Nanostructured Metal Hydrides for Hydrogen Storage

Review Cite This: Chem. Rev. XXXX, XXX, XXX−XXX

pubs.acs.org/CR

Nanostructured Metal Hydrides for Hydrogen Storage Andreas Schneemann,†,⊥ James L. White,†,⊥ ShinYoung Kang,‡ Sohee Jeong,§ Liwen F. Wan,‡ Eun Seon Cho,§,∥ Tae Wook Heo,‡ David Prendergast,§ Jeffrey J. Urban,§ Brandon C. Wood,‡ Mark D. Allendorf,† and Vitalie Stavila*,† †

Sandia National Laboratories, Livermore, California 94551, United States Lawrence Livermore National Laboratory, Livermore, California 94550, United States § Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States ∥ Department of Chemical and Biomolecular Engineering, Korea Advanced Institute of Science and Technology (KAIST), Daejeon 34141, Republic of Korea

Chem. Rev. Downloaded from pubs.acs.org by UNIV OF LOUISIANA AT LAFAYETTE on 10/02/18. For personal use only.

‡

ABSTRACT: Knowledge and foundational understanding of phenomena associated with the behavior of materials at the nanoscale is one of the key scientific challenges toward a sustainable energy future. Size reduction from bulk to the nanoscale leads to a variety of exciting and anomalous phenomena due to enhanced surface-to-volume ratio, reduced transport length, and tunable nanointerfaces. Nanostructured metal hydrides are an important class of materials with significant potential for energy storage applications. Hydrogen storage in nanoscale metal hydrides has been recognized as a potentially transformative technology, and the field is now growing steadily due to the ability to tune the material properties more independently and drastically compared to those of their bulk counterparts. The numerous advantages of nanostructured metal hydrides compared to bulk include improved reversibility, altered heats of hydrogen absorption/desorption, nanointerfacial reaction pathways with faster rates, and new surface states capable of activating chemical bonds. This review aims to summarize the progress to date in the area of nanostructured metal hydrides and intends to understand and explain the underpinnings of the innovative concepts and strategies developed over the past decade to tune the thermodynamics and kinetics of hydrogen storage reactions. These recent achievements have the potential to propel further the prospects of tuning the hydride properties at nanoscale, with several promising directions and strategies that could lead to the next generation of solid-state materials for hydrogen storage applications.

CONTENTS 1. Introduction 2. Classes of Nanostructured Metal Hydrides 2.1. Bonding in Metal Hydrides 2.2. Nanostructured Binary Hydrides 2.3. Nanostructured Intermetallic Hydrides 2.4. Nanostructured Complex Metal Hydrides 3. Synthetic Routes 3.1. Ball Milling 3.2. Gas-Phase Synthesis 3.2.1. Gas-Phase Condensation 3.2.2. Plasma Deposition 3.2.3. Thin Film Syntheses 3.3. Reductive Methods 3.3.1. Chemical Reduction 3.3.2. Electrochemical Reduction 3.3.3. Thermal Decomposition 3.4. Nanoconfinement 3.4.1. Host Materials 3.4.2. Impregnation Methods 3.4.3. Synthesis Inside the Pores 4. Methods to Analyze Structure and Storage Properties © XXXX American Chemical Society

4.1. Experimental Structural Analyses 4.1.1. Transmission Electron Microscopy 4.1.2. X-ray and Neutron Techniques 4.1.3. Vibrational Spectroscopy 4.2. Theoretical Structural Determination 4.2.1. Wulff Construction 4.2.2. Prototype Electrostatic Ground State Method 4.2.3. Microstructure Modeling 4.3. Thermodynamic and Kinetic Analysis 4.3.1. Thermodynamic Measurements 4.3.2. Thin Film Hydrogen Content Measurements 4.3.3. Rate Analysis and Kinetic Models 4.3.4. Thermodynamic Models 5. Effects of Morphology on Hydrogen Storage Properties 5.1. Free-Standing Nanoparticles 5.2. Thin Films 5.3. Nanoconfined Metal Hydrides

B E E F G H I I J K K L L L M M M N N O

O O O P P P P Q Q Q R S T U U W Z

Received: May 17, 2018

O A

DOI: 10.1021/acs.chemrev.8b00313 Chem. Rev. XXXX, XXX, XXX−XXX

Chemical Reviews 5.3.1. Confinement in Carbons 5.3.2. Other Rigid Porous Hosts 6. Mechanistic Effects of Nanosizing 6.1. Effects of Surfaces 6.1.1. Effects on Reaction Enthalpy 6.1.2. Effects on Reaction Pathways 6.1.3. Surface Entropy and Disorder 6.1.4. Chemical Kinetics at Defective Surfaces 6.1.5. Chemical Impurities and Oxidation 6.2. Effects of Internal Interfaces and Microstructure 6.2.1. Effects of Grain Boundaries 6.2.2. Types of Interphase Boundaries 6.2.3. Effects of Reactive Phase Boundaries 6.2.4. Effects of Nonreactive Phase Boundaries 6.2.5. Nucleation and Phase Transformation Mechanisms 6.3. Other Size-Related Effects 6.3.1. Shortened Diffusion Pathways 6.3.2. Altered Phase Coexistence Behavior 6.4. Metal Hydride-Host Interactions 6.4.1. Physical Confinement Effects 6.4.2. Electronic Coupling Effects 6.4.3. Chemical Interactions with the Host 7. Conclusions Author Information Corresponding Author ORCID Author Contributions Notes Biographies Acknowledgments List of Abbreviations References

Review

world energy problems by delivering and storing energy using hydrogen. The Hydrogen Economy offers a potential solution to satisfying global energy needs while reducing (with the ultimate goal of completely eliminating) carbon dioxide and other greenhouse gas emissions and improving energy security. Hydrogen storage is an essential component of a viable hydrogen economy.4 Whether the application is stationary or geared for the transportation sector, volumetric and gravimetric efficiencies often dictate feasibility. The mobile vehicular application is one of the most challenging applications, as it requires materials with very large gravimetric and volumetric hydrogen capacities and also mandates other stringent properties such as extended cycle-life, low impurity levels, and fast kinetics of hydrogen uptake and release. Hydrogen-powered fuel cell vehicles (HFCVs) are now commercially available and store the fuel in fiber-reinforced tanks at 700 bar. Although sufficient to achieve reasonable driving range (∼350−450 km), they still fall short of the ultimate U.S. Department of Energy (DOE) capacity targets for on-board storage for light-duty vehicles (0.065 kg H2/kg system and 0.050 kg/L, Table 1). The resulting overall fuel

Z AF AG AK AK AL AL AM AM AN AO AO AQ AQ AR AT AT AT AT AU AU AU AV AW AW AW AW AW AW AX AY AY

Table 1. Summary on the Targets for Hydrogen Storage Systems Set by the United States Department of Energy for on-Board Vehicular Applications7 storage parameter

units

2020

2025

ultimate

system gravimetric capacity: usable, specific-energy from H2 system volumetric capacity: usable energy density from H2 storage system cost

kWh kg−1 (wt %)

1.5 (4.5)

1.8 (5.5)

2.5 (6.5)

kWh L−1 (kg H2 L−1)

1.0 (0.03)

1.3 (0.04)

1.7 (0.05)

$ kWh−1 net ($ kg−1 H2) °C

10 (333)

9 (300)

8 (266)

−40/85

−40/85

−40/85

cycles

1500

1500

1500

bar (abs)

5

5

5

cycles

1500

1500

1500

min % H2

3−5 99.97%

3−5 99.97%

3−5 99.97%

min/max delivery temperature cycle-life (uptake/ release cycles) minimum delivery pressure cycle life (1/4 tank to full) system fill time fuel purity (H2 from storage)

1. INTRODUCTION The ever-growing fossil fuel-based economy has led to a number of new challenges facing our civilization in the 21st century, such as pollution and irreversible climate change due to excessive drilling, mining, and growing amounts of greenhouse gases. The dominance of fossil fuels as prevailing energy sources is clearly unsustainable in the long run because the adverse impact of a single energy source is additive and repetitive, and once the harmful consequences have accumulated beyond a critical threshold, permanent damage ensues.1,2 Diversification of energy sources and the development of alternatives to fossil fuels such as hydrogen-powered and electric vehicles are not only highly desirable, but some of them are currently technologically feasible and cost-competitive. Advanced energy materials with improved properties are a prerequisite for most technological innovations and are mandatory to meet the increasing demands of a growing global population. As a consequence, it becomes apparent that alternative energy solutions and materials are critical for a sustainable future.1 Hydrogen is considered an attractive alternative fuel, or more precisely an alternative energy carrier, particularly when used in hydrogen polymer electrolyte membrane (PEM) fuel cells.3 The concept of an ecologically clean “Hydrogen Economy” was first introduced in the mid-1970s and has been gaining momentum as a viable remedy for the growing

cost is a significant additional concern, exceeding the target due to costs of the tank, compressor, and associated fueling station hardware. The ultimate hydrogen storage material should have a high gravimetric capacity (≥10 wt % hydrogen), have high reversibility (≥1500 cycles), and release hydrogen below 85 °C, which would allow the utilization of waste heat from a PEM fuel cell for desorption. The storage medium must also be “kinetically fast,” meaning that the material is capable of releasing hydrogen at the demanded rate (up to 2 g H2 s−1) and also can uptake hydrogen fast enough so the tank can be refueled with hydrogen in the desired time (∼15 g H2 s−1). Consequently, on-board storage remains a factor limiting the widespread adoption of HFCVs. Metal hydrides, which bind hydrogen chemically, and sorbents, which store H2 through physisorption, comprised of light elements could, in principle, meet the DOE storage targets, yet each of these two classes of materials has its own advantages and drawbacks. For instance, materials such as 170 Mg(BH4)2 (14.9 wt %) and LiBH4 (18.5 B


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Figure 1. Processes accompanying hydrogen storage reactions in bulk (a) and nanostructured (c) metal hydrides. Despite differences in chemical properties, the local bonding (b) can be similar.

Figure 2. Examples of nanostructured metal hydrides, including NPs, core−shell structures, nanowires, nanotubes, thin films, and multilayers: (a) palladium NPs (adapted with permission from ref 25. Copyright 2009 Elsevier), (b) Mg@MgO core−shell NPs (adapted with permission from ref 26. Copyright 2014 Elsevier), (c) magnesium nanowires (adapted from ref 23. Copyright 2007 American Chemical Society), (d) palladium nanotubes (adapted with permission from ref 27. Copyright 2000 Elsevier), (e) Pd-capped Mg thin film (adapted with permission from ref 28. Copyright 2008 Elsevier), and (f) Mg/Cr multilayer film (adapted with permission from ref 29. Copyright 2014 Elsevier).

wt %) nominally can meet both the gravimetric and volumetric targets.5 Unfortunately, other factors limit successful use of hydrides as storage materials, and none currently can simultaneously meet the DOE targets for minimum delivery pressure, charging/discharging rates, and capacity. Sorbents, in contrast, have rapid adsorption/desorption rates and are readily reversible, but their capacities are very low except at cryogenic temperatures, typically 77 K.6 The reasons for the inadequate performance of any given hydride are complex, but in general these can be broken down into thermodynamic and kinetic factors. Considering thermodynamics, hydrogen storage materials (including both sorbents and light metal hydrides) fall on either side of an ideal, but effectively empty, range of hydrogen binding energies, with sorbents binding H2 too weakly through physical adsorption and hydrides binding it too strongly with chemical bonds. Metal hydrides suffer additionally from slow H2 uptake and release kinetics and limited reversibility, due in the case of some complex hydrides to the intricate reaction pathways and myriad intermediates involved. Many strategies have been

considered for improving these properties with varying degrees of success.5,8,9 These include entropic effects,10 altering the morphology of hydride particles,11 mechanical strain,12 and dopants or other chemical additives.13 Although one cannot say that the potential of any of these has been exhaustively evaluated, the improvements achieved so far are modest. A promising approach that has garnered substantial attention in the past decade is nanostructuring of the metal hydrides. Many recent experiments and models that will be described in this review suggest that nanostructured metal hydrides have significantly, rather than incrementally, different thermodynamics and kinetics than their bulk counterparts. For example, a substantial decrease in the H2 desorption enthalpy is predicted for very small MgH2 particles.14 Similarly, force field modeling predicts a decrease in thermodynamic stability of MgH2 as particle size decreases from 2.0 to 0.6 nm.15 Barriers for H2 desorption from NaAlH4 are also predicted and experimentally found to depend on particle or cluster size.16−18 Nanoconfinement of Li3N was shown to fundamentally alter both the hydrogenation and dehydrogenation reaction pathC


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recesses into which the confined material is introduced, bound, and then restricted from movement or agglomeration with other particles. The pores, therefore, determine the shape and dimensionality of the nanoscale material. Nanoencapsulation is defined here as incorporation of a nanoscale material inside a secondary material which is not necessarily porous, named as the matrix or shell (e.g., graphene,30−33 reduced graphene oxide,34−36 and polymers22,37−39). Encapsulation can also occur in core−shell NPs26,40,41 and multilayered thin films29,42 and involves a preformed nanostructure that acts as a barrier to particle/grain growth and serves as protection against disruptive conditions that could cause agglomeration or pulverization upon cycling. Theory and modeling are playing a critical role in elucidating nanoscale phenomena associated with thermodynamics and kinetics of metal hydrides. The advances in computational power and the emergence of supercomputers enabled quantum mechanical, empirical, mesoscale, and continuum approaches to modeling various phenomena in metal hydrides.43 Classical molecular dynamics (MD) simulations are widely used to discover mechanisms and diffusion pathways, examining moderate-to-large sized systems (i.e., ∼103 to 106 atoms). Density functional theory (DFT) calculations are used to quantify energy barriers and end-state phase energies for various phases and reactions. Combining MD and DFT calculations enable direct simulations of time-evolution (picoto nanoseconds) to identify mobile species and energy barriers. Kinetic Monte Carlo simulations can cover much longer time scales (up to seconds), whereas continuum methods can, in principle, cover time scales from seconds to hours. Experimental confirmation of predictions is challenging for several reasons, although there is abundant evidence indicating that nanoscale metal hydrides have lower H2 desorption temperature than bulk hydrides.24,44−49 First, the sizes suggested by the calculations at which changes in properties occur are much smaller than those produced by mechanochemical approaches, which are fraught with problems, including the presence of impurities, lack of size control, and poor reproducibility. Alternatively, bottom-up approaches to nanoscaling, such as the formation of dendrimers, aerogels, and inorganic templates (e.g., zeolites), are unsuitable due either to the highly reactive nature of storage materials or the substantially detrimental impact on gravimetric capacity.44−49 Second, tuning particle size can sometimes be difficult. Pores in host materials are not always monodisperse, making it hard to achieve uniform thermodynamic behavior and identify the transition to nonbulk properties. For example, NaAlH4 particles supported on carbon nanofibers have broad size ranges (e.g., 19−30 nm), resulting in hydrogen desorption temperatures from 70 to 200 °C.50 Molten-metal infiltration of carbon aerogels, which have a distribution of pore sizes, yields 2−5 nm Mg particles,51 but wetting agents (e.g., Ni) degrade the scaffold, and blocking of the outermost pores appears to limit infiltration. Third, the chemical environment surrounding the particle may play an important role. For example, colloidal stabilization of Mg particles yields 5 nm diameter particles that desorb H2 as low as at 85 °C (>150 °C lower compared to bulk).52 In general, the factors leading to improved properties at the nanoscale are still poorly understood, however, and there are numerous challenges to be resolved or overcome. Free NPs are not stable, so some form of stabilization or hosting is required. Unfortunately, nanoconfinement adds dead mass and volume

ways as a direct consequence of solid−solid nanointerfaces within the material.19 These are just a few examples to show that there is ample evidence for altered physical and chemical properties of nanostructured metal hydrides compared to bulk (Figure 1). Nanostructured materials have sizes intermediate between molecules and microscopic structures. They can be a wide variety of shapes and sizes (Figure 2), but they share the characteristic of having at least one dimension be on the nanoscale, which is typically defined as being between 1 and 100 nm (1 nm = 10−9 m). Nanoparticles (NPs), which have three nanodimensions, are among the most common nanostructures in materials science due to their relative ease of synthesis, often through nucleation and restricted growth processes, display large fraction of surface atoms, which makes them attractive for numerous applications, particularly in catalysis and energy-related applications.20 Some NPs are composed of two different materials, either in a Janus structure, in which each covers a portion of the exterior of the particle, or a core−shell structure, with one of the materials completely encapsulating the other (typically designated as core material@ shell material).21,22 Hollow nanotubes and solid nanorods (also called nanowhiskers, nanowires, or nanopillars, often depending on the aspect ratio of length to width) are elongated in one dimension.23 Two-dimensional materials, which have only one nanoscale dimension, include planar and sometimes curved structures including nanosheets, nanoflakes, nanoribbons, and thin films.24 Some films that have been prepared possess alternating layers of two or more discrete materials; even though these multilayers may be grown thicker than 1 μm, the discrete layers typically are individually less than 100 nm and thus remain classified as nanostructured. A further distinction among 2D materials is that thin films are generally synthesized by deposition onto a flat (occasionally monocrystalline) substrate of another material, whereas nanosheets, nanoflakes, and nanoribbons are usually prepared independently of a synthetic foundation. Nanoconf inement (Figure 3) is a broad term that involves either the formation of a nanostructured material inside a host (nanoscaffolding) or the encapsulation or coating of a material with a rigid matrix (nanoencapsulation). We define nanoscaf folding as confinement of a material inside of a scaffold with permanent porosity. The host material has nanoscale pores or

Figure 3. Schematic depicting the two major subclasses of nanoconfinement discussed in this review: nanoscaffolding (left), with the hydride inside a host with permanent porosity, and nanoencapsulation (right) with the hydride wrapped by a nonporous material, such as graphene. D


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Metal hydrides in general, and nanostructured metal hydrides in particular, can be grouped into three large classes: (1) binary hydrides MHx (M= main-group or transition metal, such as in LiH, MgH2, PdH0.6, TiH2); (2) intermetallic hydrides, ABxHy (A is typically the hydriding metal and B is the nonhydriding metal, such as TiFeHx, TiMn2Hx, and LaNi5Hx); and (3) complex metal hydrides, MEHx (E = boron (borohydrides, e.g., LiBH4, Mg(BH4)2); nitrogen (amides, e.g., LiNH2, Mg(NH2)2), aluminum (alanates, e.g., NaAlH4, Mg(AlH4)2)). In binary ionic hydrides such as NaH and CaH2, hydrogen exists as the negatively charged hydride ion (H−). In contrast, interstitial hydrides, such as palladium hydride, are usually nonstoichiometric compounds, with hydrogen existing as atoms dissolved in the lattice of host metal atoms. Complex metal hydrides are compounds composed of metal cations and polyatomic hydrogencontaining anions such as alanates (AlH4−), borohydrides (BH4−), and amides (NH2−). These compounds possess hydrogen atoms covalently bound to aluminum, boron, or nitrogen, with the hydrogen atoms having either δ+ or δ− polarity.

that reduces storage capacity. From a synthetic aspect, some modeling suggests that nanoscale effects do not present themselves until particles are one nanometer in size or even smaller. Thin films lack capacity and the low-cost fabrication techniques required for large-scale production. Core−shell architectures may avoid this problem but are still very much at the research stage. In nanoconfined materials, it is unclear which is the critical factor: surface-to-volume ratio, particlehost interaction energy, or cluster geometry (“magic number” effects are observed when metals such as gold are nanoscaled,53 for example), though some scaffolds may also have catalytic effects independent of the size of the hydride.54 Carbon, probably the most economical nanoconfinement material, is difficult to characterize in terms of its pore chemistry, and thus it is problematic to control and synthesize reproducibly. Nevertheless, nanoscaling is a promising synthetic strategy that in recent years has attracted considerable attention.48,49,55−58 Several review articles with a narrower focus were published in recent years.44,46−48,59−62 Given the renewed interest in hydrogen storage in general and metal hydrides in particular (for storage as well as other applications), a more current and general review is warranted. In this article, we review new experimental and theoretical approaches to synthesizing nanostructured metal hydrides and probing their hydrogen storage properties. This review aims to describe, summarize, and analyze reports in the scientific literature on nanostructured metal hydrides, both simple and complex, and their use for hydrogen storage, including their capacities and their thermodynamic and kinetic properties in H2 uptake and release. Techniques for making nanostructured metal hydrides will be presented, as will computational methods for probing their properties. Previous review articles and book chapters on metal hydrides typically discuss nanostructuring only briefly and not as a comprehensive study, and some have focused specifically on a narrow class of nanostructured metal hydrides.8,24,44−46,55,56,58,60,61,63−69 In addition, the present review covers the recently discovered novel classes of nanostructured hydrogen storage materials, such as grapheneencapsulated metal hydrides and core−shell nanostructures, which present a major leap forward in improving both the thermodynamics and kinetics of metal hydrides for hydrogen storage applications.

2.1. Bonding in Metal Hydrides

Metal hydrides exhibit a great diversity of chemical bonds, spanning from metallic bonding, as in interstitial hydrides such as PdHx, ionic bonding in some binary and complex hydrides, and partial or full covalent bonding as in the lightest binaries (LiH and MgH2) and complex hydride anions (AlH4− and BH4−). A detailed discussion of these is provided recently by Jensen et al.74 However, to orient the reader we present a brief summary here. Extensive experimental thermodynamic data for metal hydrides are available, typically including the enthalpy and entropy of H2 desorption. Heats of formation for binary hydrides of the alkali and alkaline earth metals and main-group elements have also been measured; for example, see Bourgeois et al. and references therein.75 Calorimetry and other experimental techniques have been used to estimate M-H bond dissociation enthalpy (BDE) in main group and transition metal hydrides; the latter have received considerable attention in this regard.76−81 For example, the “hydricity” of transition metal complexes in solution, which is the energy required to form a hydride ion (MH → H− + M+), has been measured in a number of cases;81,82 these data are useful for understanding chemical reactivity and bond-breaking events within catalytic cycles. Critical assessments of the literature and associated thermodynamic phase modeling (CALPHAD method) have provided phase diagrams for a number of hydrides,83 and these are a valuable tool for evaluating the reactivity of these materials. However, although thermodynamic data for H2 desorption reactions are plentiful, they do not always constitute dissociation energy (ΔH°BDE) for an individual M-H bond but instead are the average bond energies relative to the condensed-phase products. Consequently, one must turn to other approaches, in particular theory, to obtain detailed insight into the bonding of metal hydrides. In addition to reaction energies, first-principles calculations can provide densities of states, atomic charges, and orbital compositions that reveal the underlying nature of the metal−hydrogen interaction. The most common approach is DFT, which has been applied to virtually all classes of metal hydrides of interest for hydrogen storage75,84−92 and used to screen hydrides for

2. CLASSES OF NANOSTRUCTURED METAL HYDRIDES Bulk metal hydrides have been known for more than two centuries: Gay-Lussac and Thenard reported the synthesis of potassium hydride by reaction of potassium metal and gaseous hydrogen upon heating in 1811.70 In 1874, Hautefeuille and Troost measured the equilibrium plateau pressures for the reactions of metallic lithium, sodium, and potassium with hydrogen.71,72 However, bulk metal hydrides started to be considered as media for solid-state hydrogen storage in the 1960s, when several intermetallic hydrides (e.g., Mg2NiHx, LaNi5Hx, and TiFeHx)73 were discovered and shown to absorb and desorb hydrogen gas reversibly in a wide temperature range (i.e., −50 to 400 °C). While many intermetallic hydrides can absorb substantial quantities of hydrogen on a molar basis, most of these alloys, composed of heavy rare-earth or transition metals, have a maximum storage of about 1−2% hydrogen by weight, with the exception of Mg alloys, which have a capacity up to 7.6 wt % hydrogen. E


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Figure 4. Isothermal hydrogenation profiles of Pd octahedrons (red) and cubes (blue) at 303 K at a hydrogen pressure of 101.325 kPa (1.0 bar H2). (b) Schematic potential energy diagrams of the Pd octahedrons (red) and cubes (blue). (Reproduced from ref 111. Copyright 2014 American Chemical Society.)

(AlH4− and BH4−) is bound ionically to an alkali- or alkalineearth metal. The B−H and Al−H bonds have strong covalent character,92 with X-H bond orders from 0.5 to 0.75 and similar ionicities for the four hydrides. The predicted bond order, B− H ≈ N−H > Al−H, indicates that B−H and N−H are more covalent than Al−H. This is consistent with a later investigation by Sholl and co-workers, who considered the bonding in Mg(BH4)2 and Mg(AlH4)2 and concluded that B− H bonds possess a significantly higher degree of covalency than Al−H bonds.108 The B−H BDE predicted by molecular cluster calculations is a high 384 kJ mol−1,86 which is consistent with the higher thermal stability of the borohydrides. These predictions are all consistent with what is known empirically about the stability of these hydrides. Overall, these results suggest that it may be difficult to influence the thermodynamics of metal hydrides by nanostructuring alone (i.e., steric effects that may alter the arrangement of ions), particularly for strongly ionic hydrides for which Coulombic forces may be difficult to alter. Alternatively, the kinetics of hydrogen desorption and uptake could be dramatically affected. As seen subsequently in this article, however, such conclusions are overly simplistic and other factors, such as surface energy, interactions with “noninnocent” pore walls, internal nanointerfaces, and mechanical strain can exert large effects that can change the behavior of these materials.

specific applications.93−96 An important conclusion obtained from these studies is that, for binary hydrides, the primary factors influencing M-H BDE are (1) the difference in the electronegativities of the metal and hydrogen79 and (2) the cohesion energy of the metal relative to the energy cost for lattice expansion upon formation of the hydride. 75,92 Consequently, the stability of binary hydrides from alkaline metal hydrides through the early transition metals (Sc and Y) is dominated by the charge transfer between the metal atom and hydrogen.75 In the case of complex metal hydrides, the bond energy of the M-H diatomic (M typically being B or Al) is also an important parameter that can be used to characterize the bonding.92 A few specific cases are illustrative. For two of the lightest hydrides, LiH and MgH2, extensive charge transfer leads to negative enthalpies of formation: −89.3 to 99.1 kJ mol−1 for LiH97,98 and −76.2 to −77.0 kJ mol−1 for MgH2.99−104 The bonds in these hydrides have mixed ionic−covalent character, whereas in hydrides of the heavier elements, such as CaH2 and SrH2, the bonding is primarily ionic.74 Proceeding down the periodic table, hydride stability begins to decrease beyond the early transition metals as a result of the competition between metal cohesion energy and electronegativity difference. Miwa and Fukumoto computed heats of formation for first-row transition metal dihydrides and showed, consistent with experimental results, that at Group 8 (Fe), both the metal cohesion energy and the energetic penalty for the lattice expansion upon M-H bond formation reach a maximum. Beyond this point, the dihydrides become unstable.105 A similar conclusion was reached in an earlier study by Smithson et al.88 Deeper into the periodic table, the bonding eventually becomes metallic, with hydrogen located at interstitial sites and allowing variable compositions (effectively, solid solutions) without a phase change. In this case, bonding is much weaker. For example, experimental formation enthalpies for bulk PdHx (x = 0.25, 0.75, and 0.8) are only −17.7 to 19.5 kJ mol−1.106,107 It should be noted that DFT calculations consistently underestimate hydride formation enthalpies; the disagreement with experiment is greatest for alkaline and alkaline earth hydrides (as much as 20 kJ mol−1 for KH).75 Nevertheless, correct trends are predicted, even for complex transition metal hydrides.95 The picture does not change dramatically for complex main group metal hydrides. Yoshino et al. performed a DFT study for the hydrides LiBH4, LiAlH4, LiNH2, and NaAlH4.92 Their analysis indicates that the factors controlling the bonding trends in these hydrides are similar to those for metal dihydrides, with the added factor that the hydride anion

2.2. Nanostructured Binary Hydrides

Most of the metals on the periodic table form stable binary metal hydrides. When the surface of a metal, e.g., palladium, is exposed to hydrogen gas, the hydrogen molecules are initially adsorbed then split into hydrogen atoms, which diffuse on the surface and into the bulk. The absorption and desorption of hydrogen by Pd occurs via cycling between two phases; at low hydrogen concentration, the α-Pd phase is stable, whereas at higher hydrogen concentrations the β-Pd phase is the dominant species. The effect of size on hydrogen storage properties has been known for a while, and cycling under hydrogen gas near the equilibrium pressures needed for the formation of Pd hydrides was reported to change with particle size.107,109,110 Recent work by Li et al. has demonstrated that the morphology of Pd may play a critical role in the absorption rates of hydrogen, and that temperature as well is critical for its uptake, absorption, and diffusion. Figure 4 shows the isothermal hydrogenation profiles of Pd octahedrons (red) and cubes (blue) at 303 K, under a hydrogen pressure of 101.325 kPa (1.0 bar H2).111 In situ electron energy-loss spectroscopy (EELS) on individual palladium nanocrystals F


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ionically, are not of substantial interest for hydrogen storage. While some, such as NaH, CaH2, and LiH, have fairly high gravimetric hydrogen capacities (4.2, 4.8, and 12.5 wt %, respectively), especially compared to the heavy metals, the strength of the ionic bonds and the high energetic costs associated with forming metallic sodium, calcium, or lithium lead to desorption temperatures above 400 °C, significantly greater than the desired release temperatures for near room temperature applicaitons.5 However, these and other metal hydrides are of significant interest for the development of hightemperature thermochemical energy storage systems, for instance, as storage media for concentrated solar thermal applications.122−129

revealed that palladium nanocrystals undergo dramatic changes during the phase transition between the α and β phase, and surface effects dictate the size dependence of hydrogen absorption pressures.112 Among bulk noble metals, palladium is the only metal that can absorb and desorb hydrogen at near-ambient conditions of temperature and pressure. Downsizing metal particles to a few nanometers changes the ratio of surface to bulk atoms and can introduce fundamental changes in the thermodynamics of hydrogen uptake. For example, bulk Ir and Rh form hydride phases only at impractical H2 pressures in excess of 4 GPa. Nanoscaling Ir and Rh to j

ϵss R ij12

(8)

where Etot is the total energy of the system. The first summation is the Coulombic energy, with Qi and Qj denoting the ionic charges and Rij is the interion distance, whereas the second term, in which ϵss is the soft-sphere interaction potential, is taken from the repulsive part of the LennardJones potential, which only contributes to the overall energy when the ions physically overlap. Both sums involve all ions in the material, including both the metal cations and the rigid complex anions. The main advantage of the PEGS approach is that it can generate crystal structures quickly, without the need for the extremely time-consuming development of interatomic potentials. In addition to accurately predicting the crystal structures of a number of bulk alanates and borohydrides,304−309 the nano-PEGS variation of the code was used to generate ground-state structures for NaAlH4 and LiBH4 nanoclusters.144,310,311 The model takes advantage of the ionic nature of metal hydrides by fixing the local arrangement of complex anion groups, for example, the BH4− anion in LiBH4. Not only does the PEGS method accurately predict many known crystal structures, but it can also predict the structure of unknown materials. The predicted structure can then be refined using more accurate DFT calculations to establish the atomic coordinates of all atoms, including hydrogen. Majzoub et al. used the predicted nano-PEGS structures to evaluate the thermodynamic stability of LiBH 4310 and NaAlH4144 nanoclusters via the grand-canonical free-energy minimization approach based on total energies and vibrational frequencies obtained from DFT calculations. Nanoclusters of LiBH4 containing fewer than 12 formula units are predicted to dehydrogenate through intermediates of the form LinBnHm (m ≤ 4n) into (LiB)n clusters (Figure 14). In the sodium aluminum hydride system, the stability of NaAlH4 and AlH3 clusters increases as cluster size shrinks. Interestingly, AlH63− nanoclusters are unstable due to a Jahn−Teller distortion arising from degenerate highest occupied molecular orbitals, leading to decomposition of the normally octahedral species. This instability of the dehydrogenation intermediate Na3AlH6 yields a one-step decomposition in nanoscale NaAlH4. The calculated PCT isotherms for both LiBH4 and NaAlH4 exhibit sloping plateaus due to finite size effects, and the decomposition temperatures of free-standing clusters are found to decrease slightly with size due to thermodynamic destabilization of reaction products.

4.2. Theoretical Structural Determination

4.2.1. Wulff Construction. Theoretical modeling is a key component of understanding the structural features of nanostructured metal hydrides. DFT is the technique of choice to predict the crystal structures and atomic arrangements in metal hydride NPs and nanoclusters. Johnson and Sholl introduced a “top-down” approach to predicting the shape and surface energy of metal and metal hydride nanoclusters.301 The crystal structure of each nanocluster is assumed to be identical to the bulk crystal, aside from relaxation of the atoms in the first few layers near the surface. For each surface exposed on a nanocluster, the energy difference between edge and bulk atoms is characterized by a surface energy. The equilibrium crystal shape of a material can be then predicted using the calculated surface energies and the Wulff construction. The Wulff construction is a method to P


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boundary effects,327 and elastic inhomogeneous effects328,329 in particles and polycrystals. The method has been actively utilized to study phase evolution in materials for electrical energy storage,330−332 but has only recently been applied to hydrogen storage materials. Voskuilen et al.333 developed a phase-field model for describing hydrogen transport and operating reactions for hydriding LaNi5 and TiCrMn. In addition, Ulvestad et al.334,335 employed the steady-state phasefield calculations to analyze the internal strain field and hydriding phase morphology in Pd NPs. Further development of phase field-models for hydrogen storage will provide a means of more systematically exploring the interplay between microstructure and reaction kinetics. 4.3. Thermodynamic and Kinetic Analysis

Accurate determination of the thermodynamics and kinetics of metal hydrides is critical to ascertaining the effect of structural nanoscaling on the hydrogen storage properties. Many of the experimental techniques employed for bulk hydrides are the same for nanostructured materials and are described below. Computational models used to deconvolute energetic contributions from various size-related properties are discussed in section 6. 4.3.1. Thermodynamic Measurements. The general reaction involving reversible hydrogen storage in metal hydrides can be written as x ABHx(s) F AB(s) + H 2(g) (9) 2 The equilibrium state of this reaction can be described in terms of the equilibrium constant K and enthalpy change by the van’t Hoff equation:

Figure 14. Atomic geometries of various B−H clusters determined from nano-PEGS and first principle calculations. (Reproduced from ref 306. Copyright 2012 American Chemical Society.)

Since the PEGS approach uses a simple energy model consisting of purely electrostatic interactions, it has some limitations in predicting the structures of metal hydrides with a high-degree of covalent character. Other methods, such as constrained evolutionary algorithms,69,312−316 have been proposed for the prediction of metal hydride crystal structures consisting of well-defined molecular units. In this approach, each structural unit is treated as a rigid body with fixed bond lengths, and then the positions of preformed molecules are found through constrained global optimization, a process of finding the most stable crystal packing given a fixed bond connectivity. This evolutionary approach to structure determination has been successfully implemented in the USPEX (Universal Structure Predictor: Evolutionary Xtallography) code.317−322 However, despite progress, prediction of compounds with large unit cells is still challenging. For instance, none of the existing theory methods are able to accurately predict the structure of γ-Mg(BH4)2 with the P61 symmetry, which has as many as 330 atoms in the unit cell.323 4.2.3. Microstructure Modeling. Beyond the crystal structure and nanoscale geometry of the particles, modeling the impact of microstructural features on hydrogen storage at the nanoscale is likewise important. Nevertheless, there have been relatively few attempts to directly model the microstructure in hydrogen storage materials. One example can be found in the work of Michel and Ozoliņs,̌ who investigated microstructural effects on diffusion kinetics in the bulk Na− Al−H system by considering different model arrangements of NaAlH4, Na3AlH6, and Al.169 Another example comes from the work of Wood et al., who considered the thermodynamics of a few different model microstructures in the nanoscale Li−N−H system.19 Although illustrative, these attempts have relied on static models without explicitly considering how the microstructure evolves during the reaction. The phase-field method has emerged as one of the most powerful methods for mesoscale modeling of material microstructural evolution.324 This method is based on the diffuse-interface description that naturally connects the essential thermodynamic and kinetic ingredients of relevant chemical and materials processes.325 Recent advances have extended the method to account for surface effects,326 grain

d ln K ΔH 0 = dT RT 2

(10)

Since ΔG 0 = −RT ln K and ΔG0 = ΔH 0 − T ΔS 0

(11)

(12)

and, because the gaseous hydrogen is the only species with an activity not equal to one, eq 10 can also be represented as d ln peq ΔH 0 = dT RT 2

(13)

or ΔH 0 ΔS 0 − (14) RT R eq 0 where p is the equilibrium pressure for reaction 1; ΔH , ΔS0, and ΔG0 are the standard enthalpy, entropy, and Gibbs free energy change of the reaction; T is the absolute temperature; and R is the universal gas constant. Equation 14 indicates that a plot of ln peq vs 1/T has a slope of −ΔH0/R and an intercept of ΔS0/R. For an equilibrium pressure of 1 bar, the decomposition temperature Tdec is ln peq = −

Tdec =

ΔH 0 ΔS 0

(15)

For a typical standard entropy change of a metal hydride of 120 J K−1 mol−1 H2, at 350 K ΔH0 = 30 kJ mol−1 H2. ΔH0 values of ∼30 ± 10 kJ mol−1 H2 are considered within the Q


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Figure 15. Changes in the rates of absorption (a,d) and desorption (b,e) and the activation energy (c,f) as functions of the hydrogen content in Nidoped (a−c) and undoped (d−f) rGO-Mg composites. (Adapted with permission from ref 35. Copyright 2017 John Wiley and Sons.)

any semblance of a plateau if the disorder is particularly substantial, as found in amorphous and some nanostructured materials.338 If a material has a multistep (de)hydrogenation reaction, as in NaAlH4/Na3AlH6, or contains two hydrides in a mixture, such as MgH2 and Mg2NiH4, multiple plateaus, corresponding to each reaction, will be present in the PCT curves. 4.3.2. Thin Film Hydrogen Content Measurements. In general, the total mass of a deposited film is very small and below the threshold for gravimetric or volumetric measurements described above that are commonly used in determining hydrogen uptake.230 Therefore, alternative analysis methods must be employed to gain insight into the hydrogen content of thin films.339 Nuclear reaction analysis (NRA) bombards the sample with 15N ions at different energies to sample different depths, and emitted γ rays, which are proportional to the hydrogen concentration at each depth, are detected. Elastic recoil detection analysis (ERDA) employs heavy ions, such as Cl7+ or Ag9+, to obtain depth profile concentrations of all elements, including H.232,340 Neutron reflectometry or reflectivity relies in part on the large neutron cross-section of hydrogen and measures the scattering angle of neutrons by a sample based on the scattering lengths of the pure metal compared to that of the metal with hydrogen.341 In addition to the more direct measurements of the presence of hydrogen, other techniques have been developed that sense changes in the parent material as the hydrogen content evolves. In hydrogenography, visible light is used on materials that transition from reflective metals to opaque or transparent hydrides upon hydrogenation, utilizing a property that also has applications in switchable mirrors and windows.342,343 IR emission also has been employed, monitoring changes in intensity over time and even over area as a material is cycled.28 The IR emissivity depends on the electrical resistivity, which decreases as a metal is hydrogenated to a semiconducting or insulating hydride, and is itself often used as a measure of uptake.230,234 XRD is another commonplace indirect analysis method and is sensitive to both small changes in lattice parameter, at low hydrogen contents, to complete structural

range of interest for near-ambient temperature reversible hydrogen storage applications. For a hydride with an equilibrium pressure of 1 bar, a 10 kJ mol−1 H2 variation in ΔH0 results in about 80 K change in the decomposition temperature. The thermodynamics, particularly the equilibrium pressure peq, are typically measured using a volumetric Sieverts-type instrument. A sample of known initial mass and composition, either with or without hydrogen, is loaded into a pressure vessel, and the volume of the sample holder is calibrated using a typically inert gas and the calibrated volumes and pressure transducers inside the instrument. The two calibrated volumes, both that exclusively within the instrument manifold and the total volume including the pressure vessel, are utilized to determine the amounts of hydrogen dosed and the final remaining gas not adsorbed or absorbed by the sample, respectively. Frequently, numerous measurements are obtained at a variety of pressures and, often, temperatures to sample a representative portion of the pressure−composition−temperature space. In general, the pressure is modulated sequentially, gradually adding more or removing more hydrogen from the system to shift the composition, while remaining at constant temperature, and several such isotherms are measured to yield the van’t Hoff plot. Gravimetric instruments yield identical information as volumetric setups, but the compositional shifts are determined by changes in the measured mass of the material, rather than from differences in the dosed and final hydrogen amounts, as it absorbs and desorbs H2. In a simple interstitial hydride such as Pd, a flat plateau in a pressure−composition isotherm is established, with little pressure change corresponding to a substantial change in the hydrogen content in the material as the solid solution of α and β phases evolves from mostly α to mostly β.336 The midpoint of this plateau, as long as it remains at constant composition, is typically chosen as the value of peq. In more complex systems, however, a sloping, rather than flat, plateau is obtained; the slope may result from atomic disorder, such as defects or noncrystallinity, or from nonequilibirium processes, as found in irreversible reactions.337 Some hydride systems completely lack R


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Figure 16. (a) Computed activation barrier for early hydrogenation of MgB2 from Arrhenius analysis. (b) Reaction energy landscape for the twostep mechanism of dissociation and diffusive adsorption, with adsorption energies derived from ab initio calculations. (c) Experimental and simulated uptake curves showing the change in rate-limiting step. (d) Schematic of the proposed mechanism. (Reproduced with permission from ref 292. Copyright 2017 The PCCP Owner Societies.)

method to analyzing multiple isothermal data for rGOencapsulated Ni-doped Mg NPs to extract the activation energy.36 By interrogating the varying activation energy with the extent of the hydrogen absorption/desorption reactions (see Figure 15), they could identify the rate-limiting mechanisms for different reaction stages of (de/re)hydrogenation.

transformations from the metal crystal structure to that of the hydride. For thin films, grazing angle, also known as glancing incidence, XRD is employed so that the X-ray beam samples the film preferentially, rather than the substrate. 4.3.3. Rate Analysis and Kinetic Models. Kinetic modeling and analysis of hydrogen storage materials is critical to developing an understanding of rate-limiting processes during (re/de)hydrogenation. In many instances, kinetic insights can be obtained by fitting time-dependent reaction data to well-known models of (de)hydrogenation that exhibit characteristic uptake behavior.344 Dominant rate-limiting mechanisms can then be straightforwardly identified from model parameters derived from the fit.345 Alternatively, more complex mechanisms may require advanced approaches that directly extract activation energies for relevant kinetic processes. One way to obtain kinetic data for a metal hydride is to measure the time-dependence of the hydrogen uptake or desorption during acquisition of a PCT isotherm. Small perturbations in the dosed gas pressure, generally though not necessarily while on the equilibrium plateau, lead to recovery over time toward the initial sample pressure, as the stored capacity and gaseous H2 content shift relative to each other, restoring equilibrium according to Le Chatelier’s principle.337 The rate of this restorational change can then be used, along with those measured at other temperatures, to fit the Arrhenius eq 16, in which k is the rate constant, A is the pre-exponential factor, and Ea is the activation energy.346,347 The activation energy for the (de)hydrogenation reaction can thus be extracted from the slope of the plot of the natural logarithm of the rate (or rate constant) vs the reciprocal absolute temperature. In principle, this process can be extended across the entire range of (de)hydrogenation, yielding Arrheniusderived activation barriers for each degree of hydrogen content in the metal hydride. The activation energy as a function of reaction progress can then be utilized to analyze dynamic changes in rate-limiting processes and/or to inform the more detailed kinetic models. For example, Cho et al. applied this

ln(k) = −

Ea ij 1 yz jj zz + ln(A) R kT {

(16)

Another common approach to obtaining Ea is the Kissinger method.348 This method is usually used to determine the activation energy from nonisothermal (de)hydrogenation reaction data with different heating rates, typically using differential scanning calorimetry (DSC) systems, and has been widely used in hydrogen storage research.173,176,194,294,344,349 The linearized Kissinger eq 17, in which A is a constant, assumes a constant heating rate, β, leading to a temperature at which the maximum rate occurs, Tm, seen as a peak in a DSC curve which shifts to higher values as β increases. Kissinger analysis has the advantage of requiring fewer measurements overall. However, it assumes only a single rate-limiting process, so the composition range of interest must be chosen carefully if multiple processes are kinetically competitive. Though many standard DSC instruments operate at ambient pressure with gas flow, high-pressure variants are available to maintain a H2 overpressure and thus interrogate compositional spaces otherwise inaccessible. ij β yz ij AR yz E ij 1 yz lnjjj 2 zzz = lnjjj zzz − a jjj zzz jT z j Ea z R jk Tm z{ k { k m {

(17)

Kinetic models and analyses have been applied to analyze a wide scope of kinetic processes ranging from the surfacerelated initial stage of (re/de)hydrogenation to the deeper hydriding/dehydriding phase transformations. For deeper hydrogenation processes, one of the most common approaches S


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if rotational modes associated with molecular anions are sufficiently anharmonic at operating temperatures. Indeed, in NPs, the larger free volume associated with surfaces, interfaces, and grain boundaries is likely to exacerbate anharmonic effects. Zarkevich et al.357,358 considered the anharmonic contributions to the enthalpies of phase transformations and hydrogen storage reactions in LiBH4. By computing the degrees of freedom originating from anharmonic motion using ab initio molecular dynamics and classical statistical mechanics, they showed that including these effects brought phase transition temperatures and melting enthalpies into better agreement with experiments. DFT methods for computing enthalpy and entropy can be extended to predict the thermodynamic stability of NPs by combining calculations of bulk and surface regions. The most straightforward method involves determining the effective surface energy of the particle by considering the surface energies of various individual orientations and terminations, then weighting them relative to their expression within an equilibrium Wulff shape. The effective surface energy is then scaled according to the surface area and combined with the bulk energy to estimate the Gibbs free energy of the particle.19 Note that this approach assumes that the particle assumes an ideal Wulff shape with well-defined facets, and that the core region behaves like the perfect bulk material. As a result, it is likely better suited for larger NPs than for smaller NPs, of which disordered surface regions comprise a larger fraction of the material and core regions are difficult to delineate. Recognizing this difficulty, a number of authors have turned to explicit modeling of cluster geometries within DFT. For clusters, the size-dependent particle geometry of nanoscale metal hydrides has been predicted by a variety of methods, including molecular dynamics simulations,15 cluster expansion,359 and PEGS combined with a genetic algorithm.310 For example, using cluster expansion on NaAlH4, Mueller et al. proved that particles larger than 5 nm in diameter expose prominent low energy surfaces, as expected, but surface facets and structures become less clear in smaller particles.359 In addition to surface energies, the energy of internal interfaces and grain boundaries can contribute meaningfully to the thermodynamics of metal hydride NPs, particularly when phases coexist during (de)hydrogenation. Interface energies can be computed by explicitly attaching two phases within periodic boundary conditions in DFT360−362 or by using appropriately parametrized classical force fields. However, because determining lattice commensurability and preferred interface orientation can be difficult, this approach is best suited for simpler intermetallic and binary hydrides. For complex hydrides and other systems featuring highly disordered interfaces, the chemical composition and structure of the phase boundaries can be difficult to model explicitly (see section 6.2). To this end, Wood et al. introduced a simpler method for approximating the internal interface energetic penalty in the Li−N−H system as a weighted sum of surface energies of involved phases.19 Because the weights of individual surface energies were difficult to discern a priori, the authors instead employed a statistical sensitivity analysis to determine the likelihood of phase expression as a function of the weighting factor. From this analysis, it was possible to demonstrate the critical roles of internal nanointerfaces in determining the (de)hydrogenation reaction pathways of the Li−N−H system. Further developments in accurate modeling of interface energetics are recommended.

applies the Johnson-Mehl-Avrami−Kolmogorov (JMAK) model, which describes the hydrogenation process as a phase transition operated by simultaneous nucleation, growth, and impingement of a new phase inside the parent phase.344 A recent review article by Pang and Li documented the JMAK model, as well as a number of other predefined kinetic models, such as geometrical contraction models for interface- or diffusion-controlled processes.344 Their review also provides a comprehensive overview of the assumptions and derivations of these kinetic models, as well as several examples of application to hydrogen storage. Within certain reaction regimes, it may be difficult to identify the specific rate-limiting steps. One example is the earliest stages of (de)hydrogenation, in which governing processes may differ significantly from later stages dominated by diffusion, nucleation, and/or growth. A recent example of this can be found in the work of Ray et al.292 They employed a multiscale approach that combines quantum-mechanical calculations, the Arrhenius analysis method, and a kinetic model to verify the relevant heterogeneous multistep chemical and materials processes during the initial stages of hydrogenation of MgB2. In their integrated model, the energy landscape for the entire process (Figure 16b) was constructed by combining the hydrogen binding energetics from DFT calculations and the activation energies (Figure 16a) from the Arrhenius analysis of isothermal hydrogen uptake curves for three different temperatures. The energy landscape parametrized the corresponding kinetic model consisting of coupled differential equations that describe coevolving kinetic processes, which verified the rate-limiting processes (Figure 16c). 4.3.4. Thermodynamic Models. Fundamental phase diagrams for several metal−hydrogen systems have been constructed by conventional CALculation of PHAse Diagrams (CALPHAD) modeling method,83,350,351 which relies on inputs from experimental measurements and ab initio calculations of thermodynamic quantities. Other similar techniques informed by DFT enthalpy calculations have also been proposed for computing thermodynamic or thermochemical equilibria of multicomponent hydrogen storage materials. Kim et al. demonstrated a thermochemical equilibrium calculation approach using data obtained from DFT computations.352 They investigated the thermodynamically stable solid and gas phases of LiNH2, LiBH4, Mg(BH4)2, as well as their mixtures with destabilizing compounds. Akbarzadeh et al. determined the hydride phase diagrams of the Li−Mg−N−H system353 by combining DFT calculations with the grand canonical linear programming (GCLP) method, which determines phase expression by minimizing the grandcanonical Gibbs free energy across several simultaneous competing reactions.354 This method was also used to investigate the phase equilibria of the Li−Ca−B−N−H system.355 Wood et al. introduced an extension of the GCLP method that further incorporates phase coexistence by considering the entropy of mixing within an ideal mixing model. Within DFT, the temperature dependency of the Gibbs free energy is usually included by employing the quasi-harmonic approximation, which extrapolates thermal effects from zerotemperature phonon frequencies.306,353,356 This approximation is well suited for considering thermal effects of high-frequency modes such as bond stretching; however, it has been suggested that inaccurate results may be produced for complex hydrides T


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Figure 17. Illustration of supporting materials for nanoconfinement: (a) porous host matrix (porous carbon, silica, MOF, carbon fiber, and aerogel), (b) ligands (surfactants and polymers), (c) wrapping carbon (graphene and graphene oxide) and synthetic methods for producing nanomaterials: melt infiltration, solvent impregnation, and solution-based synthesis.

5. EFFECTS OF MORPHOLOGY ON HYDROGEN STORAGE PROPERTIES For nanostructured materials, morphology has special significance since it dictates observed physical and chemical properties. Morphology is attained during particle growth through a self-assembling process dictated by the interplay of size and molecular interactions.363 Deviations from bulk properties become prominent as the sizes of nanomaterials start to be comparable to the size of constituent blocks. Properties of metal hydride nanocomposites are controlled not only by morphology of individual nanomaterials, but also by the nature of interactions, which, in turn, are determined by the distribution of the nanomaterials in the hydride matrix. Reducing the particle size of metal hydrides to nanoscale dimensions offers new possibilities to optimize the performance of hydrogen storage. Ball milling, one method of nanoscaling, yields freshly prepared surfaces that are highly unstable, such that the newly created NPs (10−50 nm or larger) can easily aggregate into much larger particles (200− 300 nm).11 An alternative approach to creating stable NPs is to

use supporting materials, which can stabilize NPs to 50 nm23 and hexagonal 50−500 nm Mg particles210 show no change in thermodynamic properties, but dramatic improvements in the kinetics, with activation energies being decreased from 120 to 38 kJ mol−1 H2.23 Urban et al. reported air-stable 5 nm Mg NPs with reduced activation energies for hydrogen absorption (25 kJ mol−1) and desorption (79 kJ mol−1), and a cycling temperature of 200 °C.21 Sub-10 nm Mg NPs show an even more dramatic improvement in H2 storage properties, with fast kinetics and improved thermodynamics which lead to hydrogen cycling below 120 °C.52,367 Aguey-Zinsou reported that LaNi5 NPs of about 26 ± 8 nm have an enthalpy decrease of 5 kJ mol−1 H2 for absorption and 8 kJ mol−1 H2 for desorption, while the entropy decreased by 15 J K−1 mol−1 H2 for absorption and 19 J K−1 mol−1 H2 for desorption.368 Methods to synthesize free-standing NPs of complex metal hydrides are scarce and have only been demonstrated for a few materials. Varin and Parviz explored the influence of the addition of TiN, TiC, and ZrC NPs on the activation energy of LiAlH4 dehyrogenation.369 LiAlH4 was mixed with separately 5 wt % of each of the compounds in a high impact ball mill. Interestingly, the activation energies for the first dehydrogenation step are not strongly influenced and remain in the known range for the bulk material, amounting to 87.7, 89.5, and 96.4 kJ mol−1 H2 for the composites containing TiN, TiC, and ZrC, respectively. However, a drastic decrease is observed for the second dehydrogenation step, with activation energies of 76.6, 79.5, and 63.4 kJ mol−1 H2 for the TiN-, TiC-, and ZrCcontaining composites, respectively. Liu et al. studied the effect of doping LiAlH4 with micro- and nanosized TiH2.370 The doping was achieved by ball milling and resulted in a marked decrease of the activation energy by 20 and 24 kJ mol−1 for the micro- and nano-TiH2 doped LiAlH4, respectively. Christian and Aguey-Zinsou used an antisolvent approach to synthesize NPs of sodium borohydride coated with a shell of nickel.371 In contrast to bulk NaBH4 which releases hydrogen >500 °C, the NaBH4/Ni nanocomposite can be cycled as low as 350 °C with a 5 wt % capacity; 80% of hydrogen can be desorbed and absorbed in less than 60 min, and full capacity is reached within 5 h. A similar approach used to LiH/Ni core− shell particles resulted in a dramatic improvement in the kinetics of hydrogen desorption with about 10 wt % hydrogen being absorbed and desorbed at 350 °C, a significant improvement compared to bulk LiH (700 °C). Pang et al. developed a mechanochemistry/PVD approach for the synthesis of 20−40 nm Mg(AlH4)2 nanorods and 10−40 nm LiBH4 nanobelts185 with improved kinetics of hydrogen desorption; importantly, the morphology of the Mg(AlH4)2 nanorods persisted after hydriding/dehydriding cycles. Li2NH hollow nanospheres prepared by plasma metal reaction of lithium with ammonia show a surface area of 79 m2 g−1. The porosity coupled with particle morphology leads to enhanced hydrogen storage kinetics. The Li2NH hollow nanospheres absorb 6.0 wt % hydrogen in 1 min at 200 °C; in addition, the desorption temperature is decreased by about 115 °C compared with the bulk sample (Figure 19).223 The same group reported nanocubes of Mg3N2 which can be converted to Mg(NH2)2 hollow nanospheres by reacting with ammonia gas.224 The Mg(NH2)2 nanospheres milled with MgH2 release about 6.5 wt % hydrogen be heating to 400 °C, with the initial 2.5 wt % hydrogen released between about 75 and 225 °C. Table 3 summarizes the synthesis and hydrogen storage properties of numerous freestanding metal hydride NPs.

that the hydrogen binding energy differs for Pd atoms on the surface compared to those in the core. Kitagawa et al. investigated the Pd/Pt core/shell-type bimetallic NPs and found a new hydrogen absorption site in the nanointerface between the Pd core and Pt shell.364 Yamauchi et al. observed the size dependencies of the hydrogen-storage properties of coated Pd NPs with diameters of 2.6 ± 0.4 and 7.0 ± 0.9 nm.107 The authors chose poly(Nvinyl-2-pyrrolidone) (PVP) to passivate the Pd NPs because it binds well to metal surfaces. The carbonyl groups of PVP partly coordinated to the surface Pd, which limited the growth of the Pd NPs. As a result, the particle size of Pd increased as the amount of PVP decreased. The Pd NPs showed significantly different isotherms from those of Pd bulk; the NPs displayed higher hydrogen uptake at low pressure and a smaller change in enthalpy (Figure 18).

Figure 18. TEM images of Pd NPs with diameters of (a) 2.6 ± 0.4 nm (Pd−S) and (b) 7.0 ± 0.9 nm (Pd-L). (c) Absorption pressure− composition isotherms of the Pd−S (red), Pd-L (blue), and bulk Pd black (black) at different temperatures from 303 to 393 K. (Adapted with permission from ref 107. Copyright 2008 American Chemical Society.)

Magnesium hydride has been studied extensively as it has one of the highest capacities of all binary metal hydrides. NPs of MgH2 typically display accelerated kinetics of hydrogen release and/or uptake.23,121,21,365,241 Liu et al. isolated 8−25 nm Mg NPs using an electroless approach and found that the thermodynamics significantly deviate from that of bulk MgH2 with a significant decrease of both enthalpy (63.5 kJ mol−1 H2) and entropy (118.4 J K−1 mol−1 H2) for particle sizes below 25 nm.366 Experiments by Paskevicius et al. on MgH2 NPs embedded in a LiCl matrix show that the enthalpy decreases by 2.8 kJ mol−1 H2, coupled with an entropy decrease of 3.8 J K−1 mol−1 H2.186 V


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Figure 19. (a) Hydrogen sorption curves of Li2NH hollow nanospheres (marked by N) and micron-sized particles (marked by M). (b) TEM image of the as-prepared hollow nanospheres. (Adapted from ref 223. Copyright 2008 American Chemical Society.)

5.2. Thin Films

lattice parameter, and the electronic effect of making them less metallic had a lesser absolute effect.234 Several other transition metals have been investigated in thin film format. Palladium is often used as a capping layer in these systems because it facilitates hydrogen activation from the gas phase and protects the highly oxophilic metals against oxidation from oxygen and moisture. Iron−titanium was evaporated onto silicon and covered with 20 nm of Pd.228 The resulting material was partially amorphous and could be charged to 0.35 H per metal atom, with a discharge down to 0.12 H/M, indicating deep trapping sites. In addition, it was possible to charge the amorphous film with hydrogen at higher temperatures and lower pressures than microcrystalline films due to the high disorder. Zirconium was cosputtered with Ni, Co, Fe, and Cu onto glass and cycled from vacuum to 1 bar.230 The change in electrical resistance upon hydrogenation was most pronounced in the Ni-containing samples, and higher desorption temperatures restored the low resistivity better due to increased ripening and growth of the grain sizes. Hamm et al. aimed to destabilize metal hydrides using the clamping stress from the substrate (estimated to be −1.1 kJ mol−1 H GPa−1) but needed the strong adhesion of niobium on sapphire Al2O3 to prevent delamination and dislocation formation.378 Only films thinner than 6 nm were able to avoid plastic deformation and achieve the high stress desired up to 1 H/M, though 10 nm films could deform elastically after an initial plastic deformation cycle due to hardening from interactions of dislocations. Magnesium has received much attention in thin film research because of its lightweight and high gravimetric capacity (7.6 wt % H); it also absorbs hydrogen very poorly in the bulk due in part to the slow H diffusivity in Mg, making research into nanostructured magnesium especially relevant.67,379−381 Dura et al. prepared 100 nm Mg thin films with 50 nm of Pd as a cap on monocrystalline Al2O3 and found, using both scanning electron microscopy (SEM) and the scattering length from neutron reflectivity, that the porosity increased from 4% to 25% upon cycling.382 Gautam et al. sputtered rather thick films (500 and 1000 nm) of Mg capped with Pd.231,232 They found that the γ-MgH2 phase forms as a result of compressive stresses while annealing and that the absorption amount, mean crystallite size, and number of defects all increase with absorption temperature in the 150 to 250 °C range. A significant drawback to the capping of Mg with palladium is the formation of a variety of intermetallic species at the interface between the two. Transition metals have been added into thin films of magnesium to act as catalysts as well, either codeposited or in

Thin films differ from NPs in both the synthetic methods and the resulting structure, necessitating unique analytical methods for determining hydrogen sorption properties. Unlike free NPs, though similar to confined or encapsulated particles, thin films have a rigid substrate upon which they are deposited. Additionally, films have two macroscopic dimensions, with only the thickness being on the nanoscale. These materials, therefore, are dominated by the surface facing away from the substrate, leading to highly anisotropic communication with the gas phase. However, thin films are rarely monocrystalline but instead typically highly textured, with small crystallites, numerous grain boundaries, and occasionally even some porosity, depending on deposition conditions and method. A wide variety of metals have been formed as thin films for numerous different hydrogen-based applications, including hydrogen-switchable windows and mirrors339,372−374 and metal-hydride batteries.375 While these are intriguing fields, these properties are not germane to this review, which instead will focus on the changes in H2 sorption relevant to hydrogen storage, particularly in lighter materials that are of potential interest for on-board vehicular storage. Palladium has often been studied due to its simplicity, fast kinetics, and isostructural hydride. Di Vece et al. prepared MBE and nanocluster Pd films on glass and investigated their behavior with both extended X-ray absorption fine structure (EXAFS) and optical transmission spectroscopy to determine the effect of hydrogen cycling.226,227 The authors found that hydrogenation led to Ostwald ripening, even at ambient temperature, due to reduction of the binding energy in the hydride compared to the metal. Additionally, the smaller nanoclusters lacked the miscibility gap, evidenced by a plateau in the pressure as H content changes, between the α and β Pd−H phases that was evident for the MBE film and is more typical of the bulk material. Similarly sized films electron-beam deposited on quartz showed a narrowing of the plateau, particularly for 10 nm thick films, due to the induced stress from clamping of the film to the substrate.376 Khanuja et al. compared both thin films and layers of discrete NPs, deposited on Al on glass, to investigate differences in resistivity changes.377 Hydrogenation was slowed in the films due to in-plane strain and clamping effects that were absent in the NP layers which had interparticle gaps. Using sub-10 nm Pd NPs synthesized by reduction from PdCl42− with ethylene glycol and dropcast onto glass, Gupta et al. found that the change in resistivity was dominated by the geometric effect of the Pd particles expanding to form the hydride, which has a larger W


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Table 3. Hydrogen Storage Properties of Freestanding Metal Hydride NPs material

size (nm)

shape

synthesis method

H2 wt % ∼0.4 ∼0.5

Pd Pd Pd

2.6 7.0 14

sphere sphere cube

solution-based chemistry/PVP solution-based chemistry/PVP solution-based chemistry/CTAB

Pd

32

cube

solution-based chemistry/CTAB

Pd

65

cube


Pd

110

cube


Pd Pd Pd Pd Pd Pd Pd Pd Pd Pd Pd Pd Pd Pd Pd Pd Pd Rh Rh Rh Ir Ir Li Mg Mg Mg

18 24 30 32 40 42 43 45 54 60 63 85 85 85 38 × 137 38 × 215 47 × 185 2.3 2.4−7.1 7.1−10.5 1.5 1.5 4 2−7 15 30−50

cube cube cube cube cube cube cube cube cube cube cube octahedron octahedron octahedron nanorod nanorod nanorod irregular irregular irregular irregular irregular spherical sphere irregular nanowire

solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry Solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry solution-based chemistry

Mg

80−100

nanowire

gas phase condensation

Mg

150−170

nanowire

gas phase condensation

Mg Mg

6.0 40

irregular hexagon

Mg(n-Bu)2 hydrogenation acetylene plasma metal reaction

5.7 5.5

Mg

5

sphere

Cp2Mg reduction with Li-naphthalide in PPMA

5.97

Mg Mg

7−15 25

irregular irregular

spark discharge Cp2Mg reduction with potassium

4.2 6.7

Mg

32

irregular

Cp2Mg reduction with potassium

6.6

Mg

38

irregular

Cp2Mg reduction with potassium

6.1

Mg MgH2

8−25 11−20

spherical fibers

∼6.2 ∼2.6

Ni-doped MgH2

5.4

irregular

electroless reduction stabilization with azide-terminated poly (styrene) Mg(n-Bu)2 hydrogenation

Mg/Ti Mg/Ti

10−20 12

irregular irregular

gas phase condensation spark discharge

solution-based chemistry, PVP solution-based chemistry solution-based chemistry ball milling in LiCl electroless reduction gas phase condensation

18.0b 47.1c 20.2b 50.6c 22.6b 54.4c 25.3b 60.5c

|ΔH| (kJ mol−1)a

|ΔS| (J mol−1 K−1)a

35b 31b 27.4b 33.4c 28.5b 35.5c 29.9b 37.6c 35.4b 40.9c 28−36b 27−32b 32−36b 31−37b 30−35b 31−35b 29−32b 28−34b 32−36b 35−41b 33−42b 27−31b 30−38b 29−32b 33−38b 34−38b 33−39b

83b 67b 61.5b 75.8c 66.9b 82.3c 72.6b 88.2c 82.4b 94.4c 64−86b 61−78b 79−90b 74−94b 74−86b 75−87b 70−80b 65−85b 79−92b 88−107b 83−108b 62−75b 71−96b 68−78b 81−96b 84−98b 78−97b

71.2b 63.5b 65.3c

129.6b 118.4b

∼0.15 ∼0.1 ∼0.1 ∼0.09 ∼0.09 2.0 142.8c 33.5b 38.8c 38.7b 46.5c 70.3b 81.1c

5.4

4.0 X

Ea (kJ mol−1)

61.6b 114.0c 25b 79c 60−120b 122b 126c 118b 131c 115b 160c

109 109 109 418 418 418 418 418 418 418 418 418 418 418 418 418 418 418 418 418 113 114 115 116 117 419 186 273 23

65.9c

23

67.2c

23

64.8c 65.5c

122.7

420 121 21 365 241 241 241

63.5

22.7b 64.7c

ref 107 107 109

118.4

62.1b 68.1 45

366 421 420

119 84

119 422


Chemical Reviews

Review

Table 3. continued material

size (nm)

shape

synthesis method

H2 wt %

Ea (kJ mol−1)

|ΔH| (kJ mol−1)a

ref

Mg/Ti

50

irregular

spark discharge

4.0

170.9c

Mg/V

50−150

hexagons

hydrogen plasma metal reaction

5.2

71.2b 119.4c

Mg/Co Mg/Ti/Pd Mg/Ni

50−150 60 10−20

hexagons nanodots irregular

hydrogen plasma metal reaction molecular beam epitaxy chemical reduction

4.9

hexagons

arc plasma

5.1−6.1

irregular irregular irregular irregular irregular irregular irregular irregular irregular spherical nanorods hollow nanospheres hollow nanospheres irregular nanobelts spherical spherical hollow nanospheres spherical

mechanical milling mechanical milling mechanical milling mechanical milling reduction of precursor sintering Mg/Fe NPs chemical reduction chemical reduction chemical reduction stabilization with 1-dodecanethiol physical vapor deposition plasma metal reaction

4.2 4.6 6.1 6.3 4.95 5.0

plasma metal reaction

6.0

223

milling under H2 physical vapor deposition stabilization with PMMA stabilization with PMMA stabilization with PMMA

4.0

435 185 436 436 436

solvent precipitation

5.0

371

Mg@Ni Mg/Ti,Fe,Ni MgH2/TiH2 MgH2/TiH2 MgH2/Nb2O5 Mg2FeH6 Mg2FeH6 Li/Mg Na/Mg K/Mg LiAlH4 Mg(AlH4)2 Mg(NH2)2/ MgH2 Li2NH

5−10 5−10 100 100 8−40 7−52 8−32 2−16 20−40 100 100

MgB2 LiBH4 LiBH4 NaBH4 Ca(BH4)2

10−40 30−50 50−120 100−300

NaBH4/Ni

10−30

6.0

5.7 9.0 6.5

139.1c 57.4b 88.9b

73.0b 75.8c 74.3b 74.4c 82.3c 70 70.0b 70.7c 73.7b 81.3c

|ΔS| (J mol−1 K−1)a 130.9b 134.8c 155.8b 152.9c 138.8c 129

423

105.5b 104.5b

427

45.7 16.4

165.7c 159b 148b 134b 97.3c

68.2c

127c

89c 82.4c 66.9 71.1 63.2

145c 140.2c 122.1 130.9 117.8

424 425 426 242

428 429 430 431 432 433 434 434 434 417 185 224

The absolute magnitude of ΔH and ΔS. bValue obtained from absorption measurements. cValue obtained from desorption measurements.

a

kinetics and capacity under cycling because the Mg/MgH2 particles could not grow as large and limit the transport of hydrogen. Ti, like Fe, is immiscible with Mg and, like Pd, has been used to protect the Mg and dissociate H2, while forming TiH2. Baldi et al. made films with layers of 10 nm Ti and 20 nm Mg, capped with Pd, on glass.343,384 Magnesium sandwiched between two layers of titanium absorbed H2 at lower pressures than Mg between Ti and Pd (Figure 20), which they explained as due to clamping stress and alloying with the palladium. Mooij and Dam conducted follow-up experiments, again with Ti-sandwiched Mg, that showed that, in a 10 nm Mg film, the hydride phase nucleates in the same locations repeatedly, but upon dehydrogenation, the metallic phase has no visibly obvious nucleation points but appears to fade in because the barrier to Mg nucleation is small and thus the number of nuclei is large.385,386 Bannenberg et al. continued the study with neutron reflectivity and hydrogenography measurements on deuterated Ti−Mg films, showing that the nucleated β-MgH2 regions grow at a constant radial rate, with plastic deformation of the encompassing Ti and Pd layers to fit the larger phase.341 As in the study by Dura et al., the Mg layer was thicker after dehydrogenation, indicating increased porosity and enabling faster rehydrogenation kinetics. An X-ray photoelectron and Auger spectroscopy investigation on cosputtered Mg−Ti films was conducted by Jensen et al. and found, in addition to some partially oxidized

discrete layers. Nickel forms Mg2Ni, which can be hydrogenated to Mg2NiH4. Oguchi et al. synthesized 100 nm thick combinatorial Mg1−xNix thin films capped with Pd.28 The compositions with more Mg showed greater hydrogenation via IR emission imaging, and using XRD, the authors found that a nickel-stabilized fcc Mg phase enabled faster kinetics and a lower temperature of reaction by transforming to a similarly structured fluorite-type hydride. Fry et al. made a multibilayer film, capped with Pd, with 150 pairs of 16.5 nm Mg and 2.5 nm Ni−Fe−Cr (predominantly nickel).42 PCT measurements, made possible by the large amount of material, showed faster hydrogenation than ball-milled Mg and two different plateaus: one for MgH2 and one for Mg2NiH4, indicating mixing between the layers. After cycling, the delaminated films broke down into powder containing Fe NPs as well as Mg−Ni and Mg−Pd intermetallics. A 4% Fe-doped Mg thin film was deposited on silicon in a wedge shape to study the effect of film thickness on hydrogenation.229 The rate of hydrogenation, as determined from IR emissivity, was limited by the diffusion of hydrogen through MgH2; however, PCT and van’t Hoff studies showed that the enthalpy and entropy for the reaction were less than those for pure magnesium films. A pair of multilayered Mg/FeTi thin films with different layer thicknesses retained the intact FeTi layers upon cycling, while the Mg layers formed NPs confined between the sandwiching layers.383 The thinner layers had more stable Y


Chemical Reviews

Review

ultrasmall pore size (from 2 to 0.5 nm) that can be achieved via various synthetic methods.50,51,349,388−393 De Jongh et al. have prepared high-purity and oxygen- or nitrogen-modified ACs in order to synthesize Mg/C nanomaterials.51 The highpurity carbon support has a nonpolar nature, which might impede the infiltration of molten magnesium into the pores. To improve the wetting on the surface of the porous carbon, nitrogen, or oxygen was introduced. Mg nanomaterials were prepared by melt infiltration, and the average size of Mg decreased from a few nanometers for hierarchical micro and mesoporous AC with ∼10 at. % O), to ∼2 nm for high surface area graphite with mainly 2−3 nm pores to less than 1−2 nm for microporous AC-based samples (Figure 21a). It is demonstrated that AC is able to mechanically confine Mg nanomaterials inside the pores and prevent oxidation of Mg shortly after preparation (Figure 21b). The use of AC as a confinement medium has been extended to other materials, such as metal alanates and borohydrides.349,388−391,394 Nanosized NaAlH4 materials in highpurity porous carbon HSAG-500 with relatively small pores (mainly 2−3 nm) were prepared by the melt infiltration method.388 After melt infiltration, the NaAlH4/C nanocomposites were examined by physisorption measurements because the metal alanate and borohydride materials decompose under the electron beam in TEM. The Brunauer−Emmett−Teller (BET) surface area and the Barrett−Joyner−Halenda (BJH) total pore volume decreased from 500 to 156 m2/g and from 0.57 to 0.32 cm3/g, respectively, after infiltration, indicating that the NaAlH4 efficiently filled (or blocked) most of the pores. In a followup study by the group of de Jongh, the possibility of changing the reaction pathway by nanosizing it in HSAG and AC was presented.395 In addition, NaBH4 and LiBH4 nanostructures were also confined in HSAG-500, which was confirmed by broadening and reduction in the intensity of XRD peaks and the decrease in the total pore volume of the carbon, which is in good agreement with the materials completely filling in the carbon pores.389,390 By modifying a conventional method through changing the pH of the initial solution, Nielsen et al. produced several distinct nanoscale pore dimensions in carbon aerogels, which distinctly influenced hydrogen release and uptake properties in the NaAlH4 system.257 Nielsen et al. explored the influence of the pore size of resorcinol-formaldehyde carbon aerogels on the hydrogen desorption kinetics of nanoconfined NaAlH4.16 With increasing pore size of the NaAlH4-loaded scaffold, the onset temperature increases, linking the desorption kinetics to the particle size, which, in turn, is controlled by the pore size. Stephens et al. were able to alter the hydrogen desorption temperature of NaAlH4 by melt-infiltrating the hydride into an aerogel with 13 nm pores.396 Dehydrogenation to NaH was achieved at 150 °C with reasonable kinetics. Additionally, it was possible to rehydrogenate the material fully at 160 °C under 100 bar H2 pressure. Jensen et al. infiltrated a TiCl3doped resorcinol-formaldehyde aerogel with NaAlH4 via melt infiltration.397 The resulting nanocomposite hosted NaAlH4 particles with a size of 37 nm. Gross and co-workers described the synthesis of LiBH4 inside pyrolyzed resorcinol-formaldehyde aerogels with varying pore size (13 and 25 nm) by melt intrusion.398 The activation energy was lowered to 103 and 111 kJ mol−1 for the 13 and 25 nm LiBH4-loaded aerogels, respectively. In comparison, the activation energy for bulk LiBH4 amounts to 146 kJ mol−1.

Figure 20. Mg sample deposited on a glass wafer and partially covered with a 2 nm thin Ti layer. (a) Sketch of the sample architecture. (b) Pressure−temperature isotherms (PTIs) measured during loading at 333 K. (Reprinted with permission from ref 384. Copyright 2009 AIP Publishing.)

Mg, the presence of H atoms trapped in interstitial sites between Ti atomic clusters and the Mg matrix.233 Multilayers of 18 nm magnesium with 1 nm chromium or vanadium as catalysts showed different mechanisms from each other, though the two activation energies were similar.29 Both catalyzed systems were much faster than the Mg control, but the thermodynamics were comparable. A set of Mg2Si−Mg thin films, each capped with palladium, was prepared by sputtering to determine the limitations on hydrogenation of Mg2Si, especially compared to ball-milled particles.387 Adsorbed hydrogen diffused through Mg2Si and reacted with metallic Mg without substantially hydrogenating the silicide; only regions immediately proximal to either the Mg or Pd were affected, so very small particles (350 °C), a distinct change in morphology was observed due to sintering of the Ni shell. Various other metals can be used as a shell (Co, Cu, Fe, Ni, and Sn) in order to AF


Chemical Reviews

Review

retain the shape of NaBH4 core during desorption.416 The Sn shell (melting point = 232 °C) fell apart at 300 °C, the Ni shell was unstable at 600 °C, and the remaining three metals maintained their structures even at the higher temperature (Figure 27). Wang et al. expanded this approach to LiAlH4@Ti and showed significant improvements in hydrogen storage properties of LiAlH4.417

preparation of nanomaterials has been also studied. Table 5 contains details about synthesis and hydrogen storage properties of nanoscaffolded and nanoencapsulated metal hydride NPs.

6. MECHANISTIC EFFECTS OF NANOSIZING Across a vast range of applications, nanomaterials exhibit their own unique chemical and physical properties compared to their bulk counterparts.469 It is well known that various kinds of physicochemical properties can be tuned by manipulating material dimensions. As to metal hydrides for hydrogen storage, many studies have reported that their thermodynamics and kinetics can be improved via nanostructuring. Nanoscale metal hydrides possess high surface areas, leading to an increased surface energy, which is expected to improve hydrogen sorption thermodynamic and kinetic characteristics. Also, nanostructuring plays a key role in generating and enhancing the effects of controlled interfaces between different material phases, which can induce certain forms of lattice deformation or distortion that can be beneficial for cycling. Furthermore, confinement effects can fundamentally alter materials properties, either through chemical or electronic interactions with the confining medium or else through mechanical confinement stress. Potential effects of nanostructuring or nanoconfinement on the thermodynamics of metal hydrides are usually represented in terms of the enthalpy change upon cycling. For most metal hydrides, ΔH is negative (exothermic) upon hydrogenation, indicating the relative stability of the hydrided form, with the notable exception being AlH3. Furthermore, in complex metal hydrides and most light metal hydrides, the magnitude of the reaction enthalpy tends to be too large, which in turn signifies impractically high operating temperatures. This fundamentally originates from the high lattice enthalpy and strong ionic and covalent bonding that characterizes these materials classes.14 Accordingly, one of the primary goals of nanosizing and nanoconfinement is to destabilize the hydride, thereby reducing the magnitude of the reaction enthalpy and facilitating practical use within a reasonable range of operating temperature. For complex metal hydrides, the reaction entropy may be altered in addition to the reaction enthalpy, providing a second avenue for thermodynamic tuning. Likewise problematic are the often sluggish hydrogenation and dehydrogenation kinetics of metal hydrides. Poor kinetics may be associated with slow diffusion, chemical bond activation, nucleation of reaction products, or formation of stable intermediate phases in multistage reactions. As a result, cycling of metal hydrides often requires hydrogen release temperatures and uptake pressures far in excess of what is predicted thermodynamically. Nanoscale and nanoconfined metal hydrides can offer significant advantages in this regard by reducing diffusion lengths, facilitating chemical kinetics and nucleation or altering reaction pathways. Review articles by Berube et al.,11 Pundt and Kirchheim,470 Pundt,471 de Jongh and Adelhelm,45 Vajo,47 and Wang472 systematically surveyed earlier efforts exploring the effects of nanosizing, nanoconfinement, and microstructure on the interactions between hydrogen and solid-state materials. These articles comprehensively documented various effects of surfaces, phase boundaries, grain boundaries, dislocations, vacancies, and other structural defects at the nanoscale. In the following sections, we review some of their conclusions, while focusing heavily on more recently reported/proposed mech-

Figure 27. TEM micrographs of the core−shell structures in the assynthesized (left), 300 °C annealed (center), and 600 °C annealed (right) particles: (a) NaBH4@Co, (b) NaBH4@Cu, (c) NaBH4@Fe, (d) NaBH4@Ni, and (e) NaBH4@Sn. (f) Temperature-programmed desorption of the core−shell NPs at a rate of 10 °C min−1 monitored by mass spectrometry. (Adapted with permission from ref 416. Copyright 2013 Royal Society of Chemistry.)

These porous matrixes are able to decrease the particle size of metal hydrides to less than 10 nm and prohibit the aggregation into larger particles. However, there are a few disadvantages, including the decrease in gravimetric and volumetric hydrogen densities, because this approach has thus far exhibited only limited loading of metal hydride into the dead weight of inert supports. Thus, the development of new synthetic methods diminishing the usage of supports for AG


Ni-CMK-3 RF-aerogel

RF-aerogel

RF-aerogel activated carbon fiber RF-aerogel(Ni) RF-aerogel(Cu) RF-aerogel HSAG activated carbon SNU-90 RF-Aerogel RF-Aerogel RF-Aerogel RF-Aerogel Ni nanobelt Ni-doped hollow carbon HSAG 0.13 TiCl/RFaerogel RF-Aerogel CMK-3 pyrolyzed aerogel pyrolyzed aerogel CMK-3 HSAG RF-aerogel RF-aerogel Activated carbon NbOF5−CMK-3 CNT HSAG

MgH2 MgH2

MgH2

MgH2 MgH2

AH

LiBH4MgH2 LiBH4 LiBH4 LiBH4 LiBH4 LiBH4 LiBH4−Ca(BH4)2 LiBH4−Ca(BH4)2 LiBH4 LiBH4 LiBH4 LiBH4

AlH3 LiBH4MgH2

MgH2 MgH2 MgH2 MgH2 MgH2 MgH2 MgH2 MgH2 MgH2 MgH2 MgH2 MgH2

CMK-3-template CMK-3-CVD CMK-3 MIL-100(Al) MIL-101(Cr) COF-102 HSAG

porous host

PdHx PdHx PdHx PdHx PdHx PdHx Mg6Pd

hydride

1.27 1.3 0.8 1.38 1.3 0.65 1.21 3.13 0.87 1.25 0.66

2.2 4.5 13 25 4.5 2−3 30 30 1.75−3.2 3.8

2−3

0.66 1.3

1.53

3.4

2−3 26

0.8 0.8 0.8 0.65 0.45 1.64 1.32 1.32 1.85 1.89

0.8 0.79

0.65

0.66 1.27

1.14 1.13 1.1 0.65 1.64 1.35 0.69

Vtot (cm3 g−1)

13 13 13 2−3 100

2−3 2−3 13 2−3 10 11 10 8 8 2−3

2−3

0.41 0.41 0.593

0.36

0.66 0.66 0.8 0.65 0.91 1.3 2.21 2.11 1.93 0.65 0.36 0.36 0.36 0.69 0.91 0.86 1.33 1.39 1.3 1.09 0.26

0.66

0.66 0.66 0.28

2−3 2−3 4.55

HSAG CNT HSAG HSAG TiO2/Porous carbon HSAG graphite (300 mesh) HSAG HSAG RF-aerogel HSAG RF-aerogel RF-aerogel RF-aerogel RF-aerogel TiCl3/RF-aerogel HSAG carbon nanofiber carbon nanofiber RF-aerogel RF-aerogel RF-aerogel RF-aerogel RF-aerogel RF-aerogel RF-aerogel RF-aerogel RF-aerogel CNT Ti/Carbon nanofiber HKUST-1 HKUST-1 MOF-74(Mg) SBA-15

Li/LiH/LiBH4 LiAlH4 LiAlH4 LiAlH4 LiAlH4

0.35−0.83 1.63 1.21 1.2 0.65

4.7−6.5 3.5 30 3

SBA-15 CMK-3 RF-Aerogel carbon nanofiber

LiBH4 LiBH4Ca(BH4)2 2LiBH4−NaAlH4 LiBH4−LiAlH4

Vtot (cm3 g−1)

2−3

Davg (nm)

porous host

hydride

Table 5. continued

solvent impregnation solvent impregnation melt infiltration solvent impregnation

solvent impregnation melt infiltration melt infiltration melt infiltration melt infiltration melt infiltration melt infiltration melt infiltration melt infiltration melt infiltration solvent impregnation solvent impregnation melt infiltration melt infiltration melt infiltration melt infiltration melt infiltration melt infiltration melt infiltration melt infiltration melt infiltration solvent impregnation solvent impregnation

physical mixing ball milling

solvent impregnation ball milling + melt infiltration melt infiltration sequential solvent impregnation/ melt infiltration melt infiltration solvent impregnation solvent impregnation solvent impregnation solvent impregnation

infiltration method

THF THF 195 °C, 250 barH2 THF

H2O 520 °C, 5 bar H2 189 °C, 183 bar H2 H2 189 °C, 210 bar H2 189 °C, 210 bar H2 189 °C, 210 bar H2 189 °C, 210 bar H2 189 °C, 210 bar H2 H2, 180−190 bar, 180 °C THF THF 189 °C, 210 bar H2 189 °C, 210 bar H2 189 °C, 210 bar H2 189 °C, 210 bar H2 189 °C, 210 bar H2 189 °C, 210 bar H2 189 °C, 210 bar H2 189 °C, 210 bar H2 189 °C, 210 bar H2 THF THF

Ar LiBH4, NaCl

200 °C, 100 bar H2 THF THF THF THF

MTBE 200 °C, 3 bar H2 310 °C, 110 bar H2 Et2O/310 °C, 110 bar H2

solvent or atmosphere

4 4 21 20

25 25 48.6 20 34.1 41.5 51.3 56.5 48.6 5−80 9 2 12.8 27.3 34.1 31.8 41.9 43 41.4 37.2 32 5.68 8

33

18.9 37

20 1.2

25.3 66

33−50

loading/wt %

ref

460 136 458 461 462

414 265 459 266

4.5a 4.8

0.45

53 57.4

45

413 17 18 272

389 389 396 388 16 16 16 16 16 395 407 50 257 257 257 257 257 257 257 257 257 136 145

47.1

117

|ΔS| (J mol−1 K−1)

6.7a 7.9a 5.5 7.4a 1.3 1.9 2.4 2.7 2.1 6.3 2.84

47.3

|ΔH| (kJ mol−1)

463 408

58

64

Ea (kJ mol−1)

3.8−4.4a 6.8

0.19 10a 0.6 6.2

11a 9.52a 6.8a

H2 wt %

Chemical Reviews Review


2

4.5 4.5 9 4 7.5

CMK-3 CMK-3-Li RF-carbon-cryogel MCM-41 SBA-15

20

9 4.80 20

1.2

1.29 1.29 0.7

1.1

0.128

0.5 0.86 0.128

0.52

0.61

0.65

Vtot (cm3 g−1)

solvent impregnation (Mg(BH4)2·Et2O)) solvent impregnation solvent impregnation solvent impregnation solvent impregnation solvent impregnation

solvent impregnation

melt infiltration (MgH2) solvent impregnation solvent impregnation

ball milling

solvent impregnation

melt infiltration

infiltration method

MeOH MeOH THF THF MeOH

NH3

THF

120 °C, B2H6/H2 THF THF

B2H6/H2

Et2O

180 °C, 120 bar H2

solvent or atmosphere

50 50 24 33−75 50

50

6 (27% MgB12H12) 10

44

20

loading/wt %

3.5 9a

9.3

7.21a

H2 wt %

67

98

21.3

176.2; 189.5c 102

Ea (kJ mol−1)

120

2.1

|ΔH| (kJ mol−1)

|ΔS| (J mol−1 K−1)

ref

467 467 468 468 271

466

289

464 465 289

275

349

460

Hydrogen desorption values are normalized as weight percent of the infiltrated hydride. bHydrogen content is determined by absorption measurements on the corresponding metal/host complex. c Enthalpies for two consecutive reaction steps occurring at different temperatures.

a

Mg(BH4)2 • 6 NH3 NH3BH3 NH3BH3 NH3BH3 NH3BH3 NH3BH3

Ca(BH4)2

Ni-RF-Aerogel Ni-CMK-1 Cu2S hollow spheres Cu2S hollow spheres activated carbon

Mg(BH4)2 Mg(BH4)2 Mg(BH4)2

10

RF-aerogel

Mg(BH4)2

2−3

Davg (nm)

activated carbon

HSAG

porous host

Na/NaH/ NaAlH4 Mg(BH4)2

hydride

Table 5. continued

Chemical Reviews Review

AJ


Chemical Reviews

Review

0.55 J m−2 for Mg.473 Wagemans et al. studied the surfaceinduced destabilization effect in MgH2 using DFT calculations on clusters, predicting a significant reduction of reaction enthalpy (>10%) for clusters below 2 nm in size due to surface effects alone.14 These results generally agreed with the later DFT calculations of Kim et al.,301 as well as direct hydrogenation simulations by Cheung et al. using the ReaxFF method.15 Experimentally reported values for reaction enthalpy reduction in MgH2 upon nanosizing show a somewhat greater variation. Paskevicius et al. reported a reduction of

Nanostructured Metal Hydrides for Hydrogen Storage

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