where K is the surface exchange coefficient and D is the tracer diffusion coefficient. In generally, when the membrane is thick, the rate-determining step is solid-state diffusion, and when it is thin, the rate-determining step is the surface dissociation/reassociation process. Fig. 19 shows the hydrogen flux through Ni–BZCY7 membranes as a function of membrane thickness, in the temperature range of 600–900 °C using wet 4% H2 (balance He) as the feed gas [41]. Fig. 19 shows that the hydrogen permeation flux of Ni–BZCY7 is inversely proportional to the membrane thickness, indicating that flux is limited by bulk diffusion in the thickness range of 0.25– 1.00 mm. Song et al. [178] investigated Ni–BaCe0.8Y0.2O3d to determine how hydrogen permeability depended on membrane thickness; they found that hydrogen permeation was dominated by ambipolar diffusion throughout the experimental temperature range of 550–900 °C, and that bulk diffusion was the rate-determining step down to a thickness of 80 lm. As shown in Fig. 20, the hydrogen flux through the Ni–BCY membranes was inversely proportional to the thickness, above 80 lm [178]. Another notable observation to be derived from this figure is that the linear extrapolation of the

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Fig. 19. Hydrogen flux through Ni–BZCY7 membranes as a function of thickness, from 600 to 900 °C using wet 4% H2 (balance He) as the feed gas. (Reprinted with permission from Ref. [39]. Copyright 2006 Elsevier.)

Fig. 20. Reciprocal thickness dependence of hydrogen permeation flux of Ni–BaCe0.8Y0.2O3d in 3.8% H2/balance N2 gas (P H2 O = 0.03 atm) feed mixture and 100 ppm H2/balance N2 gas sweep mixture. (Reprinted with permission from Ref. [178]. Copyright 2008 Elsevier.)

reciprocal thickness dependencies does not cross the origin, indicating that the surface exchange kinetics should not be neglected. Another important consideration for increasing hydrogen flux is to optimize the membrane design through carefully controlling the microstructure and the fabrication process. When bulk diffusion is

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the rate-determining step in proton transport, according to the equation above, reducing the thickness of the hydrogen permeation membrane may be effective in reducing resistivity. Some researchers [74] have therefore developed asymmetric membranes for hydrogen permeation, as shown in Fig. 21. From Fig. 21 it can be seen that this asymmetric membrane consists of a thin, dense membrane and a porous support layer. The support layer provides enough mechanical strength, and the thickness of the dense layer can be controlled by varying the amount of powder used in it. In general, the requirements for porous substrate materials are good chemical compatibility and thermal expansion with a thin film. Hamakawa et al. [74] successfully fabricated thin layers of either SrZr0.95Y0.05O3 or SrZr0.95Yb0.05O3 supported on a porous SrZr0.95Y0.05O3 substrate by spin-coating colloidal suspensions of the powders. The hydrogen permeation rates of asymmetric membranes with thicknesses of 2– 140 lm were found to be 500 times greater than of the rate for a 1000 lm-thick membrane with otherwise identical composition. For a 2 lm-thick membrane, the hydrogen flux reached 6  104 mol H2 cm2 min1. It was also observed that even with a 2 lm-thick SrZr0.95Yb0.05O3 membrane, the H2 transport rate was still controlled by bulk diffusion rather than the dissociative chemisorption of hydrogen. Cheng et al. [179] introduced a cost-effective dry-pressing method to fabricate supported, dense, thin membranes for hydrogen separation. In their fabrication process, SrCe0.95Tm0.05O3 was first separated according to particle size; then, the powder with a small grain size was used for the top thin layer, and the powder with a larger grain size was used for the substrate. Using this asymmetric membrane, with a thickness of 150 lm, a hydrogen flux of 9.37  108 mol cm2 s1 was obtained at 900 °C when a mixture of 10% H2–He was applied at the upstream side and air was used at the downstream side. The results are shown in Fig. 22.

Fig. 21. Scanning electron micrographs of several asymmetric SrCe0.95Yb0.05O3d membranes. Top layer, SrCe0.95Yb0.05O3d; substrate, SrZr0.95Y0.05O3d. (Reprinted with permission from Ref. [74]. Copyright 2002 Elsevier.)

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Fig. 22. H2 permeation flux as a function of temperature for different membrane thicknesses. The membrane contained a dense thin layer and a porous support layer; the materials for both layers were SrCe0.95Tm0.05O3. (Reprinted with permission from Ref. [179]. Copyright 2005 Elsevier.)

However, in this work, hydrogen permeation through the SrCe0.95Tm0.05O3 membrane was also limited by bulk diffusion. Note that a key factor in determining the shrinkage and porosity of the layers in a membrane is particle size. However, it is not easy to control shrinkage when fabricating thinner membranes based on particle size selection. It is necessary to develop innovative synthesis methods to fabricate asymmetrical membranes. To address this problem, NiO was selected as a second material for fabricating asymmetrical membranes [180]. In the fabrication process, NiO was used as a pore former because of volume reduction (NiO could be reduced to Ni by H2 when the membrane was operated in a hydrogen atmosphere at high temperatures). Another benefit of using a Ni layer is that it has catalytic activity towards H2 dissociation. Both NiO and Ni have high melting points (>1950 °C, >1453 °C), so they are stable at high temperatures during the sintering process. In addition, soluble starch is normally used as an extra pore former to control shrinkage of the substrate so that it matches the top layer in size. Hence, asymmetrical membranes with different thicknesses can be prepared using this method. Zhan et al. [180] fabricated dense, crack-free SrCe0.95Y0.05O3d (SCY) membranes with a thickness of 50 lm Shrinkage of the SCY/NiO/SS substrate was controlled by adjusting the amount of soluble starch and NiO in the substrate, and the thickness of the dense layer was controlled by adjusting the weight of SCY used in the top layer. The results showed that the hydrogen flux could be increased by decreasing the membrane thickness, yielding a significant improvement in surface exchange kinetics. As a result, a H2 flux as high as 7.6  108 mol cm2 s1 was achieved at 950 °C with a 50 lm membrane when 80% H2/He was used as the feed gas. Yoon et al. [181] prepared SrCeEuO3-based thin membranes for hydrogen separation. These dense layers were supported on NiO–SrCeO3 tubes by tape-casting and rolling methods. A hydrogen flux of 2.2 mL min1 cm2 was achieved at 900 °C using a feed gas containing 25% H2, 3% H2O, and 75% Ar. Wachsman et al. [56,182] prepared a SrZr0.2Ce0.8xEuxO3 thin layer supported on a NiO– SrZr0.2Ce0.8O3 porous structure. The thickness of the dense layer was varied from 17 to 50 lm, as shown in Fig. 23(b). Using these asymmetric membranes, a high hydrogen flux of 0.35 mL min1 cm2 was achieved at 900 °C when 100% hydrogen was used as the feed gas. In addition, a single-phase thin La0.5Ce0.5O2 membrane, supported on a substrate, with both electron and proton conductivity was developed as a hydrogen separation membrane by Zhu et al. [57]. An acceptable H2 flux of 2.6  108 mol cm2 s1 was achieved at 900 °C using 20% H2/N2 as the feed

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Fig. 23. SEM images of (a) SrZr0.2Ce0.7Eu0.1O3d and (b) SrZr0.2Ce0.65Eu0.15O3d asymmetric membranes. (Reprinted with permission from Ref. [54]. Copyright 2009 Elsevier.)

Fig. 24. Backscattered electron images of a cermet prepared by different methods. (Reprinted with permission from Ref. [44]. Copyright 2010 Elsevier.)

gas and dry, high-purity argon as the sweep gas. Zhu et al. [45] also fabricated a Ni– BaCe0.7Zr0.1Y0.2O3d (BZCY) dual-phase asymmetric membrane. Yan et al. [44] obtained uniform and more sintering active NiO–BZCY powders using a co-synthesis method to fabricate Ni–BZCY asymmetric membranes, as shown in Fig. 24. The hydrogen flux of a Ni–BZCY membrane obtained using this co-synthesis method (Ni–BZCY-1) was about 4.20  108 mol cm2 s1 at 900 °C, which was about 18% higher than that of a Ni–BZCY membrane fabricated using a traditional mechanical mixing method (Ni–BZCY-2). The difference in performance was explained using the concept of three-phase boundaries (TPB).

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Fig. 25. Flowchart of preparation method for an asymmetric membrane. (Reprinted with permission from Ref. [45]. Copyright 2011 Elsevier.)

Fig. 24 clearly shows that the homogeneity of the two phases in Ni–BZCY-2 was better than in Ni– BZCY-1, leading to a higher hydrogen exchange rate. Another explanation is that the homogenous mixture of Ni and BZCY was geometrically better for proton transport, as it provided more continuous routes. The activation energy (Ea) of H2 transport through these membranes was determined to be about 39.3 kJ mol1 for Ni–BZCY-1 and about 32.4 kJ mol1 for Ni–BZCY-2. The difference in activation energy was probably due to differences in surface exchange rate and homogeneity. Based on their experiments, Zhu et al. [45] fabricated a Ni–BZCY thin layer supported on a substrate using more uniform and sintering active NiO–BZCY powders. The fabrication procedure is presented schematically in Fig. 25. As shown in Fig. 25, the mixture containing NiO–BaCe0.7Zr0.1Y0.2O3 powder and 25 wt.% soluble starch was ball-milled with stabilized zirconia medium in ethanol for 24 h, followed by drying at 100 °C for 5 h. The resultant mixture was then pre-pressed uniaxially at 100 MP into pellets with a diameter of 20 mm. After that, a certain amount of NiO–BZCY was uniformly distributed on the substrate and then co-pressed under 400 MP. The obtained green disk was calcined in ambient air at 600 °C for 3 h to remove the soluble starch and form a porous substrate. Subsequently, the pre-sintered NiO–BZCY bi-layers were sintered at 1400 °C for 10 h in 5% H2/Ar to reduce the NiO to Ni metal. Finally, an asymmetric Ni–BZCY metal ceramic membrane with a porous substrate and a dense top membrane was obtained, as shown in Fig. 26. The results obtained using the asymmetric membrane shown in Fig. 26 indicated that it had a high hydrogen permeation flux when using 80% H2/N2 (with 3% H2O) as the feed gas and dry, high-purity argon as the sweep gas. A maximum flux of 2.4  107 mol cm2 s1 was achieved at 900 °C. Recently, Liu et al. [183] fabricated dense Ni–BZCYYb membranes (44 lm thick) supported on a porous Ni– BZCYYb substrate. A hydrogen permeation flux of 1.12 mL min1 cm2 was obtained at 900 °C when pure hydrogen was used as the feed gas and N2 as the sweep gas. The hydrogen permeation rates were much higher than those reported in the literature, indicating that this Ni–BZCYYb membrane has the potential to be used in practical applications.

3.9. Synthesis, characterization, and performance of ceramic hollow fibers Compared to other configurations, such as tubular or disk-shaped membranes, hollow fibers have several advantages, including facile high-temperature sealing, the formation of thin membranes due to the fibers’ asymmetric structure, and larger membrane area per unit of packing volume [183,184]. Therefore, a high hydrogen separation rate can be achieved using a hollow fiber membrane.

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Fig. 26. Cross-sectional SEM image of an asymmetric membrane. (Reprinted with permission from Ref. [45]. Copyright 2011 Elsevier.)

Fig. 27. SEM images of Ni–BZCY hollow fiber membrane. (Reprinted with permission from Ref. [185]. Copyright 2012 Chinese Society of Physics.)

Moreover, cracks or peeling between the two structural layers can be avoided, since the dense separation film and the porous support are made in a single step from the same ceramic material. These appreciable advantages have made these proton-conducting membranes promising candidates for practical applications. For example, Tan et al. [184] fabricated a gas-tight hollow fiber membrane from BaCe0.95Tb0.05O3d (BCTb) perovskite material by a combined phase inversion and sintering technique. Yang et al. [185] also prepared Ni–BZCY hollow fibers with metal–proton dual phases using modified oxides, as shown in Fig. 27. It can be seen that the hollow fiber membrane has a ‘‘sandwich’’ composition: finger-like structures form near the inner and outer walls, while a dense layer is present at the center. The hydrogen permeation flux through the Ni–BZCY hollow fiber membrane at 900 °C was found to be 0.53 lmol cm2 s1 when using 200 mL min1 wet 20%H2/80%N2 as the feed gas and 150 mL min1 dry, high-purity argon as the sweep gas [185]. However, neither of the hydrogen separation fluxes of these hollow fibers was as high as expected. SrCe0.95Yb0.05O3d (SCYb) hollow fiber membranes were prepared by Liu et al. [186] using a combined immersion-induced phase-inversion method. But the hydrogen permeation fluxes of these SCYb hollow fiber membranes were still lower than those obtained from a composite SCYb membrane. Song et al. [187] investigated the hydrogen permeation performance of cobalt-doped BaCeTbO3 hollow fibers. The hydrogen flux reached 0.385 mL cm2 s1 at 1000 °C when the flow rates of the 50% H2–He feed gas mixture and the nitrogen sweep gas were 60 mL min1 and 100 mL min1, respectively. It was also found that cobalt doping damaged the mechanical stability of the BaCeTbO3 perovskite membrane in a hydrogen atmosphere.

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3.10. Modifying the surface of a hydrogen separation membrane To improve the H2 transport rate across a ceramic membrane, one can reduce the membrane’s thickness, but to maximize the rate it may also be necessary to reduce the interfacial resistance at the gas/membrane interface. This can be done by modifying the surface structure of the membrane materials or by using coatings of catalytic materials to speed up the H2 dissociation/reassociation reaction rates. It has been reported [188] that hydrogen permeation rates with a Pt-modified membrane were over an order of magnitude higher than those on un-modified SrCeO3-based membranes under similar experimental conditions (10% hydrogen on the feed side, inert sweep gas on the other side). The higher hydrogen permeation flux of the membrane with a Pt coating was attributable to the catalytic properties of Pt on the feed-side surface in promoting the dissociation of H2 molecules and facilitating the subsequent protonation of the membrane. Li et al. [189] found that the H2 flux through a BaCeMn0.1O3 membrane was appreciably higher when porous platinum was coated on the membrane surface. Similarly, Zhang et al. [72] reported that a significantly improved hydrogen permeation flux could be obtained with the application of a thin Pd film on the membrane surface. 3.11. Stability of hydrogen separation materials At the current state of technology, hydrogen is mainly produced through the steam reforming of natural gas. The separation of hydrogen from gases containing CO2 is a crucial step in the mass production of pure hydrogen using a hydrogen separation membrane. However, gases containing a mixture of CO2 and moisture as well as traces of H2S, NOX, and SOX can produce an acidic environment. Many membrane materials, including Ce-based perovskite oxides such as (Ba,Sr)CeO3, become unstable in such an acidic environment, which diminishes the value of their proton conductivity and good sinterability. Zr-based ceramic materials such as zirconates ((Ba,Sr)ZrO3) have high chemical stability and good mechanical strength but insufficient proton conductivity. To improve the chemical stability of ceramic materials, it is possible to replace the desired fraction of Ce in (Ba,Sr)CeO3 with Zr; the resulting materials may have both high proton conductivity and good chemical stability in acidic gas mixtures. As discussed previously, a membrane composed of SrCe0.95Tm0.05O3 perovskite mixed conductor offered high hydrogen selectivity and sufficient mixed proton–electron conductivity at temperatures above 600 °C [179]. However, if SrCeO3-based membranes are to be used in industrial applications, their chemical stability needs to be improved. Kniep et al. [190] investigated how zirconium doping affected the chemical stability, lattice structure, proton and electron conductivity, and hydrogen permeation properties of SrCe0.95xZrxTm0.05O3-based membranes. Their results showed that doping zirconium into SrCe0.95Tm0.05O3 decreased the materials’ proton and electron conductivity and significantly decreased the hydrogen permeability of the membrane in CO2-free streams. Notably, SrCe0.75Zr0.2Tm0.05O3 membranes had a larger steady-state hydrogen flux and better chemical stability than a SrCe0.95Tm0.05O3 membrane in a CO2-containing environment [190]. Zuo et al. [38] investigated, under various conditions, the effect of Zr doping on the hydrogen permeation and chemical stability of cermet membranes composed of barium cerate and Ni. Their study showed that B-site substitution of Ce with Zr in BaCeYO3 dramatically enhanced the chemical stability of Ni–BZCY membranes in a CO2and H2O-containing atmosphere, even though the hydrogen permeation rate decreased slightly at higher temperatures, as shown in Fig. 28. Among all the Ni–Ba(Zr0.8xCexY0.2)O3d compositions studied, Ni–BZCY7 (x = 0.7) showed both high hydrogen permeation flux and adequate stability; a Ni–BZCY7 hydrogen separation membrane was therefore thought to have potential for practical application. Zuo et al. [39] also systematically investigated the hydrogen transport properties and chemical stability of Ni–BZCY7 cermet membranes under various water pressure conditions. They found that hydrogen flux could be increased by adding moisture to the feed gas or increasing the hydrogen partial pressure gradient. A new Ni–BaCe0.4Zr0.4Nd0.2O3d (Ni–BZCN4) cermet membrane has been explored to separate hydrogen from a mixed gas containing hydrogen [191]. The permeation stability of Ni–BZCN4 membrane was investigated under atmospheres containing H2O or 30% CO2. The membrane maintained a steady permeation flux after 100 h of operation, as shown in Fig. 29, suggesting that it should be

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Fig. 28. Time dependence of hydrogen flux through 0.75 mm-thick Ni–BZCY (0.4 < x < 0.8) membranes in wet 20% CO2 (balance 40% H2/He) feed gas at 900 °C. (Reprinted with permission from Ref. [38]. Copyright 2006 ACS.)

Fig. 29. Time dependence of the hydrogen permeation flux of a 0.75 mm-thick Ni–BZCYYb membrane and absolute humidity in feed gas passing through the reactor. The feed gas consisted of 20 mL min1 H2, x mL min1 CO2, and 80–x mL min1 He. (Reprinted with permission from Ref. [192]. Copyright 2014 ACS.)

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suitable for separating hydrogen from a mixed gas containing H2O and CO2. Recently, unprecedented CO2-promoted hydrogen permeation in a Ni–BaZrCeYYbO3 membrane was investigated by Fang et al. [192]. In their research, hydrogen flux was promoted rather than diminished by CO2 (diminishment having been observed for other high-temperature proton-conducting membranes). The hydrogen flux enhancement in this membrane was attributed to the increase in the moisture content of the feed gas after CO2 was introduced. The rise in water partial pressure in the feed gas upon the introduction of CO2 was due to the reverse water–gas shift (RWGS) reaction:

H2 ðgÞ þ CO2 ðgÞ $ H2 OðgÞ þ COðgÞ

ðXÞ

This reaction consumes H2 and CO2 but also generates equal amounts of H2O and CO, leading to an increase in proton conductivity and, thus, a hydrogen permeation flux. Fang et al. [193] also investigated the influence of high content of H2O and CO2 on hydrogen flux of Ni–BaZrCeYYbO3, and demonstrated that the high content of H2O could promote the reaction between CO2 and BZCeYYb, forming insulating secondary phases and causing performance degradation. Although there are extensive researches in improving the chemical stability of BaCeO3 based materials, the improvement is limited. BaZrO3 was thought to be the proton conductor with excellent stability in CO2 and H2O containing atmosphere. Recently, BaZrO3 based electrolyte solid oxide fuel cells have been extensively investigated [194–218]. Three different methods including EDTA–citric method (CEC) method, solid state reactive sintering (SSRS) method and solid state reaction (SSR) method were employed to fabricate dense Ni–BZY membranes with large BZY grains [193]. Except for SSR method, BaY2NiO5 phase was also observed using other two method, and the dense Ni–BZY membranes containing BaY2NiO5 was found to be instable because that BaY2NiO5 could react with H2, H2O and CO2 to form insulating BaY2O4, Ba(OH)2 and BaCO3 phases, resulting in the degradation of hydrogen flux [219,220]. Zhu et al. [221] systematically evaluated the chemical stability and hydrogen permeation flux using a novel Ni–BPZY dense membrane, and found that this membrane exhibited satisfactory tolerance to CO2 and H2O under their operation conditions. It is well known that the reforming and partial oxidation of natural gas usually results in H2S. Although Ni–BaZr0.1Ce0.7Y0.2O3 cermet shows adequate performance and rather good stability in H2O and CO2, its stability under a reforming gas containing H2S is also an important concern. Under industry-standard hydrogen separation conditions (e.g., high temperature), H2S may poison Ni-based cermets by sulfur adsorption on the nickel surface or by the formation of nickel sulfides. Fang et al. [51] investigated H2S poisoning and regeneration of Ni–BaZr0.1Ce0.7Y0.2O3d at intermediate temperatures (Fig. 30). The hydrogen permeation flux decreased by about 30% and then stabilized for 40 h after the introduction of H2S. When H2S was removed, the hydrogen permeation flux recovered by approximately 13% over several hours and then remained stable. Thermodynamic calculations indicated that at 973 K, the reaction between BZCY and H2S was very weak, and Ni3S2 was unstable in 60 ppm H2S. These results suggested that sulfur adsorbed on the Ni surface was the poisoning mechanism at this temperature. However, at 1073 K, the reaction between BZCY and H2S was the main cause of the sulfur poisoning of Ni–BZCY. Fang et al. [51] also suggested that the hydrogen permeation flux of Ni–BZCY could be stabilized in 30 ppm H2S, and that H2S tolerance might be improved by increasing the water partial pressure in the feed gas. Liu et al. [40,222] tested the chemical stability and hydrogen permeation performance of Ni–LDC under an atmosphere containing CO2 and H2S. Fang et al. [40] measured the hydrogen performance regeneration behavior of Ni–La1.95Ca0.05CeO7, as shown in Fig. 31. Clearly, hydrogen performance decreases by about 10% after the introduction of CO2, then completely recovers within one hour of CO2 being removed. After (1) analysis of the feed gas composition and (2) thermodynamic calculations, it was suggested that the performance loss in the presence of CO2 was caused by the reverse water-shift reaction. However, as mentioned above, the enhanced hydrogen flux for the Ni–BZCYYb membrane was attributed to the increase in the moisture content of the feed gas after CO2 was introduced. According to the RWGS reaction (Reaction (X)), adding more CO2 generates an equal amount of H2O and CO. Ni–LDC is not stable in atmospheres containing H2S, as Fig. 32 demonstrates. Obviously, the hydrogen flux decreased after H2S was introduced into the feed gas. Thermodynamic calculations indicated

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Fig. 30. Sulfur poisoning and regeneration behavior of Ni–BZCY in a feed gas containing 60 ppm H2S balanced with 40% H2 and 1.5% H2O at 973 K. (Reprinted with permission from Ref. [51]. Copyright 2009 Elsevier.)

Fig. 31. Hydrogen permeation flux of Ni–CLC125 after the introduction and removal of 40% CO2 in the feed gas at 900 °C. (Reprinted with permission from Ref. [40]. Copyright 2010 Elsevier.)

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Fig. 32. Hydrogen permeation flux of Ni–CLC125 after the introduction of different H2S concentrations. (Reprinted with permission from Ref. [222]. Copyright 2009 Tans Tech Publications Inc.)

that the reaction between LDC and H2S was the main cause of the decreased hydrogen performance, and also showed that the reaction between H2S and LDC was not feasible until the concentration of H2S was lower than 50 ppm. Therefore, if one wants to use Ni–LDC as a hydrogen separation membrane under an atmosphere containing H2S, the H2S concentration must be below 50 ppm. 4. Ceramic membrane-based H2 separation systems design and fabrication To date, hydrogen separation membranes made from high-temperature ceramic proton conductors can only be utilized at the laboratory level. Hydrogen permeation measurements have been performed using a testing system, as schematically shown in Fig. 33. As for most ceramic membranes, the most critical issue is sealing, which is done using a glass ring under high-temperature operation. In general, although a glass ring yields a seal efficiency higher than 95%, the properties of glass make it difficult to achieve 100% sealing efficiency. Another disadvantage is that a glass ring cannot meet the requirements of operation cycles. Also, under high sealing temperatures, the reaction between the glass ring and the ceramic membrane and the tolerance of strict practical atmosphere including reducing and acid gas should be considered. Research scientists at Argonne National Lab employed a gold ring as the sealant [127], improving the sealing efficiency under high operation temperatures. However, using gold rings increases the system cost. Ultra-torr fittings have also been used to seal an end-capped ceramic tubular cell. Compared to a traditional disk membrane, a tubular membrane has the advantages of high surface area and excellent seal efficiency. Recently, Wachsman et al. [181,226] developed dense ceramic, hydrogen-permeable thin films on the inner side of tubular supports. Fig. 34 shows the basic design of these thin-film hydrogen membranes. It can be seen that a thin ceramic film is coated on the inner side of a catalytic, tubular-type, porous support. On the outer side of the tube, a methane and steam mixture reacts to form CO, CO2, and H2. Finally, hydrogen permeates the thin film due to the hydrogen partial pressure

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Fig. 33. Schematic diagram of a H2 permeation test system. (Reprinted with permission from Ref. [183]. Copyright 2011 Elsevier.)

Fig. 34. Basic design of a tubular-type hydrogen membrane cell for the methane steam-reforming process. (Reprinted with permission from Ref. [181]. Copyright 2009 Wiley.)

gradient between the outer and inner sides of the membrane. Fig. 35 presents the configuration of a hydrogen permeation reactor that uses 15 cm-long, one-end-closed, tubular-type hydrogen membrane cells. Regarding sealing, the difference between this tube cell and disk membrane is that ultra-torr fitting is used to avoid using a glass sealant. More importantly, this sealing process can be completed at room temperature, avoiding possible unexpected reactions between the glass sealant

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Fig. 35. Configuration of a hydrogen permeation reactor using a 15 cm-long, one-end-closed, tubular-type hydrogen membrane cell. (Reprinted with permission from Ref. [181]. Copyright 2009 Wiley.)

and the membrane at high temperatures. This may be a promising method for sealing hydrogen separation membranes.

5. Challenges of H2 separation using dense proton–electron conducting ceramic materials/ membranes Although great progress has been made in recent years regarding H2 separation using dense proton–electron conducting ceramic materials/membranes, in terms of both technology and commercialization, several challenges remain: (1) Low proton and electron conductivities of these ceramic materials/membranes for high H2 flux generation. (2) Low H2 dissociation/reassociation reaction rates at the membrane/gas interfaces. (3) Low chemical stability of the materials in an acidic gas atmosphere at high temperature. (4) Low mechanical strength and durability for long-term operation. (5) Low thermal stability for long-term operation. (6) Insufficient optimization of design and fabrication for the materials and membranes and their associated systems. (7) Insufficient fundamental understanding of material/membrane performance and degradation. (8) Relatively high cost with respect to commercialization. Despite the several advantages that dense proton–electron conducting ceramic materials/membranes have over other types of materials/membranes, they are limited by the challenges listed above. Practical application of these materials/membranes in H2 separation thus remains unrealized, particularly for industrial-scale implementation. Hence, developing highly active, selective, and stable

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ceramic materials and their associated membranes for H2 separation is still the major focus in this area. Below, we summarize the various challenges. (1) Low proton and electron conductivities for high H2 flux generation. Proton and electron conductivities are probably the most important properties characterizing the H2 separation performance of a ceramic membrane. As discussed above, some ceramic materials have high proton conductivity but low electron conductivity, while some have the converse. It is still a challenge to fabricate materials with simultaneously high proton and electron conductivities. A ceramic material normally contains two phases, proton-conducting and electron-conducting. Optimizing the phase composition ratio and the distribution of the phases inside the membrane is also not easy. Although tremendous work has been put into improving proton and electron conductivities through optimizing the materials’/membranes’ properties and compositions, as well as reducing membrane thickness by developing supported and self-supporting multilayered structures, at the current state of technology these conductivities remain insufficient for commercial use of these materials in H2 separation. (2) Low H2 dissociation/reassociation reaction rates at the membrane/gas interfaces. As discussed above, the H2 dissociation/reassociation reactions are two necessary steps in the entire H2 transport process. They are mainly controlled by the surface activities on the ceramic membrane. Although several approaches have been attempted—such as surface modification and depositing a thin layer of metal (Pt, Ni, Ru, etc.) onto the membrane surface—these surfaces still have insufficient activity and show high resistance to H2 dissociation/reassociation reactions. Furthermore, the modified/catalyzed surfaces are sometimes easily poisoned by trace gas components such as H2S, CO, NOX, SOX, and organic sulfur compounds, leading to gradual performance degradation due to high surface resistance. (3) Low chemical stability in acidic gas atmospheres at high temperatures. As previously mentioned, all ceramic materials and membranes for H2 separation must be usable in acidic gas mixtures that contain acidic and reducing gas components, such as H2S, CO, NOX, SOX, and organic sulfur compounds. Tolerance of an acidic/reducing environment is very important, so materials and their corresponding membranes must have high chemical stability for practical applications. Although significant progress has been made in improving the chemical stability of these materials and membranes, by exploring new materials and optimizing their composition and structure, they still remain insufficiently stable in acidic and reducing environments over a practical operational lifetime. (4) Low mechanical strength and durability for long-term operation. Mechanical strength and durability are essential for ceramic materials and their associated membranes to be used in H2 separation. Tremendous progress has been made in (i) creating strong, durable porous metals or ceramic materials as supports, (ii) incorporating these materials into membranes, (iii) combining proton–electron phases with refractory oxides to form composite materials/membranes, and (iv) eliminating the delaminating issue that arises with multilayered membranes. However, further efforts are still needed to achieve the commercial use of these ceramic materials and membranes. (5) Low thermal stability for long-term operation. As previously mentioned, ceramic materials and their dense membranes must work at temperatures above 600 °C to generate enough proton conductivity. Although some efforts have been made to eliminate spontaneous combustion and decomposition phases, and phase/multilayer delamination at high temperatures, some of the materials and membranes explored for H2 separation are still not sufficiently stable from a lifetime perspective. (6) Insufficient optimization of design and fabrication for materials/membranes and their associated systems. Normally, optimizing the design and fabrication of materials/membranes and their associated operation systems is a challenge for practical applications. Optimal marching among material, component and fabrication process is strongly related to the system design and performance validation. Further work is definitely necessary in this area to achieve the commercialization of H2 separation technology using dense proton–electron conducting ceramic membranes.

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Table 2 Summary of the challenges and possible research directions for hydrogen separation membranes. Challenges

(1) Low proton and electron conductivities of these ceramic materials/membranes for high H2 flux generation (2) Low H2 dissociation/reassociation reaction rates at the membrane/gas interfaces (3) Low chemical stability of the materials in an acidic gas atmosphere at high temperature (4) Low mechanical strength and durability for long-term operation (5) Low thermal stability for long-term operation (6) Insufficient optimization of design and fabrication for the materials and membranes and their associated systems (7) Insufficient fundamental understanding of material/membrane performance and degradation (8) Relatively high cost with respect to commercialization

Research directions

(1) Enhancement of proton and electron conductivity as well as chemical, mechanical, and thermal stability by exploring innovative ceramic materials and their associated membranes (2) Further fundamental understanding through both experiments and theoretical modeling (3) Optimizing material/membrane structures, system designs, fabrication processes, and operating conditions for practical applications

(7) Insufficient fundamental understanding of material/membrane performance and degradation. Although the literature contains attempts to fundamentally understand the performance and degradation of ceramic materials/membranes during H2 separation, through both experimental and theoretical modeling approaches to improve proton and electron conductivity, the work in this area remains insufficient. (8) Relatively high cost with respect to commercialization. The high cost of materials, components, and systems is always a limiting factor in technological commercialization. In H2 separation technology using dense proton–electron conducting ceramic materials/membranes, some of the promising materials that demonstrate high performance contain rare and therefore expensive elements, which may limit the materials’ use in commercial systems. The advantages, challenges and possible research directions for hydrogen separation membrane were summarized in Table 2.

6. Summary and proposed research directions This paper has provided a comprehensive overview of several decades of development as well as recent trends in H2 separation technology using dense proton–electron conducting ceramic-based membranes, with the intention of facilitating the research, development, and commercialization of this technology. Various proton–electron conducting materials and their associated membranes that have been explored and reported in the literature were summarized and classified into several important categories, such as Ni composite proton-conducting materials, as well as tungstate-based, BaPrO3-based, LaGaO3-based, and niobate/tantalite composite metal oxide-based ceramic materials/membranes. Membrane designs were also discussed, including ceramic asymmetric membranes (supported and self-supported) and surface-modified membranes. We also reviewed important properties of ceramic materials and membranes, such as proton and electron conductivity, performance, H2 transport flux, and lifetime stability. To highlight the technical progress to date, Table 3 summarizes all possible ceramic materials and membranes, along with their associated properties and performance, so readers can quickly locate the information they are seeking. We also discussed industrial attempts and achievements during recent decades in the fields of designing and fabricating H2 separation methods for practical applications, to give a clear picture of the current state of technology. It can be seen that H2 separation technology using proton–electron conducting ceramic membranes to obtain industrially usable hydrogen fuel is not yet close to reaching the requirements for commercialization, due to a number of major technological challenges. To overcome these challenges, we propose several future research directions.

Table 3 Summary of dense proton–electron conducting ceramic materials/membranes and their associated properties and performance for H2 separation. Film thickness (lm)

Normalized flux (mL min1 cm2)

SrCe0.9Eu0.1O3d SCTb SrCe0.95Yb0.05O3d SrCe0.95Sm0.05O3d BaCe0.9xY0.1RuxO3a (x = 0.075, 0.1) Sr0.97Ce0.9Yb0.1O3d (SCYb) BaZr0.8Y0.15Mn0.05O3d SrCe0.75Zr0.20Tm0.05O3a BaCe0.95Nd0.05O3d BaCe0.9xY0.1RuxO3a (x = 0.01) SrCe0.75Zr0.20Tm0.05O3d SrCe0.95Tm0.05O3d SrCe0.95Y0.05O3d (Nd5/6La1/6)5.5WO12 La5.5WO11.25 Nd6WO12 (La5/6Nd1/6)5.5WO12d La5.5W0.8Re0.2O11.25d Nd5.5W1xMoxO11.25d La27Mo1.5W3.5O55.5 La27W5O55.5–LaCrO3 La0.87Sr0.13CrO3d SrCe0.9Eu0.1O3d/NiO–SrCeO3 tubular SrCe0.7Zr0.2Eu0.1O3d SrZr0.2Ce0.8xEuxO3d (x = 0.1 and 0.15) BaCe0.2Zr0.7Y0.1O3d/ Sr0.95Ti0.9Nb0.1O3d Ni–BaCe0.95Tb0.05O3d

30 1600 670 1720 500 1160 900 1200 1000 1000 1600 1600 110 900 900 510 900 700 1000 650

0.6 0.016 0.009 1.21  109 mol cm2 s1 6.5  108 mol cm2 s1 (x = 0.1) 33 lmol cm2 s1 0.03 0.042 0.017 6.47  108 mol cm2 s1 0.005 3  108 mol cm2 s1 0.023 0.03 0.136 0.023 0.046 0.02 0.3 6  104 1.1_10_3 108 mol cm2 s1 2.2 0.23 0.35 0.026 mmol cm2 s1

SrCe0.9Eu0.1O3d/Ni–SrCeO3 SrCe0.95Y0.1O3d SrCe0.95Tm0.1O3d asy SrCe0.95Yb0.1O3d Sr(Ce0.6Zr0.4)0.85Y0.15O3d Ni–BaZr0.1Ce0.7Y0.1Yb0.1O3d Ni–Ba(Zr0.7Pr0.1Y0.2)O3d Ni–BaZr0.1Ce0.7Y0.2O3d

550 1133 33 17 4000 90 30 50 150 2 500 44 440 30

0.914 0.6 7.6  108 mol cm2 s1 0.16 6  104 mol cm2 s1 0.184 mmol min1 cm2 1.12 5.44  1010 mol cm2 s1 2.4  107 mol cm2 s1

Testing temperature (°C) 900 900 900 850 850 804 1000 900 825 800 750 900 800 1000 1000 1000 1000 760 1000 900 700 600 900 900 900 800

Chemical stability

Ref.

– – – – – – Good Stable – – Good – – Stable Stable Stable Stable Stable Stable Stable

[121] [124] [223] [47] [49] [188] [224] [225] [123] [122] [190] [119] [110] [54] [67] [66] [70] [68] [161] [160] [160,161] [229] [181] [182] [56] [232]

in CO2

in in in in in in in

CO2, H2O; stability in H2S unknown CO2, H2O, H2S CO2 CO2, H2O, H2S CO2, H2O, H2S CO2, H2O, H2S CO2, H2O

– – Good Good Good

850

–

[131]

900 850 900 950 K 900 900 900

– – – –

[227] [180] [179] [74] [230] [183,192,193] [221]

900

Good Stable in CO2, H2O Stable in CO2, H2O, H2S

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Material

[45] (continued on next page) 41

42

Material Ni–La2xSmxCe2O7 Ni–CayLa0.5yCe0.5O2d Ni–La0.5Ce0.5O2d Ni–BaZr0.1Ce0.7Y0.2O3d Ni–BaCe0.8Y0.2O3d Ni–BaCe1xYxO3d Ni–BaCe0.9Y0.1O3d Ni–BaZr0.8Y0.2O3d Ag–BaCe0.5Zr0.3Y0.16Zn0.04O3 Hollow fiber BaCe0.85Tb0.05Co0.1O3d BaCe0.8Y0.2O3d BaCe0.95Tb0.05O3d Ni–BaZr0.1Ce0.7Y0.2O3d

Film thickness (lm) 600 600 48 266 80 190 400 –

Normalized flux (mL min1 cm2) 2.86  108 mol cm2 s1 1.9  108 mol cm2 s1 6.83  108 mol cm2 s1 0.805 0.25 0.06 0.76 2  108 mol cm2 s1 – 0.385 0.38 0.422 lmol cm2 s1 0.53 lmol cm2 s1

Testing temperature (°C) 900 900 900 900 900 900 800 900 – 1000 1050 1000 900

Chemical stability

Ref.

Good Good Good Good –

[46] [40] [231] [38,39,127] [178] [37] [227] [219,220] [233]

– – Poor Poor Poor Good

[187] [228] [184] [185]

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Table 3 (continued)

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(1) Enhancement of proton and electron conductivity as well as chemical, mechanical, and thermal stability by exploring innovative ceramic materials and their associated membranes. Almost all of the possible ceramic materials and their associated membranes that have shown some proton and electron conductivity for H2 transport, as well as limited chemical, mechanical, and thermal stability, have been explored, and some progress has been achieved in terms of these properties and material/membrane performance. However, the technological advances thus far achieved are still not sufficient for practical H2 separation applications. For example, proton and electron conductivities in excess of 0.1 S cm1 are required for industrial H2 separation. The current materials/membranes are not close to reaching this level for both proton and electron conductivity at the same time. Hence, to achieve breakthroughs in this area, a priority must be to develop new material synthesis technology that yields innovative materials/membranes with optimal performance. Two important types of ceramic materials should be emphasized here: (1) doped composite materials, which are synthesized by combining several different materials, and (2) multilayered dense thin membranes. Doped composite ceramic materials should have different properties and transport performance than their individual components because the individual substances in the doped composites can experience a synergistic effect, facilitating electron and proton conduction and protecting active materials from chemical and mechanical degradation. As a result, the obtained composites may have high proton–electron conductivity and high chemical/mechanical/thermal stability during the H2 separation process. In the case of multilayered dense thin membranes, both sides of the membrane can be coated with stable catalyst layers to speed up the H2 dissociation/reassociation reactions for the H2 transport process. These multilayered membranes can contain a robust, porous support layer, so the active membrane can be made very thin. Dense thin membranes can effectively reduce proton and electron conduction resistance, leading to high proton and electron conductivity. A robust support can improve the membrane’s mechanical and thermal stability. (2) Further fundamental understanding through both experiments and theoretical modeling. For the down-selection of ceramic materials and their associated membranes, including the design and optimization of new structures to improve conductivity and stability, better fundamental understanding through both experiments and theoretical modeling is necessary. For example, we need to fundamentally understand the mechanisms of proton and electron transport and their relationship to active site structures and composition, using both theoretical calculations (molecular- and electron-level modeling) and experimental approaches, to guide new materials development. Furthermore, to mitigate material/membrane degradation, it is necessary to understand the degradation mechanisms and failure modes using both experimental and theoretical modeling approaches. For example, a variety of instrumental analysis methods (e.g., SEM, TEM, XRD, XPS, NMR, HPLC, GC, and so on) and electrochemical methods (CV, RDE/RRDE, and EIS) can be used to characterize materials and catalysts before and after lifetime tests. With better understanding, it will be possible to develop new mitigation strategies. (3) Optimizing material/membrane structures, system designs, fabrication processes, and operating conditions for practical applications. Besides improving ceramic materials in terms of their properties and performance (e.g., conductivity and stability), for overall performance it is also important to improve the membrane assembly and its associated hardware, along with the system design, and to optimize operating conditions. For example, in recent years, some innovative membrane assembly and transport cell designs have appeared to offer the right approaches for performance improvement. Cost reduction should also be a consideration in material/membrane creation as well as in system design, fabrication, and operation. In summary, H2 separation from reforming and gasified gases to produce pure H2 fuel is an important research and development subject for the global energy economy and many industrial processes. We strongly believe that with continued, extensive efforts focused on developing innovative, composite, and multilayered dense proton–electron conducting ceramic materials/membranes to overcome the challenges of insufficient conductivity and stability, H2 separation technology using such materials and membranes will become practical in the near future.

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Acknowledgments This work is supported by the National Natural Science Foundation of China (Grant No. 21406190, 21173039), the National Key Technology R&D Program of China (Grant No. 2013BAC13B01), the Natural Science Foundation of Jiangsu Province (Grant No. BK2012248), Natural Science Foundation of the Higher Education Institutions of Jiangsu Province (Grant No. 13KJB430023), the Specialized Research Fund for the Doctoral Program of Higher Education of China (Grant No. 20110075110001) and the Innovation Program of the Shanghai Municipal Education Commission (Grant No. 14ZZ074). All of above financial supports are gratefully acknowledged. References [1] Marban G, Vales-solis T. Towards the hydrogen economy. Int J Hydrogen Energy 2007;32:1625–37. [2] Kennedy D. The hydrogen solution. Science 2004;305:917. [3] Conte M, Iacobazzi A, Ronchetti M, Vellone R. Hydrogen economy for a sustainable development: state-of-the-art and technological perspectives. J Power Sources 2001;100:171–87. 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