Metal–organic frameworks in fuel cell technologies

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Metal—organic frameworks in fuel cell technologies Yuqian Ren, Guo Hui Chia, Zhiqiang Gao ∗ Department of Chemistry, National University of Singapore, Singapore 117543, Singapore Received 13 May 2013 ; received in revised form 9 October 2013; accepted 28 November 2013

KEYWORDS Metal—organic frameworks; Fuel cell; Electrocatalysis; Hydrogen storage; Proton conductor

Summary The human appetite for energy is constantly growing and becoming increasingly difficult to satiate. Fossil fuels are quickly becoming unsatisfactory substrates due to the undesirable side effect of pollution and their finite expectancy. Over the past decade, numerous important technological advances in nanotechnology have opened up new frontiers in materials science and engineering, leading to the creation of new materials to meet the energy challenge. Metal—organic frameworks (MOFs), in particular, have proven to be indispensable for clean and efficient energy conversion as well as storage in fuel cells. MOFs offer several advantages as electrocatalysts, electrolyte membranes, and fuel storage materials — they possess remarkable design flexibility, ultra-large surface-to-volume ratios, and they allow functionalization with multivalent ligands and metal centers to increase avidity for fuel cells. Considerable efforts have been made to utilize the unique properties of MOFs as energy materials in developing high performance fuel cells. This article reviews the progress in the research and development of MOFs for applications in hydrogen fuel cells with an emphasis on fuel generation, catalysts for cathode, electrolyte membranes, and H2 storage, along with some discussion on challenges and perspectives in this exciting and promising field. © 2013 Elsevier Ltd. All rights reserved.

Introduction It was estimated that the world will need to double its energy supply by 2050 [1]. The need for clean and alternative sources of energy is and will continue to be the most compelling task of 21st century science. H2 powered fuel cells are one of the best alternatives to fossil fuels as they

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offer a clean and carbon-free method of converting chemical energy directly into electrical energy. Furthermore, they are almost twice as efficient as fossil fuels (60% efficiency for fuel cells versus 34% for fossil fuels) [2]. The low efficiency of combustion engines that utilize fossil fuels can be attributed to their acquiescence to the Carnot cycle laws — fuel cells are not under the jurisdiction of such laws [2,3]. If H2 is to become the utopian fuel, logically there must be functional fuel cells to serve as the vessels to actualize its potential. To date, many types of fuel cells have been proposed, such as polymer electrolyte membrane fuel cells (PEMFCs), direct methanol fuel cells (DMFCs), alkaline fuel cells, solid oxide fuel cells, phosphoric acid fuel cells, and molten

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Figure 1 Water—hydrogen cycle: H2 and O2 are generated by injecting energy into water and water is produced when energy is released from a hydrogen fuel cell at the point of use.

carbonate fuel cells [4]. The precise mechanisms of these fuel cells differ from one another, but the fundamental working principle — of directly converting chemical energy into electricity — remain unchanged. Most of these fuel cells utilize H2 as the fuel, but some, such as the DMFCs, utilize other substrates as fuels. This article will focus on fuel cells that use H2 as the fuel, with special emphasis on PEMFCs. H2 is regarded as an ideal fuel because the water—hydrogen cycle is closed (Fig. 1); the H2 and O2 produced from the splitting of water can be recombined into water with the release of energy, hence replenishing the water used in their production [2]. However, such sustainability is only maintained if renewable methods such as photocatalysis and electrocatalysis are employed for the splitting of water: fossil fuels cannot be involved in the splitting process. Inevitably, more energy is required to split water than what can be generated upon its recombination — current techniques can split water with about 80% efficiency but only generate electricity at about 60% efficiency. However, it should be noted that this is still much more efficient than the 34% efficiency obtained for fossil fuels, meaning that H2 is still a very enticing option. The most common type of fuel cells, both in research and in commercial use, is the PEMFCs. They are low temperature fuel cells that operate at temperatures of 85—105 ◦ C. A schematic of the components of a PEMFC, the reactions involved, and its working principle is shown in Fig. 2 [2]. As seen in Fig. 2, the PEMFC is essentially a galvanic cell that uses H2 as its fuel. H2 enters the anode where it is oxidized to protons. The electrons produced travel to the

Figure 2

Operating principle of PEMFCs.

cathode through an electrical circuit bearing an external load while the protons move through an electron-insulating but proton-conducting polymer membrane at the center of the cell. O2 at the cathode accepts the electrons together with protons and it is reduced to water. The PEMFCs are presently encumbered by a few limitations. Both reactions at the anode and cathode are inherently slow — therefore, fuel cells are forced to employ expensive noble metal catalysts such as platinum or platinum alloys to accelerate the reactions. However, the extremely high cost of noble metal catalysts impedes their widespread commercial usage. Furthermore, platinum-based catalysts are easily poisoned by carbon monoxide, which can be present in the fuel if it originates from fossil fuels. Thus, considerable attention is paid to the development of non-platinum electrocatalysts. The polymer electrolyte membrane (PEM) also plays an instrumental role in the PEMFCs. It needs to be proton-conducting and yet electron insulating, so as to allow protons to flow from the anode to the cathode while forcing the electrons to travel through the external circuit. To be industrially viable, the membrane needs to be cheap and resistant to corrosion, in addition to being a good proton conductor. Another issue is the source of H2 and O2 — chemicals that are required as reactants for the PEMFCs. In order for fuel cells to be truly carbon-free and environmentally friendly, H2 must not originate from fossil fuels. Furthermore, using fossil fuels as the source of H2 will pry open the closed water—hydrogen cycle — one of initial benefits that motivated the use of H2 as the fuel in the first place. As such, the ideal means of generating H2 are renewable methods such as photocatalysis or electrochemical splitting water. On-going research is currently directed toward synthesizing suitable catalysts for such water spitting reactions. However, several hurdles obstruct the large-scale commercialization of fuel cells, including the high cost of electrocatalysts, the need for adequate production of the fuels, as well as portability issues. Thus, there is a need to lower the cost and improve the performance of fuel cells before they are ready for widespread applications. Being an indispensable enabling technology in fuel cells, safe storage, and effective delivery of H2 has been actively pursued in fuel cell technologies. As an extremely volatile gas under ambient conditions, the volumetric energy density of H2 is too low for practical applications, and considerable attention must be paid to denser storage methods if H2 is to be used as a source of energy for mobile applications. Currently, high-pressure compression and liquefaction are the two technologies used in commercial H2 storage. Energy economy in the liquefaction and compression of H2 and safety concerns related to compressed H2 largely limit widespread applications of these two H2 storage technologies. Materials research holds the answers to many of the above-mentioned problems [5]. In particular, as a group of newly emerged porous materials, metal—organic frameworks (MOFs) have shown tremendous potential as versatile materials for a number of technical applications such as gas separation and storage [6,7], catalysis [8], drug storage and delivery [9], and energy conversion [10]. Comparing to conventional fuel cell materials, the unique attributes of MOFs, such as exceptionally high surface area, tunable porosity, uniform structured nano and microscale cavities,

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Metal-organic frameworks in fuel cells availability of in-pore functionality and outer-surface modification, good thermal stability, and diversity in the range of possible metal and functional groups of MOFs, render them very attractive for fuel cell applications. Moreover, a semi-infinite number of possible combinations of metal ions and organic linkers make it possible to synthesize MOFs with a wide range of properties. Combined with the other advantages of MOFs such as ease of synthesis and facile functionalization of organic linkers, MOFs are in a favorable position for achieving unprecedented performance in fuel cells. To date, over 20,000 MOFs have been reported in the literature, but only a small fraction of them have been thoroughly examined for their porosity, and far fewer have been tested for their possible usage in fuels [11]. Besides being tested as electrocatalysts and H2 storage materials in fuel cells, contrary to popular belief, a number of studies have shown that some carefully-structured MOFs are surprisingly good proton conductors. This article reviews the progress in the research and development of MOFs for fuel cell applications, especially on the advantages they could bring to fuel cells.

MOFs as promising materials in fuel cells MOFs are a new class of materials that have garnered significant attention in the last decade and several MOFs have been in industrial production [12—15]. MOFs are open networks consisting of metal-centered secondary building units (SBUs) joined together by organic linkers to form large one-dimensional (1-D), two-dimensional (2-D) or three-dimensional (3-D) networks. The structures are crystalline in nature, displaying long range order. Exceptionally uniformed pores or channels exist intrinsically within the framework, and they usually host guests such as solvent molecules (incorporated during the synthesis process) or counter ions which balance the overall charges on the framework (caused by the charged metal nodes). Unlike other types of porous materials such as zeolites and carbons, the vital and distinguishing property of MOFs that renders them unique and highly functional is that by careful selections of the metals and organic linkers, a ‘‘directed

3 tailoring’’ of their properties can be accomplished [9—16]. This flexibility allows us to engineer MOFs to cater to our specific needs, making MOFs highly versatile and adaptable materials. This tenability is significantly different from that of traditional porous materials like zeolites whose pores are confined by rigid tetrahedral oxide skeletons that are difficult to alter. The chemical structure of MOFs has a prevailing influence on their properties and applications. Moreover, because MOFs are constructed through the formation of networks of strong coordinative bonds between metal centers and organic linkers, they usually show good thermal stability up to 500 ◦ C. Also, good chemical stability has been observed with several MOFs although it remains challenging to make chemically stable MOFs due to their high susceptibility to ligand-exchange [12—15]. Therefore, a few keystone components of MOFs will be briefly introduced below, before proceeding to their applications in fuel cells.

Structural highlights of MOFs SBUs SBUs command the topology of MOFs (Fig. 3). The geometry and chemical attributes of the SBUs and organic linkers greatly help in the design and synthesis of MOFs. It was observed that under carefully selected conditions, multidentate ligands could aggregate and lock metal ions at certain positions, forming SBUs. Subsequently, these SBUs are joined by rigid organic linkers to produce extended frameworks of high structural stability [17]. SBUs have three elements essential to their polymerization into modular networks: (1) their core structural geometries are dictated by metal ions (also referred as metal nodes) that also serve as points of extension to other SBUs, leading to a predictable topology; (2) the coordination mode of organic linkers provides important geometric and conformational information to predict the topology of the resulting network; and (3) solvent molecules can serve as weak ligands filling the voids of the framework — evacuation of the solvent molecules produces periodically arranged open metal sites. In addition

Figure 3 SBUs commonly occurring in metal carboxylates (a) the square ‘‘paddlewheel’’, (b) the octahedral ‘‘basic zinc acetate’’ cluster, and (c) the trigonal prismatic oxo-centered trimer (blue: metal ions, black: carbon, and red: oxygen). Reproduced with permission from Ref. [16]. © 2004 Elsevier.

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to the structure of the ligand and type of metal utilized, the geometry of the SBUs is also dependent on the metal to ligand ratio, the solvent, and the source of anions required to balance the charges of the metal ions [18,19].

Open metal sites Metal sites in MOFs have a tremendous influence on their properties. The metals in MOFs often act as Lewis acids and they coordinate to labile solvent molecules or counter ions during synthesis. A subsequent activation of MOFs removes these weakly bonded species, resulting in the formation of unsaturated metal centers (sometimes referred to as open metal sites). Several MOFs have coordinatively unsaturated open metal sites that are built into the pore ‘‘walls’’ in a repeating and regular fashion [20]. These metal sites, such as those found in HKUST-1 (Cu-btc, btc = 1,3,5-benzenetricarboxylate) [21] and MIL-100 (metal carboxylate) [22], have been shown to impart catalytic activity to the materials [23,24]. The Lewis acid nature of the open metal sites can activate the coordinated organic substrates for subsequent organic transformations. Several groups have reported that MOFs with open metal sites act effectively as heterogeneous catalysts in organic synthesis [23—25]. For instance, the electron-deficient nature of some open metal sites in MOFs allows them to act as efficient oxidation catalysts for alcohols [26] and thioethers [27]. Additionally, the partial positive charges on the metal sites in MOFs have been demonstrated to possess the potential to enhance general adsorption properties, and this is explored as a strategy for increasing H2 adsorption in MOFs.

Pores in MOFs Pores are the void spaces formed within MOFs upon the removal of guest molecules. Mesopores (diame˚ ) are advantageous for conducting heterogeter = 20—500 A ˚ ) with high neous catalysis, while micropores (diameter <20 A specific surface areas are excellent materials for interacting with gases and vapors. Permanent porosity is more difficult to achieve in mesoporous MOFs than in their microporous counterparts. Generally, the likelihood of structural collapse increases with increasing pore size. In some cases, MOFs are able to interpenetrate one another to maximize packing efficiency, resulting in a large reduction of pore size despite the fact that the ligands are designed to generate large pores [28]. Nevertheless, such structures may lead to improved performance in some applications such as H2 storage [28].

Functional groups in MOFs It has been established that the choice of SBUs and organic linkers can simultaneously dictate the framework topology as well as impact the chemical functionality of the framework [15,29]. MOFs can incorporate a large number of different functionalities on organic linkers in a way that homogenously mixes them, rather than forming separate domains. Such an arrangement can result in properties that are distinctly different from the individual isolated domains [30]. In contrast, a minor modification to the side chains

of block copolymers often alters the entropy of the system and results in major and usually undesirable changes in the overall structure [31]. Therefore, it is believed that the ability to modify various functional groups in MOFs possesses great potentials across a wide range of technical applications [32]. More importantly, the facile functionalization of organic linkers (pre-synthesis route) represents another advantage of MOFs. It offers the option of splicing various functionalities and refining the physicochemical properties of the resulting MOFs. A number of MOFs with diverse pendant groups such as methyl, amine, halogen, as well as many other simple substituents lining the pore channels, have been synthesized [33—37]. Despite the great success in synthesizing a variety of functionalized MOFs via the pre-synthesis route, the scope of functional groups within the pores of MOFs remains relatively limited, owing to the fact that not all functional groups are compatible with the designed architectures and/or stable under solvothermal conditions. One critical factor that one must consider is that the functional groups in MOFs must not alter the topology of the MOFs. Such undesired interference might occur when the functional groups can interact with metal ions in a way that interferes with MOF synthesis [38]. Another viable approach toward installing functionalities in MOFs is by postsynthetic modifications (PSMs) [39—42]. Instead of synthesizing functional MOFs directly from functionalized organic linkers, functional MOFs can also be constructed in a heterogeneous manner via chemical modifications of the solid lattices of MOFs after they have been synthesized [39—42]. PSM could be more advantageous than the pre-synthesis route because both the SBUs and organic linkers can be modified without affecting the topology and stability of MOFs. Furthermore, functional groups that might interfere with the formation of the framework owing to their propensity to bind metal ions can be easily incorporated into the framework after it has been formed. In addition, as compared to the presynthesis route, PSM offers great flexibility over the types and number of functional groups that can be incorporated into MOFs. For instance, different functional groups can be incorporated topologically identical MOFs [43], while MOFs with different topologies can be modified using the same functional groups [44]. In this way, structure—function relationships of MOF topology versus MOF functionality can be methodically examined. In general, covalent, coordinative, and a combination of coordinative and covalent are the three principal PSM strategies for functionalizing MOFs. The covalent strategy deals with the modifications of organic linkers [45,46], whereas the coordinative strategy focuses on the modification of the coordination environment of SBUs within the framework without altering the framework topology [37,47—49]. Both the unsaturated metal modes of SBUs and the free coordinative groups of organic linkers can serve as the receiving ends during coordinative modifications. In addition, by combining the covalent and coordinative modification, additional ligands and metal ions can be appended to MOFs [50,51]. For example, by converting free amine groups into Schiff base ligands, vanadyl acetylacetonate was introduced into IRMOF-3 (Zn4 O(bdcNH2 )3 , bdc-NH2 = 2-amino-1,4-benzenedicarboxylate), and the resulting material was used as a heterogeneous oxidation catalyst [50].

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Metal-organic frameworks in fuel cells

Development of MOF porous structures MOFs are essentially formed via reticular synthesis by connecting SBUs with organic linkers [13]. The following inherent attributes are required to classify a solid as a MOF: strong bonding providing robustness, presence of linking units, and a geometrically well-defined crystalline structure [16]. Therefore, preparation of MOFs is conventionally achieved through self-assembly where metal ions and organic linkers are neatly organized to afford crystalline, porous networks. Most synthesis recipes for MOFs are based on solvothermal treatments that facilitate and encourage the formation of metal—organic SBUs. The ability to predict the network geometry is instrumental in understanding the subsequently observed properties of MOFs. Despite the apparent simplicity of reticular chemistry, one of the greatest challenges in the synthesis of MOFs lies in the optimizing experimental conditions. A slight change in the experimental conditions may have a significant impact on the formation of the desired MOFs. Moreover, single metal coordination geometry can propagate into more than one type of network and hence create more complexities in understanding the overall structure [52]. In addition, highly labile metal ions do not impose a strong preference for any particular geometry, leading to lack of predictability of the network structure. Likewise, the flexibility of organic linkers can result in a number of possible conformations, which confounds the anticipation of the framework geometry. The use of organic linkers with rigid backbones that reduce the free orientation of ligand lone pairs may offer a solution to the problems mentioned above [53]. Exploitation of recurring coordination motifs, either from existing polymeric structures or discrete model complexes, is also helpful in identifying the network geometry. Attention should also be paid to the influence of other factors such as charge-balancing anions and ring-opening polymerization on structural predictability [53]. Several advanced techniques have been developed to aid the design of MOFs. For example, the PLUXter method [54], which employs both hierarchical cluster analysis and principal components analysis, is able to provide grouping identities for further structure—property studies on representative structures from the clusters. MOFs grant access to porous structures with network topology and connectivity that are not usually observed in traditional porous materials such as zeolites and carbons. Their capability to generate unusually large diameter channels and cavities is particularly attractive. However, great care must be taken in the design and synthesis of mesoporous MOFs, because interpenetration may occur due to the use large linkers to create mesopores, thus resulting in the reduction of the effective pore volume created [55]. A feasible alternative is to design a framework in which the spaces occur topologically; for instance, the boron net in CaB6 , graphene nets, and diamonds are frequently used in the construction of highly porous MOFs [56—58]. This designed ‘‘node and spacer’’ approach has been proved to be very successful in the preparation of a wide variety of topologies, allowing the development of families of MOFs having the same connectivity but tenable pore dimensions. The presence of bulky counter ions within the cavities may also suppress interpenetration [59]. Additionally, if the

5 interaction between guest molecules and the framework is strong enough, framework cavities may be preferentially filled with the guest molecules rather than with additional polymer strands. Evidence can be found in the MOF Ni(L)2.5 (H2 O)2 (ClO4 )2 ·1.5(L)·2H2 O (L = ligand). X-ray crystallographic experiments revealed the existence of square Ni4 (L)4 pores which are large enough for interpenetration to occur [60,61]. However, these pores were filled with hydrogen-bonded guest aggregates comprising water molecules, perchlorate anions, and uncoordinated ligand without any network interpenetration. The use of bulky organic linkers coupled with synthesis under dilute conditions may also help to diminish the interpenetration effect. A common problem encountered during the synthesis of MOFs is that the frameworks are not stable to the loss of solvent molecules as they diffuse into the cavities and even template the structure around them. Exploitation of the chelate effect may assist in maintaining the frameworks during guest exchange [62]. Several MOFs were reported to undergo desolvation with the retention of the framework structures, while some of the desolvated materials have been found to reabsorb small molecules with considerable selectivity. For instance, Kitaura and co-workers reported the synthesis of a Cu-based MOF with 1,2-dipyridylglycol as a bridging ligand, in which the empty MOF supports the retention of the structure with only a slight change in interlayer distances [63]. Powder X-ray diffraction patterns of the MOF suggested that the adsorption and desorption processes are closely associated with the structural readjustment. The inherent properties of MOFs discussed above — crystallinity, porosity, functionality, and large surface areas — make them highly adaptable materials with applications in gas adsorption and storage [64], sensors [65], separation [66], luminescence [67], and catalysis [68,69]. Their highly structured and organized frameworks confer precious predictability upon their structures. This offers a systematic approach toward modifying the frameworks by rational alteration of the individual components in order to achieve desired outcomes. For instance, the size of the pores or channels in MOFs is easily tunable via altering the length of organic linkers; if small pores are required, a short organic linker can be used and vice versa for larger pore requirements. Similarly, the hydrophilicity of MOFs can also be controlled by choosing between polar or non-polar organic linkers. These intrinsic properties of MOFs imply that they have huge potentials for fuel cell applications, as MOFs can be customized according to the particular demands of fuel cells. This review aims to summarize the progress of the applications of MOFs in fuel cells. It is foreseeable that this research niche will gain momentum in the near future.

MOFs in H2 production H2 production from water splitting using MOFs In recent years, researchers have succeeded in utilizing MOFs as photocatalysts to generate H2 . One of the first reports was made by Kataoka et al. [70]. It was observed that Ru-based MOFs are capable of producing H2 from water through photocatalysis. The Ru-MOFs functioned as activity sites for the photochemical reduction of water into H2 in the

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Figure 4 Schematic illustration of catalytic mechanism of Ru-MOFs (X = Cl and Br). Reproduced with permission from Ref. [71]. © 2011 Taylor & Francis.

presence of Ru(bpy)3 2+ (bpy = 2,2 -bipyridine) as a photosensitizer, methyl viologen (MV2+ ) as an electron relay and EDTA as a sacrificial donor. To improve the photocatalytic efficiency, the Ru-MOFs were further modified by manipulating the counter ions within the framework [71]. The photochemical production of H2 from water using three heterogeneous MOFs (Ru2 (bdc)2 X), bdc = 1,4-benzenedicarboxylate (X = Cl− , Br− or BF4 − ) was studied. The Ru—Cl and Ru—Br complexes have a 3-D jungle-gym-like structure, where a 2-D square grid sheet of dinuclear RuII,III paddlewheel motif containing a bdc linker is bridged by the Cl− or Br− counter ions, extending the 2-D structure into a 3-D motif. When the Ru-MOFs were mixed together with Ru(bpy)3 2+ , MV2+ , and EDTA, H2 production was effected. The MV2+ molecules were adsorbed on the surface of the Ru-MOFs, not in the cavities. Hence, the difference in catalytic performance between the three MOFs is a result of surface modification, which in turn can be attributed to the selection of counter ions (Fig. 4). The most active catalyst reported in this study, Ru2 (bdc)2 Br, showed a turnover number (TON) of 18.7. This activity was 16.5 times greater than that of anatase TiO2 and 1.37 times better than that of Pt(bpy)Cl2 . Similar to the Ru-MOFs, photochemical production of H2 from water under irradiation from visible light was observed with a Rh2 (bdc)2 MOF in the presence of Ru(bpy)3 2+ , MV2+ , and EDTA [72]. The MOF was stable in the catalytic reaction and that the termination of the reaction was caused by other factors such as the hydrogenation of MV2+ . In another report, a photochemically active 3-D interpenetrated Zn-Pd MOF (ZnPd(ina)4 , ina = isonicotinate) showed a catalytic activity of 2.1 times greater than that of Pt(bpy)Cl2 [73]. However, prolonged irradiation resulted in the partial decomposition of the Zn-Pd MOF. Recently, Kataoka et al. synthesized another Zn-Pd MOF of having a 3-D porous structure constructed from Zn2 (H2 O)3 units linked by trans-PdCl2 (pydc)2 (pydc = 3,5-pyridine-dicarboxylate) metalloligands [74]. A total of 50.6 ␮mol of H2 was generated

after a period of 4 h irradiation, translating to a TON of 20.2. Unfortunately, this Zn-Pd MOF was degraded and displayed no further catalytic activity after only 5 h of irradiation. Further work is needed to improve the lifespan of the Zn—Pd MOF before it can be employed in practical scenarios. Such work will be worthwhile because the photocatalytic activity of the Zn-Pd MOF is much greater than the previously reported MOFs. More recently, Kataoka’s group reported the photochemical properties of heterogeneous microporous porphyrin coordination lattices (PCLs) [75]. Two PCLs, PCL-1 (Ru2 (H2 tccp)BF4 ) and PCL-2 (Ru2 (Zntccp)BF4 ) (tccp = tetrakis(4-carboxyphenylporphyrin)), were synthesized. Similar to previous observations, physical adsorption of MV2+ on the surface of the PCLs was essential for the catalytic process, where the adsorbed MV2+ served as intermolecular electron relays from the porphyrin sites to the Ru centers. When MV2+ , PCL, and EDTA were irradiated at 320 nm, PCL-1 gave a TON of 20.8 while PCL-2 had a TON of 29.9. The porphyrin itself functioned as a photosensitizer, hence there is no need for Ru(bpy)3 2+ . The high TON values imply that the incorporation of porphyrin into the MOFs as the photosensitizer may lead to the improvement in catalytic efficiency, and this could motivate further research into porphyrin-based MOF photocatalysts. Besides the above VIIIB metal-based MOFs, Garcia and his colleagues reported that isoreticular Zr-based MOFs catalyze the photochemical generation of H2 from a water/methanol mixture [76]. UiO-66 (Zr6 O4 (OH)4 (bdc)12 ) and UiO-66(NH2 ) (Zr6 O4 (OH)4 (ata)12 ) (ata = 2-aminoterephthalate) are isoreticular MOFs that are water-stable at 1000 ◦ C. 2.4 and 2.8 mL of H2 were obtained after 3 h of irradiation for UiO-66 and UiO-66(NH2 ), respectively. The presence of NH2 groups in the linker of UiO-66(NH2 ) increased the amount of light absorbed, and therefore led to the observed enhanced activity for H2 generation.

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Metal-organic frameworks in fuel cells

H2 production from ammonia borane and organosilanes using MOFs Apart from water, ammonia borane (AB) has been gaining popularity and capturing the attention of researchers as a potential on-board fuel for fuel cells. AB qualifies as an onboard H2 storage medium because it is stable in air and has a moderate dehydrogenation temperature. Furthermore, it has high a hydrogen content of 19.6% or 140 g/L. The hydrogen stored within AB can be released via thermolysis or by reaction with water. Awkwardly, AB is plagued by sluggish H2 release kinetics at the fuel cell operating temperature of 100 ◦ C, as well as the formation of undesirable side products such as ammonia, borazine, and diborane during its decomposition. These volatile by-products are poisonous to the electrocatalysts. Thus, technical challenges that need to be addressed include reducing the dehydrogenation temperature to <85 ◦ C, significantly increasing the H2 release rate, and preventing the formation of poisonous by-products. Studies by several groups have demonstrated that MOFs have the potential to alleviate the above-mentioned problems and increase the viability of AB as an on-board fuel. Li et al. noted a synergistic effect of the nanoconfinement of AB together with the use of a MOF in improving H2 release kinetics and in preventing the formation of the poisonous by-products (Fig. 5) [77]. After AB was loaded into JUC-32Y (Y-btc) by infusion to form AB/JUC-32-Y, X-ray diffraction tests proved that the structure of JUC-32-Y remained intact. Thermal decomposition of the trapped AB occurred at temperatures as low as 50 ◦ C, as compared to 112 ◦ C required in decomposing pristine AB. At 85 ◦ C, AB/JUC-32-Y released 8% of hydrogen within 10 min and 11% in 3 h. If temperature was raised to 95 ◦ C, AB/JUC-32-Y was able to release all of its hydrogen (13%) in 3 h. Indeed, it was observed that there is no formation of the undesirable by-product ammonia. It is postulated that unsaturated Y3+ metal centers in JUC-32-Y interact with the NH3 moiety in AB and strengthen the B N bond, thereby preventing ammonia formation. As such, MOFs naturally arise as promising candidates to study this synergistic effect. This is because MOFs are intrinsically

Figure 5 Time dependence of H2 release from AB-JUC-32-Y and neat AB. Reproduced with permission from Ref. [77]. © 2010 American Chemical Society.

7 porous and hence suitable for nanoconfinement of AB. Moreover, their metal nodes can serve as the catalysts to foster the observed synergistic cooperation between nanoconfinement and the catalyst to enhance H2 release kinetics. The uniform channels present within microporous MOFs readily allow small molecules such as AB to access the interior surface and react with the active metal sites. The porosity of MOFs also promotes the uniform dispersion of reactant AB molecules throughout the framework to increase the reaction surface area. Gadipelli et al. synthesized Mg-MOF-74, (Mg2 (dobdc), dobdc = 2,5-dioxido-1,4-benzenedicarboxylate) containing 1-D hexagonal channels, for the thermolysis of AB [78]. Upon heating under vacuum, the terminal water molecules attached to the MOF were removed to yield open, unsaturated Mg metal sites for catalytic purposes. In addition, AB molecules can be confined within the pores of the MOF, with a loading of one AB molecule per Mg2+ , which is equivalent to 26% of AB, to form an AB-Mg-MOF-74. It has been shown that the electropositive boron in AB interacts with the oxygen groups in the dobdc ligand and the NH3 moiety in AB interacts with Mg2+ in the AB-Mg-MOF-74. These interactions weaken the N B and B H bonds, which ultimately results in enhanced desorption kinetics and a low dehydrogenation temperature. At low loadings of AB, AB molecules are confined within the pores of the AB-Mg-MOF-74, whereas excess AB molecules are aggregated outside the saturated pore channels at high loadings. It was observed that only those AB molecules confined in the channels possess enhanced dehydrogenation kinetics. For the low-loading AB-Mg-MOF-74, the dehydrogenation of AB was a one-step process taking place at ∼70 ◦ C without the production of ammonia, borazine or diborane. H2 was generated instantly without the need for an induction period once the temperature reached 65 ◦ C. At the optimal operating temperature of 85 ◦ C for PEMFCs, 9% H2 was generated in 90 min. This is a significant 50% improvement from the 6% H2 obtained from pristine AB at temperatures ranging from 85 to 125 ◦ C over a few hours with a few additional hours of induction time. Likewise, enhanced dehydrogenation kinetics and clean H2 generation were observed with Zn-MOF-74, (Zn2 (dhbda), dhbda = 2,5-dihydroxybenzene1,4-dicarboxylic acid) [79]. AB molecules were incorporated into the hexagonal 1-D pores of the Zn-MOF-74 to form a hybrid framework AB-Zn-MOF-74. An AB:Zn ratio of 1:1 was readily obtained in the AB-Zn-MOF-74, corresponding to 19% by weight of AB in the MOF. Similar to the Mg-MOF-74, the enhanced dehydrogenation is due to unsaturated Zn centers in the AB-Zn-MOF-74. The AB-Zn-MOF-74 was able to deliver ∼10% of H2 at 60 ◦ C in a few minutes without any induction time. More importantly, the AB-Zn-MOF-74 exhibited complete suppression the formation of byproducts, including ammonia, borazine, and diborane, thus preventing the poisoning of the electrocatalysts. The performance of the AB-Zn-MOF-74 is superior to AB-Mg-MOF-74. Hopefully, further investigations will eventually lead to further enhancements in the dehydrogenation performance of AB and the implementation of AB in on-board fuel cells. Burrell’s group discovered that ZIF-8 (Zn(mim)2 , mim = 2-methylimidazolate) promotes the release of H2 from AB [80]. In addition to the synergistic effect of the nanoconfinement and AB-metal center interaction, it is

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postulated that the homogenous dispersion of Zn ions on the surface of ZIF-8 and the disruption of the hydrogen bond network within AB are responsible for the enhanced kinetics. However, it is also possible that H-bonding and non-specific host—guest interactions are responsible for the enhanced H2 release kinetics. Furthermore, ZIF-8 was demonstrated to have a strong ability in immobilizing Ni nanoparticles, preventing their aggregation and therefore increasing their catalytic surface area [81]. In this respect, ZIF-8 is an excellent matrix to control the size and distribution of the nanoparticles because of its high thermal stability (up to 550 ◦ C in N2 ) and excellent stability in aqueous medium. In this case, ZIF-8 was not directly used for AB catalysis; instead, the immobilized Ni nanoparticles on the ZIF-8 serve as the catalytic sites. The catalytic process was postulated to take place via the formation of an activated complex species which arise from interaction between AB molecules and the metal nanoparticle surface. In a separate study, Li and his colleagues synthesized two Ni-MOFs, Ni-MOF-1 (Ni(4,4 -bipy)(btc), 4,4 -bipy = 4,4 bipyridine) and Ni-MOF-2 (Ni(pyz)Ni(CN)4 , pyz = pyrazine) which were used as precursors in generating Ni-based catalysts used in the thermolysis of AB [82]. Upon exposing the MOFs to AB in methanol, some of the Ni2+ ions were reduced to Ni0 . Due to the unique structure of the MOFs, the metallic Ni generated is different from normal Ni nanomaterials; they are likely to be present in well-dispersed tiny clusters that bestow them with high dehydrogenation catalytic activity. There was a greater degree of reduction in Ni-MOF-1, while Ni-MOF-2 had better preservation of its framework. This difference arises largely from the structure difference of the two MOFs — due to the greater number of Ni sites in Ni-MOF-1, Ni-MOF-1 lowered the dehydrogenation temperature more effectively. For instance, Ni-MOF-1 released 7.0% H2 within 40 min without any induction time while Ni-MOF-2 produced 6.0% H2 within 20 min at 90 ◦ C. This is a remarkable rate of dehydrogenation as compared to the less than 0.3% H2 generated from pristine AB under the same conditions. The activation energies for the generation of H2 were about 131 ± 5 kJ/mol for Ni-MOF-1 and 160 ± kJ/mol for Ni-MOF2, respectively. Both are lower than the 184 ± 5 kJ/mol for pristine AB. There was also an observed decrease in reaction exothermicity compared to pristine AB, which facilitates the regeneration of spent fuel. The same concept was applied to the synthesis of a highly efficient Co-based catalyst derived from a MOF for the hydrolysis of AB [83]. However, in this study H2 was generated via hydrolysis instead of thermolysis, according the following equation: NH3 BH3 + 2H2 O → NH4 + + BO2 − + 3H2 The advantage of producing H2 via hydrolysis over thermolysis is that hydrolysis of AB does not release volatile and harmful by-products. An amorphous Co0 catalyst for the hydrolysis of AB was synthesized via the reduction of a Co-MOF (Co2 (bdc)2 (dabco), bdc = 1,4-benzenedicarboxylate and dabco = 1,4-diazabicyclo[2.2.2]octane). The precursor Co-MOF has a 3-D pillared structure with an accompanying 3-D pore system. Such interconnected channels coupled with a large surface area render the metal sites in the

Co-MOF highly accessible. When AB is added to a mixture of Co-MOF and NaBH4 , the framework of Co-MOF collapses and the amorphous Co0 catalyst is formed in situ through the reduction of the Co-MOF. The resulting catalytically active Co0 sites are separated by the residue of the precursor Co-MOF and stabilized by the organic linkers. A H2 generation rate of 932 L/mol/min was obtained in 0.32 M AB. Notably, this level of catalytic activity is comparable to some noble metal catalysts. Unfortunately, Co(OH)2 is formed during the reaction and deposited on the surface of the catalyst when it is exposed to air, resulting in a significant decrease in catalytic activity. This exposure to air can be conveniently avoided in fuel cells if the catalyst is kept in an aqueous solution in a H2 atmosphere. In addition to water and AB, another plausible source of H2 for fuel cells is organosilanes. Ison et al. demonstrated the hydrolytic oxidation of organosilanes to produce H2 [84]. H2 was produced via the catalytic hydrolytic oxidation of an organic liquid (organosilanes, (R4−x Si(OH)x ), R = alkyl or aryl, x = 1, 2 or 3) using an oxorhenium(V) oxazoline coordination complex. The reaction can be carried out under ambient conditions in less than 1 h, without the need for a solvent. The amount of H2 generated decreases in the order primary silane > secondary silane > tertiary silane, because the energy required to split water originates from the formation of a Si O bond. Although the cost of organosilanes is generally high, this problem can be overcome by efficient recycling of the organic silicon by-products. This approach deserves more attention when designing MOF-based catalysts for H2 production as it allows H2 to be produced using an organic liquid and only water as a co-reagent. Despite the possible benefits of generating H2 using MOF catalysts from organosilanes, no progress has been made in this area so far.

MOFs as oxygen reduction reaction catalysts The H+ ions formed during the oxidation of H2 at the anode are reunited with the electrons they previously lost at the cathode. Together with O2 , they combine to produce water. This is the oxygen reduction reaction (ORR) that occurs at the cathode of a fuel cell. ORR is typically catalyzed by Pt-based electrocatalysts because the sluggish kinetics of ORR [2,3,85]. Compared to the H2 oxidation at the anode, ORR is the rate-determining step in H2 fuel cells. As such, slow ORR can be considered to be the Achilles heel of fuel cells and it typically prevents fuel cells from achieving their ideal efficiency. Furthermore, the extortionate price of the Pt-based ORR electrocatalysts employed in present-day fuel cells is responsible for half of the fuel-stack cost. Consequently, intensive efforts have been devoted to the development of novel electrocatalysts for the ORR in fuel cells in order to enhance their practical applicability. Among various proposed electrocatalysts, non-platinum group metals (non-PGM) materials are generally more appealing, mainly because they are more readily available and much less costly. However, the downside is that non-PGM materials are generally less efficient and greater amounts of catalysts have to be used to achieve a reasonable catalytic efficiency, resulting in a thick electrode layer that causes poor mass transport. Hence, the contemporary challenges are to achieve better catalytic

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Metal-organic frameworks in fuel cells powers and higher catalytic site densities using non-PGM ORR electrocatalysts. Additionally, the electrocatalysts should be evenly distributed throughout the electrode and easily accessible by both gaseous and liquid substrates. With the above concerns in mind, MOFs emerge as particularly viable contenders as novel materials for ORR catalysis — they contain numerous metal centers that are arranged in a highly organized and predictable manner, rendering them easy to modify and manipulate. Their well-defined 3D structures provide the highest possible volumetric density through regularly arranged cell structures. Lastly, the pore size of MOFs can be fine-tuned by modifying the length of organic linkers, allowing for optimal diffusion of gases and liquids. With these well-suited intrinsic characteristics, it is not surprising at all that MOFs have garnered a great amount of research interest as prospective materials in the pursuit of ORR electrocatalysts [85]. One of the pioneering reports on the use of MOFs as non-PGM electrocatalysts was by Dodelet’s group who experimented with Fe-based cathode catalysts for the H2 -air fuel cell [86]. Lefevre and his colleagues worked on Fe/N/Ccatalysts, using ZIF-8 as the host for Fe and N precursors. ZIF-8 was chosen because it possesses the two factors critical for ORR activity of the Fe/N/C-catalysts — high nitrogen content and high microporous surface area. Previous investigations have demonstrated that the incorporation of nitrogen in carbon materials (especially in the form of pyridinium moieties) is instrumental in boosting the electrocatalytic activity for ORR. To prepare the best performing catalyst, a mixture of ZIF-8, 1,10-phenanthroline, and ferrous acetate was ball-milled and pyrolyzed twice, first in argon at 1050 ◦ C, then in NH3 at 950 ◦ C. As a proof-of principle, experiments using this catalyst produced a power density of 0.91 W/cm2 at a voltage of 0.6 V — a value that rivals that of the state-of-the-art Pt-based electrocatalysts with a loading of 0.3 mg/cm2 . The remarkable catalytic activity is mainly attributed to enhanced mass-transport properties arising from the interconnected alveolar carbon nanostructure with numerous pores that facilitate the movement of water and oxygen molecules. Regrettably though, the ZIF-8 derived catalyst did not fare as well in durability and stability studies.

9 Instead of using iron or zinc as the metal centers in MOFs, Ma et al. synthesized a similar MOF-based electrocatalyst with Co-imidazolate frameworks that also incorporates nitrogen as a crucial component in its structure (Fig. 6) [87]. One of the reasons for choosing cobalt is because Co-based ORR catalysts have a better tolerance for the acidic environment. The electrocatalyst precursor Co-MOF was synthesized by reacting 3,5-imidazolate and Co(NO3 )2 in N,N -dimethylacetamide under solvothermal conditions. In the Co-MOF every Co atom bonds with four N atoms from four 3,5-imidazolate ligands, and each 3,5-imidazolate ligand connects with two Co atoms to form an extended 3-D porous structure. Electrochemical measurements on freshly-prepared Co-MOF revealed no ORR activity due to its insulating nature. Upon thermal activation at 600—750 ◦ C, the Co-MOF exhibited ORR activity at onset potentials ranging from 0.77—0.83 V with a dominative four-electron process with water as the main reduction product and peroxide as a minor side product. However, further increase in pyrolyzing temperatures led to deterioration of the ORR activity. There was a direct correlation between the surface area and the catalytic activity. Most of the surface area was contributed by the mesopores formed during thermal activation. The active sites for catalysis were believed to be formed by the transformation of carbon and nitrogen in the imidazolate linkers during thermal activation. At elevated temperatures, the imidazolate moieties slowly convert into carbonaceous forms, while a portion of the nitrogen is retained as pyridinic- or pyrolic-like moieties, thus producing a nitrogen-containing graphene-like structure accompanied by the formation of mesopores, with the retention of a certain fraction of the micropores. Meanwhile, the coordination of nitrogen to the Co centers forms the electrocatalytic sites. Although the original MOF structure is not retained after the thermal activation, it plays a pivotal role in ensuring the even distribution of Co and Nfunctional groups throughout the 3-D porous architecture, resulting in densely populated active sites while retaining adequate porosity and surface area — factors which are of paramount importance in non-PGM electrocatalysts. Copper-based MOFs have also been tested for possible electrocatalytic activities in ORR [88]. It was observed that

Figure 6 (a) Local Co-N4 coordination moiety, (b) structure packing of the Co-MOF along the [100] direction and (c) faradaic current density as a function of potential in O2 saturated 0.10 M HClO4 at 25 ◦ C. (䊉) Fresh Co-MOF and Co-MOF treated at () 500, () 600, () 700, () 750, () 800 and () 900 ◦ C. Reproduced with permission from Ref. [87]. © 2011 Wiley-VCH.

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HKUST-1 displays electrocatalytic activity toward ORR, but it is structurally unstable in aqueous media. Another Cubased MOF, (Cu-bpy-btc, Cu2 (OH)(bpy)2 (btc)3 ·2H2 O), which has btc as primary ligands and 2,2 -bipyridine as auxiliary ligands (that serve to stabilize the skeletal structure via strong binding to the Cu centers) demonstrated better stability. Cu-bpy-btc retained its structural integrity even after being exposed to water for 24 h. When immersed in oxygenated 0.10 M phosphate buffer, an electrode modified with Cu-bpybtc showed a substantial increase in the reduction current at −0.20 V, signifying that O2 is catalytically reduced by Cubpy-btc. The number of electrons involved in the reduction was determined to be 3.8, again suggesting that the main product of the reduction is water, a much preferred ORR process. Electrocatalysts for ORR have also been fashioned from a graphene-metalloporphyrin (5,10,15,20-tetrakis(4carboxyl)-21H,23H-porphyrin) MOF composite synthesized from pyridine functionalized graphene and Fe-porphyrin [89]. It was found that the contents of graphene and Feporphyrin can influence the crystallization process of the MOF and enhance the electrocatalytic properties of the composite when appropriate amounts are added. Again, four electrons participated in the ORR regardless of the potential tested. Furthermore, chronoamperometric tests showed that the composite electrocatalyst is more durable than its Pt and Ni counterparts. Although adaptation of MOFs as electrocatalysts in fuel cells is highly desirable, and several studies have presented encouraging results in ORR, MOFs possessing electrocatalytic activity and durability rivaling that of platinum remains elusive. Doping of MOFs with PGM nanoparticles might be unavoidable if one wants to significantly improve the performance of the MOF electrocatalysts while keeping the cost relatively low. Besides the contribution from the significantly enhanced surface area, nanoconfinement effect may also contribute to the catalytic activities of MOFs. It has been observed that Pt-based nanoporous catalysts show near Nernstian behavior, faster response times, and less hysteresis in comparison with their nonporous counterparts with similar surface areas [90]. The increased residence time, due to the extremely confined spaces within the nanoporous structure, is likely to be the extra contribution [90]. Although other factors such as the defect and

Figure 7

stress effect of MOFs may also contribute to the enhancement in the electrocatalytic reactions, it is valuable to understand the role of geometric confinement for enhancing electrochemical reactions in MOFs. In addition, the structure-induced promotional effect may also contribute to the catalytic ORR as both the inner and outer surfaces of MOFs are easily accessible to O2 , while the interconnected MOF skeleton, branching in all dimensions, provides paths for fast electron transport, leading to high current density. It is believed that these novel nanoporous composites will open up new opportunities for ORR as well as many other electrocatalytic applications.

MOFs as proton-conducting polymer electrolyte membranes The functionality of PEMFCs is dependent on the effectiveness of PEM. Apart from preventing disastrous fuel crossover which ultimately leads to voltage loss, the PEM must also be proton conducting yet electrically insulating. Proton conduction across the PEM is believed to occur via two mechanisms — the Grotthuss mechanism [91] and the vehicle mechanism [92]. Both can be thought of as involving proton carriers; however, the proton carriers in the vehicle mechanism are mobile (as the name suggests) while the proton carriers in the Grotthuss mechanism are immobile. A schematic representation of the two mechanisms is shown in Fig. 7. As depicted in Fig. 7, in the Grotthuss mechanism, the protons are transported through a network of hydrogen bonds in two steps. The first step involves the movement of a proton from one proton carrier to another without translation of the carrier. In the second step, the proton carrier is reorientated such that it is ready to receive another proton. This process is repeated and the proton is carried along by permutations of hydrogen and covalent bonds. Proton conduction by the Grotthuss mechanism is usually accompanied by activation energies between 0.1 and 0.4 eV. The most common proton carrier is water, but other molecules that are capable of hydrogen bonding, such as carboxylates and imidazoles, can also function as proton carriers. In the vehicle mechanism, the protons are not transported as H+ but instead they are attached to a vehicle such as H2 O or NH3 and transported together with the carrier as H3 O+ or

Illustration of proton conduction models: the Grotthuss mechanism (top) and the vehicle mechanism (bottom).

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Metal-organic frameworks in fuel cells NH4 + ions. The ‘‘unloaded’’ vehicles move in the opposite direction of the ‘‘loaded’’ vehicles so as to pick up more protons to ensure continuity of the process. Proton transfer via the vehicle mechanism usually has activation energies above 0.5 eV. In MOFs, both the Grotthuss and vehicle mechanisms are believed to contribute to proton conductivity. At the first look, it might seem a little counterintuitive to use MOFs as proton-conducting membranes in fuel cells. This is because MOFs are intrinsically crystalline and porous materials — properties that are not welcome in proton conductors [93]. Porosity is generally not favored in PEMs as it increases the risk of fuel crossover from the anode to the cathode which results in voltage loss. Crystallinity is similarly disfavored because transport of ions is generally favored by disordered systems instead of ordered frameworks. However, upon closer examination and analysis, MOFs are actually promising candidates for membranes in PEMFCs. This is because the guest molecules in the pores of MOFs can actually block the pores to prevent catastrophic fuel crossover [94]. Concurrently, the counter ions that are present in the channels of some MOFs can help to conduct protons if they contain water, acidic groups, hydroxyl groups, or simply functional groups that contain hydrogen atoms. To date, many groups have reported proton-conductive MOFs based on guest molecules or counter ions present in their channels. Currently, MOFs serving as PEM is the largest research area for MOF applications in fuel cells. Kitagawa’s group was among the first to investigate the proton conductivity of MOFs [95—100]. They first worked on several proton conducting Cu-based MOFs with the general formula R2 dtoaCu (dtoa = dithiooxamide) and varying the identity of the R group. These MOFs have a 2-D structure consisting of Cu-dimeric units and their bridging ligands. Water molecules are present in these MOFs and their concentration increases with increasing relative humidity (RH). Since it has been observed that conductivity also increases with increasing RH, it has been proposed that the water molecules play a part in the proton conduction mechanism. Initial studies showed considerably high proton conductivities of 1 × 10−6 S/cm at 75% RH [95] and 4.2 × 10−6 S/cm at 100% RH [98] for H2 dtoaCu and (H5 C2 )2 dtoaCu, respectively. In an attempt to further boost the proton conductivity, OH groups were introduced to the R moieties of the MOFs, with the hope that OH groups will aid in the formation of hydrogen bond networks that facilitate proton conduction. As expected, a much improved proton conductivity of 1.2 × 10−5 S/cm under RH of 83% and 300 K was obtained with (HO-H4 C2 )2dtoaCu [96]. Since more water molecules can fit into its interplanar spacing of longer ligands [97], an even higher proton conductivity should be observed. However, thermogravimetric analysis on (HO-H4 C2 )2 dtoaCu and (HO-H6 C3 )2 dtoaCu showed that the same amount of water is present per dimeric unit under 100% RH for both MOFs. A proton conductivity of 2.0 × 10−6 S/cm was obtained at 100% RH for (HO-H6 C3 )2 dtoaCu. This suggests that the length of the hydrophobic moiety might not have a significantly positive effect on the proton conductivity. Similarly, OH groups were introduced onto three oxalatebridged MOFs, NH(prol)3 [MCr(ox)3 ] (M = Mn(II), Fe(II), Co(II) and NH(prol)3 + = tri(3-hydroxylpropyl)ammonium and ox = oxalate) by engaging hydrophilic NH(prol)3 + ions during

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Figure 8 Dependence of conductivity on relative humidity of (䊉) MnCr(ox)3 , () FeCr(ox)3 , () CoCr(ox)3 and () N(n-C4 H9 )4 MnCr(ox)3 at 25 ◦ C. Inset: structural view of the bimetallic MOF. Reproduced with permission from Ref. [98]. © 2010 American Chemical Society.

synthesis (Fig. 8) [99]. The structure of the MOFs consists of oxalate-bridged bimetallic layers that are intercalated by NH(prol)3+ ions. The bimetallic layer forms a 2-D honeycomb network while the tri(oxalato)chromate(III) (Cr(ox)3 ) ligand bridges three M(II) ions via its oxalate groups. Meanwhile, the NH(prol)3 + ions are hydrogen bonded to the bimetallic layers. One water molecule is within the honeycomb cavity and another water molecule is located on the opposite side of the bimetallic layer and is hydrogen-bonded to one MCr(ox)3 . Proton conductivity of these oxalate-bridged MOFs was found to increase with RH until ∼80%, after which the MOFs began to decompose. The increase in proton conductivity with increases in RH was rationalized using water adsorption isotherms — more water molecules are adsorbed on the MOFs at higher RH, and this aids in proton conduction. The MOFs achieved a sizeable proton conductivity of about 1 × 10−4 S/cm at 75% RH. Evidently, the substantial proton conduction is caused by the 2-D hydrophilic layers formed by the NH(prol)3 + ions. These results illustrated that oxalate-bridged bimetallic complexes containing hydrophilic ions have the potential to be good proton conductors. To further improve the proton conductivity, another oxalate-bridged MOF, (NH4 )2 (adp)-[Zn2 (ox)3 ]·3H2 O (adp = adipic acid), was synthesized [100]. This Zn-based MOF consists of 2-D honeycomb sheets of (Zn2 (ox)3 )2− . The adipic acid molecules penetrate the hexagonal voids of the layers with their carboxyl groups projecting into the interlayer space of (Zn2 (ox)3 )2− sheets. NH4 + ions and water molecules occupy the pores next to these hydroxyl groups. The close proximity of the three groups results in the formation of a 2-D hydrogen bond network among adipic acids, NH4 + ions, water, and oxalate ions of the framework. Moreover, additional protons from adipic acid and NH4 + make the framework highly proton conductive. Consequently, a proton conductivity as high as 8 × 10−3 S/cm at 298 K and 98% RH was obtained, which is comparable to the proton conductivity of 2.0 × 10−3 S/cm obtained

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12 with Nafion — the gold standard of proton conducting electrolyte membrane in fuel cells [101]. The activation energy of the proton conduction was 0.63 eV, suggesting that the proton conduction occurs both by the Grotthuss mechanism via the 2-D hydrogen bond network and the vehicle mechanism, where additional protons from NH4 + and water diffuse throughout the framework. Additionally, the effect of altering the hydrophilicity of counter ions on the proton conductivity was investigated with model oxalate-bridged layered MOFs with the general formula NR3 -CH2 COOH(MCr(ox)3 )·nH2 O (R = Me (methyl), Et (ethyl) or Bu (n-butyl), and M = Mn or Fe) [102]. These MOFs have a 2-D honeycomb-based structure formed by the linkage of the metal ions to the [Cr(ox)3 ]3− units. The cationic NR3 (CH2 COOH)+ is aligned in a perpendicular manner to this 2-D layer, with their carboxyl groups occupying the honeycomb cavity. Meanwhile, the NR3 group of the cation occupies the interspace, and therefore the hydrophilicity of the interspace can be controlled by the identity of NR3 (either Me, Et, or Bu). The carboxyl groups in the MOFs function as proton carriers and their hydrophilicity is tuned by the attached NR3 moieties. It was shown that the greater the hydrophilicity, the greater the adsorption of water, and the higher the proton conductivity. Me-FeCr is the most hydrophilic MOF among them, due to the smallest size of the alkyl groups. Therefore, it adsorbed the greatest amount of water at RH above 40% as compared to the rest. A proton conductivity of ∼10−4 S/cm was observed even at low RH values. The above observations may motivate further research into fine-tuning the hydrophilicity of the counter ions to improve the proton conductivity of MOFs. Yamada et al. discovered that one of the simplest MOFs of having reasonable proton conductivity is ferrous oxalate dihydrate, commonly known as Humboldtine [103]. Humboldtine has a 1-D structure formed by ferrous ions serving as metal nodes with oxalate ligands bridging the nodes together. Two water molecules coordinate to the ferrous ion to form another 1-D array of water molecules. This ordered array of water molecules is suspected to be responsible for creating a pathway for proton conduction. It was found that Humboldtine displays a high proton conductivity of 1.3 × 10−3 S/cm at 25 ◦ C and 98% RH, comparable to the gold standard — Nafion. Humboldtine also had a very low electrical conductivity of <10−10 S/cm, further justifying its use as an electrolyte membrane. The activation energy was estimated to be 0.37 eV, thus suggesting a Grotthuss mechanism. This high proton conductivity is probably the direct consequence of the confined nanoarray column created by the 1-D nanoarray of coordination waters, which form hydrogen bonds with the oxygen atoms of the ferrous oxalate framework. Previous studies have demonstrated that the ionic conductivity of nanospaced ion conduction materials is greater than that of their bulky counterparts [104—107]. Similar to Humboldtine, it was noted that an oxalatebased 3-D chiral MOF, (NH4 )4 (MnCr2 (ox)6 )·4H2 O also utilizes water in its proton conducting mechanism (Fig. 9) [108]. (NH4 )4 (MnCr2 (ox)6 )·4H2 O harbors water molecules within its type A functionalized channels (Fig. 9 inset). And it also houses ammonium cations and two other water molecules around the framework. It was hypothesized that the proton conductivity of (NH4 )4 (MnCr2 (ox)6 )·4H2 O is boosted by situating water and ammonium ions in the framework close

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Figure 9 Dependence of the conductivity on relative humidity at 295 K. (Inset: view of the crystal structure along the c-axis of the bimetallic MOF, Cr(III) (pink) and Mn(II) (green) polyhedra, respectively.) Reproduced with permission from Ref. [108]. © 2011 American Chemical Society.

to the water molecules in the channel. The close proximity of the ammonium ions and water molecules around the framework with the water molecules in the channels lead to strong interactions between them, and results in a high proton conductivity of 1.1 × 10−3 S/cm at room temperature and 96% RH. Another example of utilizing water for proton conduction is a phosphonate-based MOF synthesized by Shimizu and his colleagues [109]. They rationalized that phosphonate groups are good candidates for proton conduction — they possess three oxygen atoms, leaving two ‘‘free’’ oxygen atoms to act as H-bond acceptors after one has been used to coordinate to the metal node. The proton conductivity in H2 was measured as 3.5 × 10−5 S/cm at 25 ◦ C and 98% RH with an activation energy of 0.17 eV, signifying that the Grotthuss mechanism is the basis of proton conduction. Sahoo et al. investigated the role of halogens and helical water chains in affecting proton conductivity within MOFs [110]. Two MOFs, Zn(L-LCl)(Cl)(H2 O)2 and Zn(D-LCl)(Cl)(H2 O)2 , (L = 3-methyl-2-(pyridin-4ylmethylamino)-butanoic acid), displayed proton conductivities of ∼4.45 × 10−5 S/cm at ambient temperature. Both MOFs have a 3-D supramolecular network containing a close packed 1-D open channel along the c-axis, which is occupied by water molecules. The walls of the channels are hydrophobic in nature due to the presence of pyridyls and isopropyl groups. In the channels, one water molecule forms a weak hydrogen bond with the Cl atom from the M-Cl group, and a second water molecule forms a hydrogen bond with the first water molecule; this process repeats itself throughout the channel, thereby forming a continuous helical water chain. This ordered helical chain allows protons to be transported via the Grotthuss mechanism. Cabeza and co-workers synthesized a flexible, cross-linked, ultra-microporous magnesium tetraphosphonate hybrid framework (CUMTHF) that similarly employs water in its proton conduction mechanism [111]. CUMTHF has a 3-D pillared open framework with crosslinked 1-D channels filled with water and DMF molecules

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Metal-organic frameworks in fuel cells and possess considerable framework flexibility by virtue of the aliphatic (CH2 )n chain in the ligand. Therefore, when H2 O and DMF molecules are removed from the 1-D channels via heating or vacuum drying, a new crystalline compound with larger water content can be formed upon rehydration. A proton conductivity of 1.6 × 10−3 S/cm was observed at a temperature of 292 K and 100% RH. The high proton conductivity can be traced back to its structure as CUMTHF has 1-D channels that are occupied by water molecules running along the c and b axes. Furthermore, the phosphonates that are directed toward the channels might also aid in proton conduction. The activation energy was 0.31 eV, which points to the Grotthuss mechanism via the water molecules. More recently, two variants of a zirconium MOF using two different cations Na+ and NH4 + , Zr(O3 PCH2 )2 NC6 H10 -N(O3 CH2 P)2 Na2 H2 ·5H2 O (Na@Zr) and Zr(O3 PCH2 )2 NC6 H10 -N(O3 CH2 P)2 (NH4 )2 H2 ·5H2 O (NH4 @Zr) [112], were synthesized and studied for their proton conducting ability. Both Na@Zr and NH4 @Zr have a 3-D open-framework structure made of inorganic polymeric units bridged by cyclohexyl groups in a ‘‘brickwall-like’’ building texture with channels that are rectangular in shape. These channels are occupied by five water molecules and two Na+ and two NH4 + ions per formula unit, respectively. The water molecules form strong hydrogen bonds with each other and have strong interaction with the P O groups on the surface of the framework too. The Na+ and NH4 + ions in the MOFs can be substituted by protons. After substitution, proton conductivity of the MOFs was found to increase from 1.3 × 10−5 to 4.9 × 10−5 S/cm when RH was increased from 60 to 95%. This illustrated that water acts as the proton carrier via the Grotthuss mechanism, as evident from the low activation energy of 0.23 eV. In another study where water was involved in proton conduction, Banerjee et al. adopted a slightly different approach toward tuning the proton conduction of MOFs. They substituted the metal ions instead of modifying the organic linkers or the guest molecules [113]. Three 2-D MOFs, namely Ca-SBBA, Sr-SBBA, and Ba-SBBA (SBBA = 4,4 sulfobisbenzoic acid) were constructed using the alkaline earth metals Ca(II), Sr(II), and Ba(II) as metal nodes and SBBA as organic linkers. Surprisingly, their structures are completely different from each other although the metal nodes are all group II metals and the same organic linker was used. This could be due to the wide range of coordination numbers for the metal nodes, as well as different binding modes. The as-synthesized MOFs did not show any proton conductivity. Proton conduction only occurred after a period of 24 h humidification. The authors postulated that upon humidification, water molecules are absorbed within the crystal via the formation of hydrogen bonds with the carboxylate bound metal clusters, as well as with the electronegative sulfone groups. Ca-SBBA showed a proton conductivity of 8.6 × 10−6 S/cm, whereas Sr-SBBA exhibited a proton conductivity of 4.4 × 10−5 S/cm at 298 K under 98% RH. In contrast, Ba-SBBA did not show any proton conductivity under similar conditions. Judging from the low activation energy of 0.23 eV for Ca-SBBA, it was deduced that proton conduction follows the Grotthuss mechanism. On the other hand, the higher activation energy of 0.56 eV for Sr-BBA suggested the Grotthuss mechanism coupled with the vehicle mechanism for proton transport.

13 In accordance with the laws of kinetics, higher temperatures afford faster rates of reaction. However, the use of water molecules as proton conductors in MOFs limits the operating temperatures of most fuel cells to a maximum temperature of 100 ◦ C. As such, researchers have developed MOFs that do not rely on water to mediate their proton conduction. This was achieved by the incorporation of non-volatile carriers into the MOFs to enable proton conduction above the dehydration temperature and boost electrode kinetics. For example, Bureekaew et al. succeeded in trapping imidazole molecules within an aluminum MOF to create a hybridized polymer with high proton conductivity [114]. Imidazoles are promising proton carriers in fuel cells because they have higher boiling points and are much less volatile than water. Further, imidazoles exist as two tautomeric forms with a proton that moves between the two nitrogen atoms, creating a proton transport pathway. The authors suggested that the proton transport could occur by the Grotthuss mechanism, via structure diffusion that depends on proton transfer between the imidazole and imidazolium ion through their hydrogen-bonded chain. After this intermolecular proton transfer step, the molecules are then re-orientated according to step two of the Grotthuss mechanism. Two different MOFs, Al(␮2 -OH)(ndc) (ndc = 1,4naphthalenedicarboxylate) and Al(␮2 -OH)(bdc), both having 3-D framework and 1-D channels were synthesized and characterized. Both MOFs host imidazole molecules within their 1-D channels. The key difference between them lies in the difference in the ligands 1,4-ndc and 1,4-bdc. Due to this dissimilarity in ligands, the channels of the two MOFs have different characteristics, causing different host—guest interactions with the imidazole molecules, and therefore leading to different proton conductivities. Al(␮2 OH)(ndc) has two kinds of microchannels, both of which with hydrophobic pore surfaces, while Al(␮2 -OH)(bdc) has only one kind of microchannel that is amphiphilic in nature. The amphiphilic nature of the surface of Al(␮2 -OH)(bdc) leads to two different forms of interactions between the imidazole molecules and the surface — one strong and one weak. In comparison, the hydrophobic surface of the channels in Al(␮2 -OH)(ndc) can only have weak interactions with the imidazole molecules. The proton conductivity of Al(␮2 -OH)(ndc) after incorporation of imidazole was only 2.2 × 10−5 S/cm at 120 ◦ C with an activation energy of 0.6 eV. This is much higher than that of Al(␮2 -OH)(bdc), whose proton conductivity after incorporation of imidazole is 1.0 × 10−7 S/cm with an activation energy of 0.9 eV at the same temperature. This difference is attributed to the divergence between the dynamic motion of the imidazole molecules: in Al(␮2 -OH)(ndc), the non-polar surface of the pores do not interact strongly with the polar imidazole molecules. Hence, the imidazole molecules are able to move freely and conduct protons. However, in Al(␮2 -OH)(bdc), due to the amphiphilic nature of the surface, half the amount of imidazole interacts strongly with the hydrophilic sites on the surface of the pores, leading to enhanced host—guest interaction and restricting the mobility of imidazole molecules. 1-Ethyl-3-methyl imidazolium (Emim+) was also tested as proton carrier in MOFs. Long et al. succeeded in incorporating Emim+ into Co2 Na(bptc)2 (bptc = 2,2 ,4,4 -biphenyl tetracarboxylate) [115]. The resulting MOF had a phase transition temperature of 343 K. Three types of channels exist

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14 in the resulting MOF: one along the c-axis, one along the [1 1 0], and one along the [−1 1 0] direction. The channels along the [1 1 0] and [−1 1 0] direction have the same microenvironment and both of them are occupied by two Emim+ cations each. The channel along the c-axis, however, is only occupied by one Emim+ cation. The proton conductivity along the [110] direction was 2.63 × 10−5 S/cm while the proton conductivity along the c-axis was 4.78 × 10−7 S/cm. The higher proton conductivity along the [1 1 0] direction is attributed to the greater number of Emim+ cations in their channels, as ionic conductivity is proportional to the number of carriers. The proton conductivity was found to decrease with increase in temperature up to 343 K, because the increase in temperature leads to an increase in the disorder of the Emim+ cations, which leads to a decrease in the concentration of Emim+ cation charge in the channel. However, further increase in temperature beyond 343 K increases the proton conductivity, because the orientation order, motion, mobility, and concentration of Emim+ cations would increase beyond the phase transition temperature. These observations illustrates the relationships between temperature, order/disorder of the Emim+ in the channels, and proton conductivity — relationships that could be useful for further research to improve proton conductivity. In a following report, much better proton conductivity was obtained using histamine instead of imidazole as the guest molecule [116]. Histamine could be an excellent proton carrier as it has three proton-hopping sites — two on the imidazole ring and one on the amine group. It was found that histamine can be incorporated into Al(OH)(ndc), to produce a fine powder while retaining the integrity of the porous framework of Al(OH)(ndc). Up to 30% by weight of histamine can be loaded to Al(OH)(ndc), corresponding to one histamine molecule per Al3+ ion. The proton conductivity of Al(OH)(ndc) is negligibly low and that of pristine histamine is low too. However, the histamine-loaded Al(OH)(ndc) had a proton conductivity as high as 3.0 × 10−3 S/cm. The drastic increase of proton conductivity, especially when compared to the above mentioned imidazole-loaded MOF, can be explained by the much higher concentration of histamine than that of imidazole in the MOFs — the concentration of histamine is twice as high as that of imidazole. Although histamine is bulkier that imidazole, the histamine molecules in Al(OH)(ndc) are more densely packed and therefore formed a potent ion transport pathway, thus implying that the concentration of the proton carrier in MOF is instrumental in producing decent proton conductivity. Furthermore, it is likely that the intramolecular proton exchange that occurs in histamine facilitates the reorientation of the histamine molecules, thereby accelerating the molecular reorientation process — the rate-determining step in the Grotthuss mechanism. Inspired by the well-known proton conductor Nafion, the potency of sulfonated MOFs as proton-conducting materials was also investigated [93]. A MOF, Na3 (thbts) (thbts = 2,4,6trihydroxy-1,3,5-benzene trisulfonate) that has channels that are originally loaded with water was modified by replacing the water molecules with 1H-1,2,4-triazole (Tz) so as to allow proton conduction under an anhydrous H2 atmosphere above the dehydration temperature of 100 ◦ C. The proton conductivity measured for the Tz-loaded MOF was in the range of 2—5 × 10−4 S/cm at 150 ◦ C. Unlike Nafion and some other reported proton-conducting MOFs that utilize water

Y. Ren et al. as a proton carrier and require low operating temperatures and high RH, the Tz-loaded MOF has the advantage of being able to function above 100 ◦ C in the absence of H2 O. The viability of utilizing the Tz-loaded MOFs as the proton conducting membranes in a hydrogen-air fuel cell was further demonstrated. An open cell potential of 1.18 V was obtained and the cell was stable for 72 h at 100 ◦ C. As a potential non-aqueous proton conductor, the very first MOF HKUST-1 was thoroughly studied in methanol by Jeong et al. [117]. As shown in Fig. 10, HKUST-1 consists of Cu(II) paddlewheel-type metal nodes with BTC ligands. The Cu(II) sites are easily accessible and therefore they are the sites at which solvent or other molecules can be incorporated. The proton conductivities of as-synthesized HKUST-1, water coordinated HKUST-1, ethanol coordinated HKUST-1, and acetonitrile coordinated HKUST-1 were investigated (Fig. 10). Compared to pristine HKUST-1, the water coordinated HKUST-1 showed improved proton conductivity, while the other two had decreased conductivity. The lack of improvement in conductivity for acetonitrile coordinated HKUST-1 is understandable because acetonitrile has no dissociable proton. The lack of enhancement in proton conductivity for ethanol coordinated HKUST-1, which is more unexpected because it contains dissociable protons, can be explained by the lower acidity of ethanol. Since ethanol is a weaker acid and a poorer proton donor than methanol, it is unlikely for ethanol to donate protons to the stronger acid methanol. Based on their results, Jeong and co-workers suggested the approach of coordinating a poor proton donor to an open metal site should be possible in other MOFs with open metal sites, opening up the possibility of synthesizing more, and perhaps better proton conducting MOFs.

MOFs as H2 storage medium Another area of research relevant to fuel cell technologies is H2 storage. Currently, high-pressure compression and liquefaction are the two technologies used in commercial H2 storage. The large amount of energy consumed during compression and liquefaction, the continuous boil-off of liquid hydrogen, and the safety concerns related to compressed H2 limit the utility of these two H2 storage technologies. Researchers have therefore turned their attention to solid storage materials [118]. New materials capable of storing H2 at high gravimetric and volumetric densities are required if H2 is to be widely employed as a clean source of energy. The storage of hydrogen in solid materials can generally be divided into two categories: (1) chemical storage — hydrogen is stored in chemical compound and release of hydrogen is through a chemical reaction and (2) physisorption/entrapment — H2 is either adsorbed onto the surface of a solid material or captured by cages. In the latter, H2 storage capacity is largely dependent on the surface area and pore volume of a porous material. With exceptionally high surface areas, highly uniformed pore densities and dimensions, diversified and highly tunable chemical properties and structures, MOFs have recently emerged as one group of the most promising candidate materials for H2 storage. In view of the explosive research activities and several excellent reviews published recently on this topic [119—127], only representative examples are mentioned in this section to

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Figure 10 Qualitative representation of proton transfer from Cu(II) centers in HKUST-1 coordinated with (a) water, (b) ethanol and (c) acetonitrile. Reproduced with permission from Ref. [117]. © 2011 American Chemical Society.

illustrate the progress and prospective of the field, the relationships between structural features and the enthalpy of H2 adsorption, and strategies for improving storage capacity. In 2003, Rosi et al. first explored the cubic carboxylatebased framework MOF-5 (Zn4 O(BDC)3 ) as a H2 storage material [128]. It turned out that MOF-5 is the best cryogenic storage material currently known. It was also demonstrated that H2 can be loaded into a cold sample of the compound within 2 min, and can be completely desorbed and re-adsorbed for at least 24 cycles without loss of capacity [129]. Encouraged by the very promising results, hundreds of other MOFs have since been tested and numerous computational studies that attempt to guide the design of MOFs for H2 storage and model H2 adsorption data in MOFs have been carried out [127]. As the uptake of H2 is closely associated to the surface area of a material in physisorption, surface area, pore size, and pore volume of MOFs are the first parameters been extensively investigated. It is well known that physisorption of gases is both pressure and temperature dependent. In the case of MOFs, it seems that H2 uptake is closely associated with the heat of adsorption at low pressures, surface area at moderate pressures, and free volume at high pressures [125,126]. A systematic approach to study the effect of pore size and geometry on the surface area and H2 uptake is to design MOFs with the same topology and different organic ligand length. For example, a series of NbO-based MOFs were

constructed from tetracarboxylate ligands. By extending the ligand length, the window size and surface area (BET surface area) are increased accordingly. As a result, there was a slight decrease in H2 uptake at low pressures but a dramatic increase at 20 bar [130]. Introduction of open metal sites on the surfaces is an effective means of improving the H2 adsorption enthalpy in MOFs. It is well established that H2 can bind to certain metals with metal—H2 bond dissociation energies as high as 80—90 kJ/mol [131]. Unfortunately, this sort of metal—H2 interaction is obviously too strong for H2 storage since a great amount of heat would be released upon loading with H2 and a great amount of heat is required to liberate the bound H2 . To achieve the desired binding energy of ∼20 kJ/mol, one should look into the charge-induced dipole interactions for H2 storage instead of the orbital interactions. Therefore, high concentrations of exposed highly-charged metal cations must be first generated on the surfaces of MOFs. It has been demonstrated that thermally-assisted evacuation of solvent molecules bound to metal nodes is a feasible method of producing metal ions on the surfaces of MOFs [132,133]. In addition, reactive ligands in MOFs can also enable metalation, whereby metal ions are brought onto the surface of the MOFs [134]. PSM of MOFs with alkali metal represents an attractive route to significantly enhance the H2 storage capacity of MOFs [135—137] as simulations have suggested that atomic or ionic lithium modified MOFs may

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Y. Ren et al.

Figure 11 (a) TEM bright field and (b) STEM dark field images of the PdNP decorated MIL-100(Al). Reproduced with permission from Ref. [148]. © 2011 American Chemical Society.

result in MOFs with H2 storage capacity surpass the United States Department of Energy (DOE) targets [138,139]. For example, by attaching a lithium alkoxide functional group on the ligand of IRMOF-8 (Zn4 O(ndc)3 ), a combination of quantum and classical calculation showed that the modified IRMOF-8 should have a gravimetric uptake of 4.5 wt% at 295 K and 100 bar and 10 wt% at 77 K and 100 bar [139]. Also, comparing to the unmodified hydroxyl-MIL-53(Al), lithium alkoxide functionalized hydroxyl-MIL-53(Al) showed a significant increase in H2 uptake capacity from 0.50 to 1.7 wt% at 77 K and doubled heat of adsorption at low coverage (5.8 versus 11.6 kJ/mol) [139]. However, care must be taken in controlling the degree of modification as over-modification with lithium ions may have detrimental effect on the H2 storage capacity [140]. MOFs with large pores tend to have low volumetric H2 uptake capacities under ambient conditions, as hydrogen molecules near the center of the pore experience little interaction with the pore walls. A large pore volume that is composed of a large number of small voids is more desirable for an efficient storage material [141]. One could possibly reduce the number of large voids in MOFs via interpenetration/catenation. For example, a two-fold interpenetrating MOF exhibited improved H2 uptake, providing some evidence for the advantages of interpenetrated MOFs [142]. On the other hand, a recent simulation study by Ryan et al. suggested that interpenetration may be beneficial only at low loadings because of the overall reduced pore size and higher number of metal sites per unit volume. At intermediate and high loadings, the surface area and total free volume become more important [143]. Hydrogen spillover could also be utilized in enhancing the H2 storage capacity. Hydrogen spillover involves the dissociation of H2 into H• on a metal surface and the subsequent migration of these atoms onto the support surface [144—146]. This reaction is reversible with hydrogen atoms spontaneously recombining to afford the diatomic gas. Hydrogen spillover can be realized by incorporating metal nanoparticles in MOFs. Indeed, Zlotea demonstrated that a nearly two-fold enhancement in H2 uptake at 298 K is achieved when MIL-100 (Al) was embedded with 2-nm palladium nanoparticles (PdNPs) (Fig. 11). The much improved

H2 uptake has been partially attributed to ␤-hydride formation at low pressures and a spillover mechanism at pressures above 4.5 kPa [147]. PdNP-embedded SNU-3 (Zn3 (ntb)2 , ntb = 4,4 ,4 -nitrilotrisbenzoate) (PdNP@SNU) also showed an increased H2 storage capacity of 1.48 wt% at 77 K, compared to the 1.03 wt% of pristine SNU-3, due to the effect of hydrogen spillover [147]. While the employment of a hydrogen spillover mechanism in MOFs has generated some very encouraging results for room-temperature H2 storage, many fundamental questions remain to be addressed. For instance, the possibility of achieving a spillover effect with non-noble metals such as cobalt or nickel and with individual metal ions also remain to be realized. Enhanced H2 storage capability can also be achieved through rational ligand design. Panda et al. recently synthesized a ZIF-type MOF with both a free amino group and a free tetrazole nitrogen [148]. A significantly higher H2 uptake of 1.6 wt% more than other members of the ZIF series at 77 K and 1 bar was observed, despite the comparatively lower surface area. In a related study, Cohen and co-workers have postsynthetically modified three different MOFs via condensation reactions between free amino groups in the MOFs and alkyl anhydrides and isocyanates [149]. The modified MOFs exhibited higher H2 uptake than their unmodified counterparts, even at higher pressures, demonstrating that aromatic substituents enhanced the H2 uptake because of the additional binding site. Hupp et al. also postsynthetically modified the Zn nodes of a MOF by coordinating a series of functionalized pyridine ligands [47]. It was shown that the PSM can substantially modulate the MOF’s internal surface area, pore volume, and ability to interact with H2 , providing a promising route to MOF property modulation in H2 storage.

Concluding remarks and future perspectives This review has illustrated the vast possibilities MOFs provide toward the enhancement of fuel cell technologies. The proof-of-concept studies reviewed above will certainly stimulate more research in this exciting area. MOFs are able to contribute to the advent of the hydrogen economy via applications in hydrogen production, oxygen reduction

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Metal-organic frameworks in fuel cells electrocatalysts, PEMs, and H2 storage. In this aspect, their exceptionally high surface areas (up to 5.0 × 103 m2 g−1 ), highly uniformed pores, tunable pore geometries, versatile structures, and rich chemistry render MOFs attractive candidates for fuel cell technologies. Moreover, the utilization of multifunctional bridging ligands and postsynthetic modifications has allowed the judicious design and functionalization of MOFs, allowing their properties to be rationally tuned. Flexible control over the pore structure, combined with surface engineering, may open up new avenues for systematically studying the effect of structural parameters and chemical compositions on the performance of the MOF-based fuel cells, thus leading to significant fundamental insights. In the context of their application to visible-lightpromoted photocatalytic hydrogen production, MOFs possess great flexibility in terms of framework design, due to the simultaneous utilization of the unique characteristics of both the metal nodes and organic linkers. In principle, it should be possible to integrate a photosensitizer and a water reduction catalyst into a single MOF system. For example, it has been observed that electron transfer can take place from the photoexcited organic linkers to the metal-clusters within the Zr-based MOFs [76]. This process, termed linker-to-cluster charge-transfer (LCCT), occurs in photocatalytic reactions and during photoluminescence, in which MOFs have been known to participate in [8,76]. In hydrogen production via the LCCT mechanism, the conduction band (CB) edge position of a titanium-oxo cluster would be more suitable for efficient charge transfer from the excited state of the organic linker, owing to the fact that the CB potential of a titanium-oxo cluster is more positive than that of its zirconium counterpart. Hence, Ti-based MOFs are expected to be excellent photocatalysts for hydrogen generation from water. It has also been reported that porphyrin-based photocatalytic systems show very impressive H2 generation from water when irradiated at 320 nm [75]. It is likely that incorporation of porphyrins into MOFs as photosensitizers will lead to an improvement in catalytic efficiency, and this could motivate further research into porphyrin-based MOF photocatalysts. However, these types of photosensitizers are typically active only in the ultraviolet range, which accounts for merely 3—5% of the solar spectrum. Therefore, the development of MOF-based photocatalysts for the water splitting reaction that responds to visible light, the major component of sun light, is fundamentally important. The most remarkable characteristic of MOFs relevant for catalytic production of H2 from on-board fuels is the full accessibility of catalytic centers to the reactants. Owing to the remarkably open architecture and extremely high porosity of MOFs, the diffusion coefficients of molecules in the pore system are only slightly lower than that in the bulk solvent. This means that mass transport in the pore system is not hindered. Additionally, the ordered structure offers the opportunity to spatially separate active centers to maximize their exposure. Lastly, as a result of their high surface areas, MOF-based catalysts contain a very high density of fully exposed active sites per volume. These features of MOFs are expected to result in enhanced catalytic activity and an overall more effective catalytic system. MOFs are also excellent matrices for studying the nanoconfinement of metal catalysts for the dehydrogenation of on-board fuels,

17 and they have potential applications in unexplored areas, as hinted by their ability to oxidize unorthodox substrates such as organosilanes to produce hydrogen. As such, MOFs are expected to help us reach the targets and benchmarks set by DOE and accelerate the practical implementation of PEMFCs in our daily lives in the near future. Comparatively, there has been relatively less research into the usability of MOFs as ORR electrocatalysts in fuel cells. As the kinetics of oxygen reduction is most often the limiting step in a fuel cell, more work definitely needs to be done in this area in order for MOFs to really contribute as revolutionary materials in accelerating the practicability of fuel cells. In-depth theoretical and experimental studies are needed to delicately design MOF-PGM electrocatalysts with desirable pore sizes and dimensions for specific applications. It is also essential to clarify the role of the MOF skeleton in the overall catalytic activity. The future trend is leaning toward improving the utilization of PGM or replacing them completely with less expensive metals such as Co, Fe, or Ni, for cost effectiveness. It is still a pioneering area to explore a facile approach to fabricate robust MOFbased electrocatalysts with less or no PGM content but yet providing good catalytic activity, high poisoning resistance, and long-term stability. In addition, the unusually large cavity sizes, now encroaching into the mesoporous regime of a few nanometers, suggest that MOFs may contribute to the bottom-up approach of synthesizing nanoscale electrocatalysts [150,151]. With the continuous advancement of preparation methodologies and characterization techniques for new MOF-based electrocatalysts, it is expected that revolutionary breakthroughs and wider technological applications are realizable in the near future. There is growing evidence to support the optimistic forecasts in the application of MOFs as PEM in fuel cells. Different research groups have employed different tactics to boost proton conduction, such as modifying the hydrophilicity of the framework and coordinating molecules that assist in proton conduction to the frameworks, leveraging on both the Grotthuss mechanism and vehicle mechanism. The use of non-volatile proton carriers instead of water has also been explored. These different and unique strategies form a repository that other researchers can tap into or even add to. Impressive improvements in proton conductivity and practical applicability have been obtained over the years. For example, MOFs with impressive proton conductivities rivaling that of Nafion have been successfully synthesized. Further improvements are expected through rational design of MOFs that can accommodate both the Grotthuss mechanism and the vehicle mechanism. On the other hand, a critical issue to be addressed urgently in the next stage of research is the long-term operating stability of MOFbased PEMs in working fuel cells. A few days of stable operation of the MOF-based PEMs falls far short of any meaningful applications in fuel cells. Unlike the high stability of zeolites due to the high energy barrier needed to break the strong Si O and Al O bonds, coordination bond strengths in MOFs are lower and thus MOFs are expected to be less thermally stable. MOF stability could be maximized by ensuring good hard—hard or soft—soft matchings between the ligand and metal to maximize bond strengths. This is strongly suggested by the fact that some of the highest decomposition temperatures for MOFs have been observed

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18 for lanthanide carboxylates, which possess good soft—soft matching [17]. Additional stability could be conferred upon MOFs by introducing covalent bonds within the MOFs with minimal implications on their proton conductivity. Another possible solution could be derived from a combination of strategies, such those in MOF-based hybrid blends. To reach these goals for proton conductive MOFs, new strategies and materials must ultimately be developed, and material scientists as well as fuel cell engineers must work side-by-side to determine which strategy works best and which materials function best under fuel cell operating conditions. With exceptionally high surface areas, highly uniformed pores, and tunable pore geometries, MOFs are a natural choice for gas storage. H2 storage in particular is important for fuel cell applications. Significant progress has been made since the first attempt to utilize MOFs as a H2 storage medium in 2003. Outstanding H2 storage capacities at cryogenic temperatures of ≤77 K and at high pressures of up to 100 bar have been realized. However, the H2 storage capacity decreases dramatically under ambient conditions. In order to apply MOFs as H2 storage media in fuel cells, MOFs should be capable of storing large amounts of H2 at ambient temperature. The DOE targets that should be achieved by 2017 are 5.5 wt% and 45 g/L of H2 at 40—60 ◦ C under the maximum delivery pressure of 100 atm [152]. In order to effectuate these targets, MOFs should ideally possess a high H2 adsorption enthalpy. Theoretical calculations have predicted that MOFs should have a 15—25 kJ/mol of isosteric heat during H2 adsorption in order to store H2 at ∼30 bar and to release it at ∼1.5 bar. However, most MOFs reported so far have isosteric heat of H2 adsorption in the range of 5—12 kJ/mol [153]. Research is currently being steered toward the synthesis of MOFs that have enhanced interaction energies with H2 . To increase the interaction energy between H2 and MOFs, various strategies, such as constraining pore size, generating open metal sites, ligand functionalization, interpenetration, alkali or alkaline-earth metal ion inclusion, doping with metal ions, and embedding MOFs with hydrogen absorbing metal nanoparticles, have been developed. Constraining pore size may be beneficial in improving the isosteric heat of H2 adsorption. However, a delicate balance must be attained as surface area generally scales down together with decreasing pore size. Therefore, further investigations should be conducted for the optimization of pore size and surface area. Since creating open metal sites definitely increases the isosteric heat of H2 adsorption, one may also design new MOFs using multifunctional organic linkers that contain open metal sites after complexing with metal ions. However, generating a high concentration of open metal sites in a MOF and achieving the ideal case whereby each metal ion binds to more than one H2 molecule remains challenging. Unexpectedly, the incorporation of small metal ions with high charge density in MOFs could only slightly increase the isosteric heat of H2 adsorption due to the very strong interaction between the metal ions and solvent molecules. Theoretical calculations have shown that the binding energies to H2 are 21.9, 34.6 and 46.5 kJ/mol for Sc, Ti and V, respectively [154]. Therefore, constructing MOFs from these metals and subsequently generating open metal sites in the resulting MOFs may be beneficial for enhancing H2 storage capacity at room temperature. When designing MOFs, one must keep in mind

Y. Ren et al. that strong metal—ligand bonds should be made so that the MOFs are stable under the working conditions of fuel cells. It is anticipated that research in H2 storage using MOFs will continue to be vigorous. We can expect that with an adequately strong initial heat of adsorption, sufficiently large surface area as well as pore volume, and a multitude of MOF—H2 interaction mechanisms, MOFs will eventually be able to reach the DOE targets for an on-board H2 storage system. Being highly versatile, MOFs have the propensity to be rationally modified according to the different needs that may arise. Novel applications of MOFs in niche areas such as electrocatalysis of ethanol have been reported [155], exemplifying the versatility of MOFs in the most unexpected of areas. It is heartening to witness the diversion of research interest in MOFs away from traditional fields such as storage or separation, and into more urgent and pressing needs such as clean energy production. Much progress has been made with respect to ramping up the performance of MOFbased components to industrial and commercial standards. However, further breakthroughs will require MOFs to outperform the standards set by conventional materials. Given the vast number of possible permutations of metal centers and organic linkers, the outlook of the applications of MOFs in fuel cells remains promising.

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Metal-organic frameworks in fuel cells Ms. Yuqian Ren received her B.Sc. in applied chemistry from Peking University in 2012. She is currently a graduate student in the Department of Chemistry National University of Singapore. Her research interests include applications of metal—organic frameworks and electrocatalysts for water splitting and fuel cell applications.

Mr. Guo Hui Chia was born in Singapore in 1989. He served as a platoon commander in the Singapore Armed Forces before pursuing his Bachelor of Science (Honors) in Chemistry at the National University of Singapore. He is now working on his honors project at the Department of Chemistry.

21 Dr. Zhiqiang Gao is an associate professor at the Department of Chemistry National University of Singapore. He received his B.Sc. and Ph.D. in Chemistry from Wuhan University. The following years he worked as a postdoctoral fellow at Åbo Akademi University and The Weizmann Institute of Science. After spending three years in the United States and seven years at the Institute of Bioengineering and Nanotechnology, he joined NUS in April 2011. Research in his laboratory currently includes electrochemistry, analytical chemistry, and materials science.

Please cite this article in press as: Y. Ren, et al., Nano Today (2013), http://dx.doi.org/10.1016/j.nantod.2013.11.004