# How to Design Hydrogen Storage Materials?
## Fundamentals, Synthesis, and Storage Tanks

Qiwen Lai, Yahui Sun, Ting Wang, Poojan Modi, Claudio Cazorla, Umit B Demirci, Jose Ramon Ares Fernandez, Fabrice Leardini, Kondo-françois Aguey-zinsou

► To cite this version:

Qiwen Lai, Yahui Sun, Ting Wang, Poojan Modi, Claudio Cazorla, et al.. How to Design Hydrogen Storage Materials? Fundamentals, Synthesis, and Storage Tanks. Advanced Sustainable Systems, 2019, 3 (9), pp.1900043. 10.1002/adsu.201900043. hal-02310497

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# How to Design Hydrogen Storage Materials?
Fundamentals, Synthesis, and Storage Tanks

Qiwen Lai, Yahui Sun, Ting Wang, Poojan Modi, Claudio Cazorla, Umit B. Demirci, Jose Ramon Ares Fernandez, Fabrice Leardini, and Kondo-François Aguey-Zinsou*

As the world shifts toward renewable energy, the need for an effective energy carrier is pressing. Hydrogen has often been touted as a universal clean energy vector and the fuel of the future. Unfortunately, mass adoption of the hydrogen economy is slow due to a lack of incentives and technical difficulties in storing hydrogen. Better materials capable of reversible hydrogen uptake/ release with hydrogen capacity surpassing 5 mass% at the ambient must emerge. So far, finding such materials has been elusive; alloys capable of ambient hydrogen uptake/release have a low storage capacity while high capacity hydrides have a very high hydrogen release temperature. From metal alloys to complex hydrides, a better understanding of the behavior of hydrogen in hydrides is essential to fine-tune their properties toward application. Herein, the latest approaches to design hydrogen storage materials based on known hydrides are reviewed with the aim to facilitate the emergence of alternative thinking toward the design of better hydrogen storage materials. Synthetic methods and conceptual approaches to achieve particular hydrogen thermodynamics and kinetics are discussed. These include metallurgical alloying, mechanochemical modification, chemical destabilization, the nanosizing approach, and theoretical modeling and machine learning techniques to guide experimental work.

## 1. Introduction

### 1.1. Hydrogen in Broader Context

The fate of human kind, is at a cross-road with significant challenges threatening several concepts that were taken for granted, e.g., technological progress will enable us to defeat the general laws of nature and grow economic wealth beyond the capability

of our planet.[1] Indeed, we have learnt through the greatest minds to master the use of fire to drive machines to do work for us. We are at the corner of taking these machines closer to artificial intelligence capability—the so called 4.0 Industry. But the idea that such technological advances can lead to further wealth and equal rights across the planet remains a debate and most likely utopia.[2] To date, despite all these progress, we are still failing to recognize the limits of our current enterprise in terms of the amount of natural resources available to us and we are brutally lacking imagination and leadership to reinvent our societal organization and models to cope with the changes ahead (Figure 1a). Some call for a “degrowth” economy[3] and a reorganization of economic models; however, this is only one aspect of the solution.

Current energy consumption is projected to double to support the economic growth of the emerging economies. Enabling readily access to large amounts of energy is thus a paramount. Indeed, the viability of our economic models and quality of life are directly related to our ability to access energy.[4] However, the majority of the energy produced is at present consumed by a very small fraction of the world population (Figure 1b). Bringing energy fairness cannot be done through the use of fossil fuels, but via renewable energy that is widely available across the planet. For example, current photovoltaic technology could easily support the energy needs of a population of ten billion.[5] Nonetheless, the large scale deployment of

Dr. Q. Lai, Dr. Y. Sun, T. Wang, P. Modi, Prof. K.-F. Aguey-Zinsou
MERLin

School of Chemical Engineering
The University of New South Wales
Sydney, NSW 2052, Australia
E-mail: f.aguey@unsw.edu.au

C. Cazorla
Material Sciences and Engineering
The University of New South Wales
Sydney, NSW 2052, Australia

Prof. U. B. Demirci
Institut Européen des Membranes
IEM – UMR 5635

CNRS
ENSCM

Univ Montpellier
Montpellier 34000, France

Dr. J. R. Ares Fernandez, Dr. F. Leardini
Laboratorio de Materiales de Interés Energético
Universidad Autonoma de Madrid
Cantoblanco 28049 Madrid, Spain

The ORCID identification number(s) for the author(s) of this article can be found under https://doi.org/10.1002/adsu.201900043.

DOI: 10.1002/adsu.201900043
renewables is still conditioned to the development of an effec-
tive energy carrier. Renewable sources are intermittent and our
current societies rely on 24/7 dispatchable energy ready for use.
This could be achieved through the local production of synthetic
hydrocarbons from renewable energy. However, the overall
efficiency remains low[6] and in the context of a decarbonized
economy, the use of hydrogen as the alternative energy carrier is
a more viable option. Hydrogen has three times the gravimetric
energy density of gasoline. It is abundant in the form of water
and can be used to generate electricity with a fuel cell, mechan-
ical work in combustion engines, or heat through catalytic com-
bustion (Figure 2). The idea of substituting current fossil fuels
with hydrogen is not new[7] but this has been the subject of
many stop and go policies that have left the impression that the
technology is not mature.

**1.2. The Hydrogen Economy: Quo Vadis?**

The drawback in the use of hydrogen as an energy carrier is
often associated with the lack of methods to effectively store it.
Hydrogen has a very low volumetric energy density (0.7 kJ L⁻¹ at
25 °C and atmospheric pressure), and the current mature storage
technology is through hydrogen compression at 70 MPa. This

leads to a volumetric energy density of ≈3 MJ L⁻¹ at the system
level by using composite tanks.[8] These high pressure hydrogen
tanks are already finding applications within the first hydrogen

<figure><img src="image_3.png" /><figcaption>Kondo-François</figcaption></figure>

**Aguey-Zinsou** completed a Ph.D. in heterogeneous catalysis at the University Pierre et Marie Curie (Paris) in 2001. He carried out postdoctoral research at The University of Queensland in bio-electrochemistry. He later moved to the research center Helmholtz Zentrum Geesthacht in Germany,

Queen Mary University London, and University College London where he supervised research projects on hydrogen storage, biofuel cells, and biomaterials. Since 2009, Francois is leading MERLin at UNSW. His research focuses on the development of advanced materials for hydrogen technologies.

<figure><img src="image_4.png" /></figure>

**Figure 1.** Current energy related challenges facing society. a) Evolution of largest glacier in France “La Mer de Glace.” In 1903, the tourist attraction was to walk over the glacier, which was at the level of the Montenvers station (2013 picture). In 2013, the ice was not visible anymore. Subjected to the effects of global warming, the glacier has moved far back. The tension between economic development and the need for new models is emerging across many rich countries as demonstrated by the riots in France in November 2018 (image bottom left); b) electricity consumption normalized per capita and the average human output of 300 W as function of the latitude. Electricity consumptions were sourced from The World Factbook 2016–17 (Washington, DC: Central Intelligence Agency, 2016). A value of 1, a country is consuming the energy that can be produced by a single man or woman. The world current electricity consumption is almost equal to 1, i.e., the earth population, but many countries do not have access to their share of electricity to ensure economic growth. The over consumption in the rich countries is an average that is not reflecting social gaps in access to energy.
<figure><img src="image_5.png" /><figcaption>**Figure 2.** The hydrogen economy. Illustration of the hydrogen renewable energy chain. The difficulty of solid state hydrogen storage remains in the finding of better materials meeting competing requirements. LaNi₅ as a hydrogen storage material (1.28 mass% H₂, and 123 g H₂ L⁻¹) was taken as the reference material. Upon hydrogen absorption heat is released owing to the exothermic nature of the reaction of hydrogen with the metal alloy. The storage materials must meet targets of cost, temperature and pressure for hydrogen absorption and release as well as high energy density, both volumetric and gravimetric. The current cost of hydrogen produced from renewables is estimated at ≈AUD $100-270 per MWh.[568]</figcaption></figure>

vehicles being commercialized. However, when considering stationary applications the use of compressed tanks is not the best option owing to their high cost (≈AUD $685 per kg H₂ at the system level[9]—Figure 3). Metal hydrides and in particular those based on the early alloys developed, e.g., LaNi₅ and TiFe and more recently Ti BCC alloys,[10] provide new and significant opportunities to effectively store renewable energy with a high volumetric density at scale (>10 MJ L⁻¹)[11] and in a cost effective manner. These metals store hydrogen via the simple reaction (1) and thus provide a means to safely store hydrogen within their structure at low hydrogen pressures.

$$M + \frac{x}{2} H_2 \leftrightarrow MH_x + \text{heat} \quad (1)$$

For a long time, further research along these alloys has been discouraged in favor of the development of new materials to meet the targets of sole vehicle application. However, with the uptake of renewable energy at levels over 20% in some countries, the emergence of alternative energy storage solutions has become critical. Intermittent renewable energy like wind and solar have fluctuations in their output which require grid integration with energy storage

<figure><img src="image_6.png" /><figcaption>**Figure 3.** Competitiveness of hydrogen technology for the storage of electricity. a) Cost of electricity per storage technology and b) breakdown cost of hydrogen for stationary energy storage applications using solid state hydrogen storage given. All costs are in AUD/MWh and the following conversion rates were used: 1 USD to 1.36 AUD and 1 EUR to 1.60 AUD. The cost of all the storage technologies besides hydrogen were obtained from the DOE/EPRI report.[569] Prices of individual components of the fuel cell system were obtained from the NREL report[13b] and were adjusted to include the metal hydride prices and production cost of hydrogen from renewable energy technologies. For this calculation, TiMn₁.₉₅ was used as the metal hydride for the solid state storage of hydrogen. High pressure hydrogen storage tanks for vehicles application cost: AUD $550 per kg H₂, which corresponds to a system cost of AUD $685 per kg H₂,[9] while DOE cost target is AUD $456 per kg H₂ for storage system cost.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>**Figure 4.** Price and abundance of elements for designing hydrogen storage materials. Periodic table summarizing the elements forming hydrides, their market price (AUD $ per kg), and abundance. Discrepancies arise when multiple processes are involved in extracting the pure elements, and upon additional market variations. This summary indicates that Na, Mg, K, Ca, Fe, Al, Si and O are the most common elements found in the earth's crust. Fortunately, metal hydrides such as NaBH₄, MgH₂, and AlH₃ with high hydrogen storage capacities can be formed using such elements.[11]</figcaption></figure>

materials to provide the required energy supply (Figure 2).[12] In the case of stationary applications, hydrogen storage technologies provide solutions through the integration of three technologies: water electrolysis, hydrogen storage, and fuel cells for electricity generation. Once stored, the hydrogen can then be transported to the point of use or even exported. Many feasibility studies and demonstration sites have been showcased to validate the benefits of hydrogen technologies for on-site energy storage and generation in stationary applications.[13] However, the understanding of the potential of hydrogen as an energy carrier remains limited. Thus to date, the current technology considered is that of battery with the hope that more mature battery chemistry will be capable of “high density energy storage” at relatively low costs.

Detailed cost analysis of hydrogen technologies integrating electrolyzers, solid state storage, and fuel cells systems reveals that these integrated hydrogen technologies are already extremely cost competitive (Figure 3). In particular, when considering that currently no large scale manufacturing capability exists along the hydrogen chain, the potential for the technology to become readily accessible as cost goes down is immense. Through this analysis it is apparent that the main cost remains at the level of the electrolyzer and fuel cell. The cost of the storage technology is minimal in comparison. As a first approximate, the cost (COA, B) of the raw hydrogen storage material (A, B) can be determined from

$$CO_{A, B_y} = \frac{CO_A \times \frac{x}{x + y} \times N_A + CO_B \times \frac{\gamma}{x + y} \times N_B}{\frac{x}{x + y} \times N_A + \frac{\gamma}{x + y} \times N_B} \quad (2)$$

where COA, COB correspond to the cost of the elements A and B ($ per kg), x and y are the respective amount of the

elements A and B, and N_A and N_B their respective molar mass (g mol⁻¹).

By using the tabulated cost of single elements in Figure 4, the cost of binary to higher quinary alloys with (non)-stoichiometric compositions, can be determined from Equation (2). The cost of typical hydrogen storage alloys is summarized in Table 1. With the cost of room temperature (RT) hydride =AUD$10 kg⁻¹, the cost per kg of hydrogen stored in a typical hydride with a storage capacity of 2 mass% H₂ would be around AUD$500, which is comparable to the cost of hydrogen compressed at 70 MPa. However, owing to the lower pressures involved, solid state hydrogen storage is safer and offers much higher volumetric densities. In the context of competitive markets, the awareness of the potential of solid state hydrogen storage must be raised and practical guidelines must be developed for matching operating hydrogen pressures, flow, and temperatures to facilitate the integration of these hydrides with electrolyzers and fuel cells. At the fundamental level strategies, which this review aims to summarize, the design and operation of hydrogen storage materials must be established to consolidate and strengthen the existing knowledge and facilitate the market penetration of hydrides. The advantages of existing low temperature solid-state hydrogen materials lie in their long cycle life, improved safety compared to compressed hydrogen, and full discharge capability with “no memory” drawbacks. Research is needed for the optimization of their operation, large scale production, safety, maintenance, and recycling.

## 1.3. Beyond Current Materials

Current hydrogen storage alloys with capability to store hydrogen at the ambient have only a low hydrogen storage capacity (typically <2 mass% H₂). The challenge in enabling a
<table><caption>**Table 1.** Common metal alloys (AUD$ kg⁻¹ of material) based on the current market price of individual elements (not including the manufacturing cost of the alloy).</caption><thead><tr><th>Metal alloys</th><th></th><th>Price of the alloys</th></tr></thead><tbody><tr><td rowspan="5">AB<sub>2</sub></td><td>TiMn<sub>1.95</sub></td><td>5.61</td></tr><tr><td>TiMn<sub>1.25</sub>Cr<sub>0.25</sub></td><td>6.95</td></tr><tr><td>Ti<sub>0.9</sub>Zr<sub>0.1</sub>Mn<sub>1.4</sub>Cr<sub>0.45</sub>Fe<sub>0.15</sub></td><td>8.04</td></tr><tr><td>TiCrMn</td><td>8.42</td></tr><tr><td>Ti<sub>1.2</sub>CrMn</td><td>8.59</td></tr><tr><td></td><td>Ti<sub>0.9</sub>Zr<sub>0.2</sub>Mn<sub>1.8</sub>V<sub>0.2</sub>Fe<sub>0.2</sub></td><td>9.19</td></tr><tr><td rowspan="5">AB<sub>5</sub></td><td>La<sub>0.8</sub>Nd<sub>0.2</sub>Ni<sub>5</sub></td><td>23.29</td></tr><tr><td>La<sub>0.8</sub>Y<sub>0.2</sub>Ni<sub>5</sub></td><td>20.76</td></tr><tr><td>LaNi<sub>5</sub></td><td>18.83</td></tr><tr><td>LaNi<sub>4.55</sub>Al<sub>0.45</sub></td><td>18.67</td></tr><tr><td>TiFe<sub>0.7</sub>Mn<sub>0.2</sub></td><td>6.05</td></tr><tr><td rowspan="3">AB</td><td>TiFe<sub>0.88</sub></td><td>5.80</td></tr><tr><td>Ti<sub>24</sub>V<sub>0.45</sub>Cr<sub>31</sub>Mn<sub>10</sub></td><td>10.11</td></tr><tr><td>V<sub>0.37</sub>Ti<sub>0.33</sub>Mn<sub>1.3</sub></td><td>14.44</td></tr><tr><td></td><td>V<sub>0.855</sub>Cr<sub>0.145</sub></td><td>24.70</td></tr><tr><td colspan="3"><strong>Common hydrides</strong></td></tr><tr><td></td><td>NaH</td><td>43<sup>[445]</sup></td></tr><tr><td></td><td>LiH</td><td>237<sup>[445]</sup></td></tr><tr><td></td><td>CaH<sub>2</sub></td><td>142<sup>[446]</sup></td></tr><tr><td></td><td>MgH<sub>2</sub></td><td>—</td></tr><tr><td></td><td>AlH<sub>3</sub></td><td>87</td></tr><tr><td colspan="3"><strong>Common complex hydrides</strong></td></tr><tr><td></td><td>NaAlH<sub>4</sub></td><td>45<sup>[447]</sup></td></tr><tr><td></td><td>LiAlH<sub>4</sub></td><td>355</td></tr><tr><td></td><td>NaBH<sub>4</sub></td><td>140<sup>[446]</sup></td></tr><tr><td></td><td>2LiBH<sub>4</sub>+MgH<sub>2</sub></td><td>818<sup>[445]</sup></td></tr></tbody></table>

wider uptake of hydrogen is in the design of better materials with much higher energy densities (>5 mass% H₂) to effectively displace fossil fuels and enable the transport and distribution of renewable energy in a very compact form including vehicles application. Nonmetals like carbon microporous materials and highly porous structures like metal-organic frameworks (MOFs), covalent-organic framework (COFs) can also be used to store molecular hydrogen via physisorption with rapid kinetics (Figure 5b).[14] However, in this case the adsorption capacity is often too low, typically ≈0.24 mass% H₂ for carbons[15] and <2 mass% H₂ for MOFs;[16] although Morris et al. recently claimed a reversible excess sorption of 10.5 mass% H₂ at the ambient in a manganese-based MOF.[17] Liquid organic carriers such as those based on NH₃ and cycloalkanes are also potential alternatives and their potential as hydrogen carrier has been summarized in several excellent reviews.[18] However, processes based on organic carriers are often energy intensive and should be improved in terms of catalyst stability, activity, and reversible hydrogen storage capacity. This review does not cover recent progress on materials for hydrogen physisorption or liquid organic carriers, but rather focuses on materials where

hydrogen is stored via chemisorption in the form of hydrides. Various elements can effectively form hydrides of extremely high hydrogen densities (Figures 4 and 5),[19] and thus have the potential to deliver a solution to store hydrogen beyond current capabilities. For example, light alkali and alkaline earth metals form ionic hydrides such as LiH and NaH,[20] where hydrogen exists as a hydride (H⁻). Compounds including LiNH₂ and NH₃[21] form high hydrogen capacity covalent hydrides, where the hydrogen is partially charged positively (H⁺). In contrast, in complex hydrides such as NaAlH₄ and LiBH₄, the hydrogen exists, again with high density, as a negatively charged H⁺ (Table 2). In metals such as Pd and Ti forming interstitial hydrides, the hydrogen bonding is considered to be metallic and these elements in addition to their hydrogen storage properties can provide a solution for hydrogen catalysis in the high hydrogen density hydrides. The variety of existing hydrogen bonds can be viewed as an opportunity of finding ways to modify the properties of existing hydrides toward specific hydrogen storage applications. For example, the combination of H⁺ and H⁻ is a means to facilitate the release of hydrogen from mixed compounds involving hydride and protide owing to the more favorable thermodynamics.[22]

To date, most of these high capacity hydrides are not suitable in their “raw” form for hydrogen storage because they only release hydrogen at temperatures way above the ambient (Table 2); and often the hydrogen release is not reversible. Furthermore, the release or uptake of hydrogen in these hydrides can lead to significant volume changes, which does not simplify their integration into real systems. For example, in a typical metal hydride, hydrogen desorption/absorption can lead to a significant volume expansion owing to the (de)insertion of the hydrogen atoms in the crystalline lattice and structural reorganization of the atomic order.[8] In particular, in intermetallic compounds these volume changes are about 1–3 cm³ mol⁻¹ H.[23] In saline hydrides like Mg/MgH₂ the expansion is of 30%. This can be overcome by leaving some space in the tank for expansion/contraction but it affects the overall hydrogen storage capacity at the system level. Complex hydrides also undergo various decomposition pathways and changes in volume, which depend upon the type of the reaction and the extent of melting/foaming of the material. For example, H₃BNH₃ foams upon thermal heating as the result of the large amount of gases released in addition to hydrogen.[24] NaAlH₄ decomposes into Na₃AlH₆ and Al with 30.3% contraction in the first stage of hydrogen release, but in the second decomposition stage of Na₃AlH₆ into NaH and Al a 13.8% expansion is observed.[25]

The current challenge in the use of high hydrogen capacity hydrides is thus in the conception of novel fundamental approaches to effectively control their properties and to simplify their behavior upon hydrogen uptake/release for facile hydrogen sorption near the ambient. This also involves taking into account important parameters including: i) the ease of activation of the material for the initial hydrogen uptake/release, ii) the rate of hydrogen uptake/release, iii) the heat of formation of the hydride and the management of that heat, iv) the hydrogen cycling stability of the material, v) the resistance of the hydride to hydrogen gas impurities, vi) the safety associated with the use of the hydride, and vii) the final cost of the material.[26]
<figure><img src="image_8.png" /><figcaption>**Figure 5.** Current potential solid states hydrogen storage materials. a) Energy density and hydrogen desorption temperature of selected hydrogen storage materials versus the US Department of Energy (DOE) targets;[570] and b) typical methods to store hydrogen and associated hydrogen binding strength. For practical applications a relatively weak hydrogen bonding strength is needed to reversibly store hydrogen with a minimum energy input.</figcaption></figure>

Herein, we review current state-of-the-art in the design of hydrogen storage materials and the fundamental understanding of hydrogen sorption in these materials with the aim of facilitating a rationalization of current approaches. A particular emphasis is made on current fundamental understanding of hydrogen behavior in hydrides across materials types, from metal alloys to complex hydrides; across length scales, from several micrometers to the nanoscale; and across phases from nano-structured to amorphous as well as methods to load/unload hydrogen from these materials. Synthesis approaches to reach certain hydrogen storage properties are also reviewed as well as methodologies for modeling guidance and machine learning techniques for materials' discovery and prediction. Additional details of these methodologies can be found in books and a range of excellent previous review articles.[8,27] The difficulties with the engineering steps associated with the insertion and hydrogen cycling of hydrides in tanks and strategies to overcome these difficulties are also discussed. In particular, the importance of integrating engineering aspects at the very early stages of hydride materials synthesis and design is highlighted. The performances of solid state hydrogen tanks are directly related to the initial hydrogen storage material design.

## 2. How is Hydrogen Practically Loaded into and Released from Materials?

Due to the nature of the reaction involving hydrogen uptake and release in materials (reaction (1)), the most common method to drive hydrogen in and out of materials has been by

tuning temperature and pressure (Figure 6a). Hence at a given temperature, by applying a hydrogen pressure above the equilibrium pressure $P^{eq}$ of the hydride phase (see Section 3.2), hydrogen can be loaded in materials. In contrast, by applying a pressure below that of $P^{eq}$, hydrogen is released from the material. One of the drawbacks of such an approach is that under operating conditions of hydrogen absorption/desorption, $P^{eq}$, which determines the pressure at which hydrogen is released or absorbed by the material, is never at a constant value of pressure but is always increasing/decreasing because of the heat generated by the reaction of hydrogen with the sorption material. Accordingly without appropriate heat management, hydrogen cannot simply be delivered at a given rate. In this case, the hydride must be heated to notably raise $P^{eq}$, or a hydrogen pressure significantly lower than $P^{eq}$ must be applied, to sustain a constant flow of hydrogen and compensate for the heat consumed by the release of hydrogen from the hydride. The reverse applies for the uptake of hydrogen and in this case, heat must be removed from the hydride bed to facilitate a rapid hydrogen uptake despite the exothermic nature of hydrides formation.

For investigating the hydrogen sorption properties of small samples (≈mg to g) this is not a major issue; however, for large scale systems involving kg of hydrides, heat management becomes critical and external systems to drive heat out or in must be added to the hydride storage tank (see Section 10). This not only adds further complications in terms of system design but also significantly impedes the practical hydrogen storage capacity at the system level. However, because of its simplicity, the use of temperature and pressure as external control of the
<table><caption>**Table 2.** Summary of the hydrogen gravimetric density, temperature of decomposition (T<sub>dec</sub>), enthalpy of formation (ΔH<sub>f</sub>), enthalpy of decomposition (ΔH<sub>dec</sub>) and entropy (ΔS) of selected hydrides. The hydrides are classified as a function of the type of hydrogen bonding.</caption><thead><tr><th></th><th>H<sub>2</sub> density [mass%]</th><th>T<sub>dec</sub> [°C]</th><th>ΔH<sub>f</sub> [kJ mol<sup>-1</sup>]</th><th>-ΔH<sub>dec</sub> [kJ mol<sup>-1</sup> H<sub>2</sub>]</th><th>ΔS [J K<sup>-1</sup> mol<sup>-1</sup> H<sub>2</sub>] or *[J K<sup>-1</sup> mol<sup>-1</sup>]</th></tr></thead><tbody><tr><td colspan="6"><strong>Protides</strong></td></tr><tr><td>LiBH<sub>4</sub></td><td>18.5</td><td>380<sup>[448]</sup></td><td>-146,<sup>[449]</sup> -177<sup>[450]</sup><br/>-184.7,<sup>[451]</sup> -190.46<sup>[452]</sup><br/>-194.44<sup>[118a,450,453]</sup></td><td>-95.8<sup>[22]</sup></td><td>*75.9,<sup>[450]</sup> *128.7<sup>[454]</sup></td></tr><tr><td>NaBH<sub>4</sub></td><td>10.7</td><td>400<sup>[455]</sup></td><td>-182~ -188<sup>[449]</sup><br/>-183.3,<sup>[454]</sup> -191<sup>[118a]</sup></td><td>-94.3<sup>[22]</sup></td><td>*101.3, *126.2<sup>[454]</sup></td></tr><tr><td>KBH<sub>4</sub></td><td>7.5</td><td>500<sup>[455]</sup></td><td>-229,<sup>[118a]</sup> -228<sup>[449]</sup><br/>-242.3<sup>[454]</sup></td><td>-120<sup>[454]</sup></td><td>*106, *162<sup>[454]</sup></td></tr><tr><td>Mg(BH<sub>4</sub>)<sub>2</sub></td><td>14.9</td><td>323<sup>[22]</sup></td><td>-99<sup>[449]</sup></td><td>-24.8<sup>[449]</sup></td><td>*91.3,<sup>[456]</sup> *128<sup>[457]</sup></td></tr><tr><td>Ca(BH<sub>4</sub>)<sub>2</sub></td><td>11.6</td><td>360, 260</td><td>-151<sup>[449]</sup></td><td>-37.8<sup>[449]</sup></td><td>158<sup>[458]</sup></td></tr><tr><td>Be(BH<sub>4</sub>)<sub>2</sub></td><td>20.5</td><td>25<sup>[22]</sup></td><td>27<sup>[118a]</sup></td><td>7</td><td>—</td></tr><tr><td>Al(BH<sub>4</sub>)<sub>3</sub></td><td>16.9</td><td>150<sup>[459]</sup></td><td>-301.8<sup>[454]</sup></td><td>-50.3<sup>[454]</sup></td><td>*289<sup>[454]</sup></td></tr><tr><td>Zn(BH<sub>4</sub>)<sub>3</sub></td><td>8.5</td><td>85<sup>[22]</sup></td><td>-16,<sup>[22]</sup> -18<sup>[449]</sup></td><td>-4<sup>[22]</sup></td><td>—</td></tr><tr><td>Ti(BH<sub>4</sub>)<sub>3</sub></td><td>13.1</td><td>25<sup>[22]</sup></td><td>—</td><td>—</td><td>—</td></tr><tr><td>Zr(BH<sub>4</sub>)<sub>4</sub></td><td>10.7</td><td>250<sup>[22]</sup></td><td>—</td><td>-10.9<sup>[449]</sup></td><td>—</td></tr><tr><td>H<sub>3</sub>BNH<sub>3</sub></td><td>19.6</td><td>100, 150, 500<sup>[460]</sup></td><td>-21<sup>[461]</sup></td><td>-7<sup>[461]</sup></td><td>—</td></tr><tr><td>LiAlH<sub>4</sub></td><td>10.5</td><td>180<sup>[22]</sup></td><td>—</td><td>45.4<sup>[22]</sup></td><td>*78.8<sup>[462]</sup></td></tr><tr><td>Li<sub>3</sub>AlH<sub>6</sub></td><td>11.2</td><td>220<sup>[22]</sup></td><td>—</td><td>-76<sup>[22]</sup></td><td>*102.6<sup>[462]</sup></td></tr><tr><td>NaAlH<sub>4</sub></td><td>7.4</td><td>247<sup>[22]</sup></td><td>—</td><td>-120<sup>[22]</sup><br/>-40<sup>[463]</sup></td><td>*89.7<sup>[462]</sup></td></tr><tr><td>Na<sub>3</sub>AlH<sub>6</sub></td><td>5.9</td><td>268<sup>[22]</sup></td><td>—</td><td>-178<sup>[22]</sup></td><td>*161.5<sup>[462]</sup></td></tr><tr><td>KAlH<sub>4</sub></td><td>5.7</td><td>300<sup>[22]</sup></td><td>—</td><td>-207.4<sup>[22]</sup></td><td>—</td></tr><tr><td>K<sub>3</sub>AlH<sub>6</sub></td><td>4.0</td><td>330<sup>[22]</sup></td><td>—</td><td>-270.8<sup>[22]</sup></td><td>—</td></tr><tr><td colspan="6"><strong>Hydrides</strong></td></tr><tr><td>LiNH<sub>2</sub></td><td>8.7</td><td>380</td><td>-186.3<sup>[464]</sup></td><td>-186.3<sup>[464]</sup></td><td>—</td></tr><tr><td>NaNH<sub>2</sub></td><td>5.1</td><td>400<sup>[465]</sup></td><td>-118.8</td><td>-118.8</td><td>*76.9</td></tr><tr><td>Mg(NH<sub>2</sub>)<sub>2</sub></td><td>7.1</td><td>350</td><td>-374<sup>[464]</sup></td><td>-187<sup>[464]</sup></td><td>*52<sup>[464]</sup></td></tr><tr><td colspan="6"><strong>Ionic hydrides</strong></td></tr><tr><td>LiH</td><td>12.5</td><td>720<sup>[22]</sup></td><td>-116.3<sup>[22]</sup></td><td>-232.6<sup>[22]</sup></td><td>-134<sup>[27m]</sup></td></tr><tr><td>NaH</td><td>4.2</td><td>425<sup>[22]</sup></td><td>-56.5<sup>[22]</sup></td><td>-113<sup>[22]</sup></td><td>-162<sup>[27m]</sup></td></tr><tr><td>KH</td><td>2.5</td><td>417<sup>[22]</sup></td><td>-57.7<sup>[22]</sup></td><td>-115.4<sup>[22]</sup></td><td>-168<sup>[27m]</sup></td></tr><tr><td>MgH<sub>2</sub></td><td>7.6</td><td>327<sup>[22]</sup></td><td>-75.3<sup>[22]</sup></td><td>-75.3<sup>[22]</sup></td><td>-134<sup>[27m]</sup></td></tr><tr><td>CaH<sub>2</sub></td><td>4.8</td><td>600<sup>[22]</sup></td><td>-181.5<sup>[22]</sup></td><td>-181.5<sup>[22]</sup></td><td>-140<sup>[27m]</sup></td></tr><tr><td>RbH</td><td>1.2</td><td>170<sup>[22]</sup></td><td>-52.3<sup>[22]</sup></td><td>-104.6<sup>[22]</sup></td><td>-170<sup>[27m]</sup></td></tr><tr><td>CsH</td><td>0.75</td><td>170<sup>[22]</sup></td><td>-54.2<sup>[22]</sup></td><td>-108.4<sup>[22]</sup></td><td>-170<sup>[27m]</sup></td></tr><tr><td>SrH<sub>2</sub></td><td>2.2</td><td>675<sup>[22]</sup></td><td>-180.3<sup>[22]</sup></td><td>-180.3<sup>[22]</sup></td><td>-156<sup>[27m]</sup></td></tr><tr><td>BaH<sub>2</sub></td><td>1.4</td><td>—</td><td>-177<sup>[22]</sup></td><td>-177<sup>[22]</sup></td><td>-144<sup>[27m]</sup></td></tr><tr><td colspan="6"><strong>Covalent hydrides</strong></td></tr><tr><td>B<sub>2</sub>H<sub>6</sub></td><td>21.7</td><td>40<sup>[22]</sup></td><td>36.4<sup>[22]</sup></td><td>12.1<sup>[22]</sup></td><td>—</td></tr><tr><td>GaH<sub>3</sub></td><td>4.1</td><td>-15<sup>[22]</sup></td><td>118<sup>[22]</sup></td><td>25<sup>[22]</sup><br/>78.7<sup>[22]</sup></td><td>—</td></tr><tr><td>BiH<sub>3</sub></td><td>1.4</td><td>-40<sup>[22]</sup></td><td>217.6<sup>[22]</sup></td><td>32,<sup>[22]</sup> 144.6<sup>[22]</sup></td><td>—</td></tr><tr><td colspan="6"><strong>Metallic hydrides</strong></td></tr><tr><td>LaH<sub>2</sub></td><td>1.4</td><td>600<sup>[27m]</sup></td><td>-210<sup>[27m]</sup></td><td>-210<sup>[27m]</sup></td><td>-152<sup>[27m]</sup></td></tr><tr><td>YH<sub>2</sub></td><td>3.2</td><td>600<sup>[27m]</sup></td><td>—</td><td>-225<sup>[27m]</sup></td><td>-144<sup>[27m]</sup></td></tr><tr><td>TiH<sub>2</sub></td><td>4</td><td>380<sup>[22]</sup></td><td>-126<sup>[27m]</sup></td><td>-126<sup>[27m]</sup></td><td>-100<sup>[27m]</sup></td></tr><tr><td>VH<sub>2</sub></td><td>3.8</td><td>35<sup>[22]</sup></td><td>-54<sup>[27m]</sup></td><td>10.8,<sup>[22]</sup> -54</td><td>-108<sup>[27m]</sup></td></tr><tr><td>PdH<sub>0.5</sub></td><td>0.47</td><td>-78<sup>[27m]</sup></td><td>—</td><td>-40<sup>[27m]</sup></td><td>-84<sup>[27m]</sup></td></tr></tbody></table>
<table><caption>**Table 2. Continued.**</caption><thead><tr><th></th><th>H<sub>2</sub> density [mass%]</th><th>T<sub>dec</sub> [°C]</th><th>ΔH°<sub>f</sub> [kJ mol<sup>-1</sup>]</th><th>-ΔH<sub>dec</sub> [kJ mol<sup>-1</sup> H<sub>2</sub>]</th><th>ΔS [J K<sup>-1</sup> mol<sup>-1</sup> H<sub>2</sub>] or *[J K<sup>-1</sup> mol<sup>-1</sup>]</th></tr></thead><tbody><tr><td colspan="6"><strong>Intermediate</strong></td></tr><tr><td>AlH<sub>3</sub></td><td>10</td><td>150<sup>[22]</sup><br>25<sup>[27m]</sup></td><td>-46<sup>[22]</sup></td><td>-30.6<sup>[22]</sup><br>-8<sup>[27m]</sup></td><td>-</td></tr><tr><td>BeH<sub>2</sub></td><td>18.2</td><td>250<sup>[22]</sup></td><td>-</td><td>-18.9<sup>[22]</sup></td><td>-</td></tr><tr><td>ZnH<sub>2</sub></td><td>3</td><td>90<sup>[22]</sup></td><td>-</td><td>-5.4<sup>[22]</sup></td><td>-</td></tr><tr><td>CdH<sub>2</sub></td><td>1.75</td><td>-20<sup>[22]</sup></td><td>-</td><td>26.4<sup>[22]</sup></td><td>-</td></tr><tr><td colspan="6"><strong>Unknown</strong></td></tr><tr><td>TiH<sub>3</sub></td><td>1.4</td><td>-95<sup>[22]</sup></td><td>245<sup>[22]</sup></td><td>164<sup>[22]</sup></td><td>-</td></tr><tr><td>CoH<sub>0.5</sub></td><td>0.84</td><td>25</td><td>-</td><td>0<sup>[27m]</sup><br>30</td><td>-</td></tr><tr><td>HgH<sub>2</sub></td><td>1</td><td>-125<sup>[22]</sup></td><td>-</td><td>55.4<sup>[22]</sup></td><td>-</td></tr><tr><td>NiH<sub>0.5</sub></td><td>0.85</td><td>20<sup>[27m]</sup></td><td>-</td><td>-6<sup>[27m]</sup><br>-58</td><td>-116<sup>[27m]</sup></td></tr><tr><td>FeH</td><td>1.76</td><td>-</td><td>-</td><td>14<sup>[27m]</sup></td><td>-</td></tr><tr><td>MnH</td><td>1.79</td><td>450-730<sup>[27m]</sup></td><td>-</td><td>-9<sup>[27m]</sup><br>-22</td><td>-116<sup>[27m]</sup></td></tr><tr><td>PtH</td><td>0.5</td><td>-</td><td>-</td><td>26<sup>[27m]</sup></td><td>-</td></tr><tr><td colspan="6"><strong>Intermetallics</strong></td></tr><tr><td>CaNi<sub>5</sub></td><td>0.99</td><td>43<sup>[26]</sup></td><td>-</td><td>31.9<sup>[26]</sup></td><td>101<sup>[26]</sup></td></tr><tr><td>LaNi<sub>5</sub></td><td>1.28</td><td>12<sup>[26]</sup></td><td>-</td><td>30.8<sup>[26]</sup></td><td>108<sup>[26]</sup></td></tr><tr><td>TiMn<sub>1.5</sub></td><td>1.15</td><td>-21<sup>[26]</sup></td><td>-</td><td>28.7<sup>[26]</sup></td><td>114<sup>[26]</sup></td></tr><tr><td>TiFe</td><td>1.5</td><td>-8<sup>[26]</sup></td><td>-</td><td>28.1<sup>[26]</sup></td><td>106<sup>[26]</sup></td></tr><tr><td>MmNi<sub>5</sub></td><td>1.24</td><td>-56<sup>[26]</sup></td><td>-</td><td>21.1<sup>[26]</sup></td><td>97<sup>[26]</sup></td></tr><tr><td>Mg<sub>2</sub>Ni</td><td>3.3</td><td>255<sup>[26]</sup></td><td>-</td><td>-64.5<sup>[26]</sup></td><td>-122<sup>[26]</sup></td></tr><tr><td>(V<sub>0.9</sub>Ti<sub>0.1</sub>)<sub>0.95</sub>Fe<sub>0.05</sub></td><td>1.8</td><td>36<sup>[26]</sup></td><td>-</td><td>-43<sup>[26]</sup></td><td>-140<sup>[26]</sup></td></tr></tbody></table>

hydrogen reaction with hydrides is still currently the dominant approach.

The possibility of using other means to drive the release/uptake from solid state hydrogen materials has been investigated, but apart from the electrochemical methods, all the other approaches (e.g., microwave, laser, and ultrasound irradiation) also involve the direct or indirect generation of heat (Figure 6a). For example, the use of microwave radiation has been investigated by several groups because it has the potential to provide rapid, energy-efficient heating. Heating by microwave involves electromagnetic radiations in the frequency range of 300 MHz to 300 GHz.[28] The interaction of charged particles within dielectric materials with the electric field component of electromagnetic radiations causes the materials to heat up due to the Maxwell-Wagner effect.[29] Several results have been reported on the possibility of using microwave to drive hydrogen desorption from metal hydrides and complex hydrides.[30] For example, hydrogen could be released from TiH₂ by microwave irradiation but the release was incomplete because of low penetration depth of the microwave irradiation in the TiH₂ particles.[30a] Transition metal hydrides, such as VH<sub>0.81</sub>, ZrH₂, and LaH<sub>2.48</sub>, also showed a rapid temperature increase upon microwave irradiation with some hydrogen release.[30c]

By using ultrasonic irradiation, it is also possible to generate shock waves and localized heating by acoustic cavitation.

Ares et al. reported the possibility of releasing hydrogen from MgH₂ through this principle.[31] The idea of using a laser was also tested. For example, starting with Ti in hydrogen atmosphere, the formation of TiH₂ was reported upon laser irradiation at 4 J cm⁻².[32] Investigations of laser heating of MgH₂ and Mg₂NiH₄ + MgH₂ were also conducted by Plasma Kinetics Corporation. It was reported that ≈80 continuous watts were needed to release about 4.5 kg of hydrogen from 119 kg of MgH₂ at rates of ≈0.9 kg h⁻¹.[33] A temperature of 390 °C was reached within 1 ms to enable the decomposition of the hydride.

Exposure of hydride materials to UV radiation was also proposed decades ago as a means to release hydrogen from hydrides, including MgH₂,[34] LiAlH₄,[35] and AlH₃,[36] but this led to little amounts of hydrogen. For example, upon UV irradiation with an intensity of 280 mW cm⁻², 100 mg of MgH₂ powder released 0.008 mass% H₂, at room temperature.[37] The mechanism for hydrogen release was proposed to be associated with an electronic charge and/or energy transfer process involving the creation of doubly excited sites.[34b] Based on the similar idea of using light to drive hydrogen absorption/release, Sun et al. recently proposed an alternative method, by using the plasmonic effect of Au nanoparticles to generate local heat at the surface of hydride particles and thus locally trigger hydrogen absorption/desorption.[38] It is expected that by such an approach, the localized generation
<figure><img src="image_9.png" /><figcaption>**Figure 6.** Principles of methods to store hydrogen. a) Illustration of existing methods to facilitate the loading and unloading of hydrogen in materials; b) principles of electrochemical hydrogen storage; and c) the various morphologies of hydrogen storage materials. P₁ and P₂ correspond to the equilibrium plateau pressure (P<sup>eq</sup>) at the temperature T₁ and T₂. Materials are usually heated to raise P<sup>eq</sup> and facilitate the release of hydrogen above atmospheric pressure to feed a fuel cell, for example.</figcaption></figure>

of heat can help to minimize the need of additional external heat management systems. This is highly desirable because it can lead to practical systems of much higher hydrogen storage capacity than currently feasible. However, to date there are very few investigations on this topic<sup>[39]</sup> and many concern the use of plasmonic effect toward the design of hydrogen sensors.<sup>[40]</sup>

Electrochemical charging/discharging of hydrides is another approach, but this is mainly applicable to hydrides with high resistance to alkaline corrosion. Accordingly, the application of this concept to hydrides of high hydrogen storage capacity such as Mg/MgH₂ has been limited.<sup>[41]</sup> Electrochemical charging/discharging relies on the oxidation and reduction of a strong alkaline electrolyte to generate H atoms that are directly absorbed by a metal alloy (M) forming a hydride (MH) (Figure 6b). The charging process follows the general reaction

$$M + nH_2O + ne^- \leftrightarrow MH_n + nOH^- \quad (3)$$

and the reverse reaction applies during discharging. During this process, a nickel counter electrode swinging between its oxide hydroxide and hydroxide is used to balance (4)

$$Ni(OH)_2 + nOH^- \leftrightarrow NiOOH + nH_2O + ne^- \quad (4)$$

It should be noted that reaction (3) is equivalent to the reaction observed upon direct use of temperature and pressure (1). Accordingly both processes are comparable and the

gas equilibrium pressure P<sup>eq</sup> and electrochemical equilibrium potential E<sup>eq</sup> are linked by the relationship<sup>[42]</sup>

$$E^{\text{eq}} = \frac{RT}{nF} \ln P^{\text{eq}} \quad (5)$$

where n corresponds to the number of electrons involved, R is the universal gas constant (8.314 J mol<sup>-1</sup> K<sup>-1</sup>), and F is the Faraday constant (96 500 C mol<sup>-1</sup>).

However, P<sup>eq</sup> in excess of 0.1 MPa is not desirable to avoid a significant self-discharge of the metal alloy at the ambient and the excess formation of hydrogen through the secondary reaction

$$H_2O + e^- \rightarrow OH^- + 1/2H_2 \quad (6)$$

A P<sup>eq</sup> comprised between 0.001 and 0.1 MPa at 25 °C is considered as optimal. In the case of the archetypical LaNi₅ alloy at the basis of Ni–MH commercial battery, this means that its initial P<sup>eq</sup> of 0.23 MPa at 25 °C had to be modified through the substitution of A and B.<sup>[42b]</sup> For a one electron process, the electrochemical charging capacity (C<sub>el</sub> in mAh g<sup>-1</sup>) is also linked to the solid–gas charging capacity (C<sub>gs</sub>, i.e., the amount of hydrogen absorbed per mol of metal alloy) by

$$C_{\text{el}} = \frac{C_{\text{gs}}F}{3.6M} \quad (7)$$

where M is the molecular weight of the alloy.
<figure><img src="image_10.png" /><figcaption>**Figure 7.** Thermodynamic of a range selected hydrides. Evolution of the enthalpy of decomposition (−ΔH°dec) of a) common metal hydrides and alanates, and b) selected borohydrides, amines and intermetallics as a function of their decomposition temperature (Tdec) for hydrogen release. The hydrides with a reasonable hydrogen storage capacity decompose at high temperatures.</figcaption></figure>

It should also be noted that the effectiveness of the method used to load/unload hydrogen in hydrides depends on the particle morphology and the initial state of the hydride particle surface (Figure 6c).

### 3. Hydrogen Storage Properties of Materials: Fundamental Principles

Today, most of the hydride materials of high hydrogen gravimetric density, e.g., MgH₂, LiH, LiAlH₄, LiBH₄, and H₃BNH₃, have hydrogen storage properties unsuitable for application, i.e., temperatures and/or pressures too high for hydrogen absorption/desorption cycling with often very poor hydrogen rates of uptake and release to meet the regimes of an electrolyzer or a fuel cell (Table 2 and Figures 5 and 7). Enabling hydrogen storage under moderate conditions of hydrogen pressure and temperature requires the modification of the heat of reaction of these materials with hydrogen as well as their rate of reaction. This necessitates a better understanding of the interaction of hydrogen with hydride forming elements and a weakening of M–H bonds and intrahydrogen pairing (H–M–H) in these elements to improve the thermodynamics of the reaction as well as the energy barriers controlling the rate of the hydrogen uptake/release. In the context of stationary applications, additional development of such knowledge, along existing low temperature hydrides, e.g., FeTi, Ti BCC (Table 2), has also become critical. In addition to a better understanding of the interaction of hydrogen with alloys, the effects of compositional adjustment on the ease of their activation and operating conditions must be progressed to effectively facilitate their application in the field.

#### 3.1. Current Understanding of the Basic Rules Guiding the Rate of Hydrogen Uptake/Release

The rate of reaction of hydrogen with absorbing materials first of all depends upon the probability of collision of hydrogen molecules in the gas phase with the material’s surface (Figure 8a). In the gas phase, hydrogen molecules have an average thermal

velocity of ≈1775 m s⁻¹ at 26 °C.[20b] Part of the initial kinetic energy is released upon interaction with the surface of the absorbing material and at speed of ≈18 Å ps⁻¹, the van der Waals (vdW) interaction comes into play to physisorb hydrogen (Figure 8a).[20b] The initial rate of reaction thus depends upon the material’s available surface area and any additional activation barriers for the chemisorption of hydrogen. In particular, a high surface area is important when considering the storage of hydrogen through physisorption processes on mesoporous carbons or MOFs, for example.[15,43] Usually the physisorption of hydrogen involves a low energy barrier of ≈1–10 kJ mol⁻¹ and thus is readily achievable at low temperatures.[27e,44] For example, hydrogen physisorption is observed in many porous materials including silicas, aluminas, zeolites, graphites, activated carbons, and carbon nanofibers at –196 °C.[45] For materials such as Mg absorbing hydrogen, a high surface area is also very important as a higher number of available surface sites help to improve the rate of hydrogen chemisorption and thus the overall reaction. For example for Mg, which is notorious for the high hydrogen dissociation barrier (≈432 kJ mol⁻¹ H₂)[44] at its surface, a reduction of the overall activation energy (Eₐ) from 156 to 120 kJ mol⁻¹ has been reported upon a tenfold increase in the specific surface area of Mg particles.[46] The absorption of hydrogen until the formation of the hydride phase is also controlled by many additional barriers involving the penetration of hydrogen into the subsurface and then further diffusion through the material (Figure 8).[27e] Indeed, assuming the formation of an hydride layer at the surface of the material upon the absorption of the first hydrogen molecules, the diffusion of hydrogen through the hydride phase becomes slower (e.g., 1.5 × 10⁻¹⁶ m² s⁻¹ in MgH₂, instead of 1 × 10⁻⁸ m² s⁻¹ in Mg at 350 °C).[44,47]

The rate (k) of the hydrogenation/dehydrogenation reactions depends on the activation energy barrier (Eₐ) following the Arrhenius equation

$$k(T) = Ae^{\frac{-E_a}{RT}} \qquad (8)$$

In this equation, A is a pre-exponential factor. For a given reaction, the reaction rate k at a given temperature depends on
<figure><img src="image_11.png" /><figcaption>**Figure 8.** Illustration of the hydrogen absorption/desorption principles in metals. a) Lennard-Jones potential diagram corresponding to the successive energy barriers encountered by hydrogen during absorption/desorption in a metal. *E*<sub>phys</sub>: energy for hydrogen physisorption, *E*<sub>dis</sub>: energy for hydrogen dissociation, *E*<sub>chem</sub>: energy for hydrogen chemisorption, *E*<sub>pen</sub>: energy for hydrogen penetration in the subsurface, *E*<sub>dif</sub>: energy for hydrogen diffusion in the bulk, and *E*<sub>nuc</sub>: energy for the nucleation and growth of the hydride; b) growth of the hydride phase upon hydrogen absorption through a shrinking core; c) hydrogen diffusion in metal leading to an expansion of the lattice structure; and d) examples of models relevant to hydrogen diffusion in metals. *D* is the diffusion coefficient, *a* the lattice parameter, *S*<sub>1</sub> the elementary jump distance, *d*<sub>j</sub> is the jump distance, *J*(*m*) is the tunneling-matrix element, *m*<sub>H</sub> the mass of hydrogen, *E*<sub>a</sub> the activation energy, and *ν*<sub>H</sub> the vibration frequency of an hydrogen atom (*ν*<sub>H</sub> ~ *m*<sub>H</sub><sup>1/2</sup>). For *S*<sub>j</sub> > 0, the hopping is incoherent. The thermally activated tunneling is considered to be non-adiabatic; which is the case for most hydrides. Reproduced with permission.[27m] Copyright 2005, Springer.</figcaption></figure>

both the pre-exponential factor (A) and the *E*<sub>a</sub> of the reaction process, i.e., a higher pre-exponential factor (A) and a lower value of *E*<sub>a</sub> favor faster reaction rates.

The physical meaning of A remains elusive.[48] Although $A=e^R$ was proposed by Wert and Zener in 1949 for diffusion processes,[49] this was found to be inconsistent with many numerical results. A is more often regarded as an empirical factor that can be affected by the reactant states, e.g., particle size and defective sites, and the reaction conditions, e.g., microwave aid.[50] The Meyer–Neldel rule, which describes the compensation correlation between thermodynamics or kinetics factors,[48] has also been used to explain the relationship between A and *E*<sub>a</sub>, and their compensation effects observed both experimentally and computationally.[50b,51] Nevertheless, predictions based on the Meyer–Neldel rule were found to lack sufficient levels of accuracy to enable the reliable determination of the pre-exponential factor from the activation energy.[48] It is noteworthy that in the case of hydrogen storage materials, such a correlation has barely been explored owing to the difficulty of separating individual energy barriers (Figure 8a). Typical efforts have mainly focused on lowering the overall energy barrier (*E*<sub>a</sub>) by facilitating the uptake and release of hydrogen along what is believed to be the slowest reaction step. For example, in the case of the Mg/MgH<sub>2</sub> system, Mg is often doped with Pd, Nb, or V[27e,52] to recombine/split hydrogen atoms at the metallic

surface of the storage material, because this is believed to be the slowest step. It is possible that defects and grain boundaries introduced in hydride materials are positive factors enhancing the pre-exponential factor for their reaction with hydrogen.[53] However, the interaction of hydrogen with solid materials includes several barriers that may become the rate-limiting steps as the reaction progresses.

Various models including the Jander[54] and Johnson–Mehl–Avrami–Kolmogorov (JMAK)[55] models have been used to better understand the rate-limiting steps involved in hydrogen sorption,[56] with the assumption that individual particles are of similar size, shape, surface reactivity, and undergo similar growth regimes (e.g., are defect free). These models summarized in Figure 9 include the “shrinking core” models which assume that the decomposition starts from the surface and proceeds inward,[57] and the “nucleation and growth” models based on the assumption of random nucleation and growth process within the bulk of the material.[55d] However, it should be noted that the interpretations based on these models are highly subjective to the quality of the fit. Accordingly, these models are often not conclusive owing to the dynamic multistep reaction paths of hydrogen in materials. For example in the case of Mg/MgH<sub>2</sub>, Liang et al. concluded on a 2D growth model as the best fit for the rate-limiting step at 300 °C for MgH<sub>2</sub> doped with transition metals.[58] In contrast, Hanada et al. proposed that the
<figure><img src="image_12.png" /><figcaption>**Figure 9.** Classical rate-limiting step models. Despite their “simplicity” these models have been found to be useful in understanding the rate-limiting steps of hydrogen sorption in metals and complex hydrides.</figcaption></figure>

hydrogen desorption reaction from Ni catalyzed MgH₂ is of first order.[⁵⁹] In another example, Pourabdoli et al. showed that for MgH₂-10 wt% (9Ni–2Mg–Y) the difference in the fit correlation

of the hydrogen desorption rate between the 2D and 3D JMAK model is about 1%.[⁶⁰]

The hydrogen kinetics of complex hydrides are somewhat more complicated because of the multidecomposition steps and reaction intermediates involved, the disproportionation of the initial complex hydride upon hydrogen release, and the resulting phase segregation (Figure 10).[⁶¹] In this case, the rate-limiting step of each individual reaction step is often modeled by using the classical models summarized in Figure 9.[⁶²] However, in addition to the rate-limiting steps associated with the adsorption and diffusion of hydrogen within the material, mass transfer leading to elemental dissociation or recombination upon hydrogen absorption must also be considered. An archetypical example of such complexity is the Ti-doped NaAlH₄ system for which the exact reaction paths and rate-limiting steps upon Ti doping remain unclear. Considering the rehydrogenation from NaH and Al particles (Figure 10a), Na₃AlH₆ as the intermediate must first be formed at the interface of NaH and Al particles, where hydrogen molecules from the gas must be “activated” (reaction (7)). The Na₃AlH₆ intermediate must then be further reacted with Al and H₂ to lead to the regeneration of NaAlH₄.

<figure><img src="image_13.png" /><figcaption>**Figure 10.** Illustration of the hydrogen desorption/absorption principles in complex hydrides. General decomposition path of a) alanates and b) borohydrides; c) hypothetic mechanisms for hydrogen release and uptake in complex hydrides and associated Lennard-Jones potential diagram corresponding to the successive energy barriers encountered by hydrogen during absorption. E<sub>phys</sub>: energy for hydrogen physisorption, E<sub>dis</sub>: energy for hydrogen dissociation, E<sub>chem</sub>: energy for hydrogen chemisorption, E<sub>pen</sub>: energy for hydrogen penetration in the subsurface, E<sub>diff(ion)</sub>: energy for ionic diffusion, and E<sub>nuc</sub> and E<sub>growth</sub>: energy for the nucleation and growth of the hydride. In complex hydrides, the reaction is believed to occur at the interface of segregated elements.</figcaption></figure>
$$Na_3AlH_6 \leftrightarrow 3NaNH + Al + 1.5H_2 \quad (1.8 \text{ mass}\% \text{ H}_2) \quad (9)$$

$$3NaNAlH_4 \leftrightarrow Na_3AlH_6 + 2Al + 3H_2 \quad (3.7 \text{ mass}\% \text{ H}_2) \quad (10)$$

If some Al particles remain trapped within the forming Na₃AlH₆ phase, and transport of Al is limited, then additional Al must be added to the system to reach the full regeneration.[63] Since Ti is a transition metal, it has been claimed from H₂/D₂ scrambling experiments[64] that Ti facilitates the recombination/dissociation of hydrogen molecules at the material/hydrogen gas interface. But this is not sufficient to explain the increase in reaction rate observed because the rate of H₂/D₂ scrambling is much faster than the rate of hydrogen release and uptake of Ti-NaAlH₄. It has thus been suggested that additional Ti dispersed within the material must help with the significant mass transfer involved in this reaction,[63,65] and thus the dissociation/recombination of coexisting NaH and Al phases.[66] In this context, vacancies have also been claimed to play a role.[67] It is noteworthy that despite the significant mass transfer involved, the dehydrogenation of NaAlH₄ and the Ti-doped counterpart have been found to follow a simple Arrhenius law (Equation (6)) and lead to a modification of both the activation energy and the pre-exponential factor (A).[68] At Ti additions >0.9 mol%, the Eₐ was found to remain constant while A increased and was thus solely responsible for the additional improvement in dehydrogenation rate observed.

A key factor for improving the kinetics is also in the understanding of the hydrogen diffusion behavior in materials. For example, it is generally admitted that as the hydrogen molecule dissociates at the surface of metals into individual atoms, these have to diffuse through the metal lattice by occupying interstitial sites within the metal lattice (Figure 8c). This results in elastic interactions between the interstitial atoms, phase transitions, lattice strains, and a modification of the electronic and magnetic properties as hydrogen migrates from one interstitial site to another upon formation of the hydride phase. To diffuse through the metallic lattice, the hydrogen atoms have to overcome the energy barrier generated from the compression of the surrounding atoms, as a consequence the diffusion coefficient D follows the Arrhenius law[11,27m,69]

$$D = D_0 e^{\frac{E_a}{k_B T}} \quad (11)$$

where Eₐ is the activation energy related to the elasticity of the lattice and the size of the atoms, k_B is the Boltzmann constant, T is the temperature, and D₀ is the probability of diffusion, which depends upon the frequency of the thermal vibration (ω₀) of the interstitial atoms of mass (m).

Once in metals the diffusion of hydrogen atoms is relatively fast (e.g., at a rate of 2 × 10¹² H jumps per second at room temperature in vanadium)[70] as compared to complex salts (e.g., 5 × 10³ jumps per second at -223 °C in NaAlH₄—which is equivalent to a relaxation frequency of 10¹¹ s⁻¹ at room temperature).[71] For example, diffusion coefficient of H in Ti₁.₂Mn₁.₈ is ≈10⁻⁷ cm² s⁻¹ at 25 °C,[72] while for NaMgH₃, a diffusion coefficient of ≈10⁻²² to 10⁻²³ cm² s⁻¹ has been determined for the negatively charged Hδ⁻ and associated vacancies.[73] Quantum effects can also be observed because of

the small size of the diffusing hydrogen atoms and this may influence the diffusion process.[27] The migration of hydrogen atoms in metals is believed to occur via different mechanisms depending on the temperature (Figure 8d). At low temperatures where the phonon contribution is minimal, hydrogen diffuses through coherent tunneling. As the temperature is increased the diffusion evolves to incoherent hopping and then with the increasing number of phonons available, the resulting lattice vibration further facilitates the diffusion of hydrogen (Figure 8d). At higher temperatures hydrogen atoms may be regarded as classical particles jumping over barriers via thermal excitations and if the temperature is further increased the hydrogen atoms are free to move as in the gas phase and do not remain within the potential wells of interstitial sites (Figure 8d).

The diffusion of hydrogen in metals can be improved by minimizing the interactions of H atoms with each other's and the metallic atoms of the host material. At low hydrogen concentration in metals, it is believed that the hydrogen occupies octahedral sites in fcc metals, e.g., Cu,[74] Ag,[75] Pd,[76] Pt,[77] Al,[78] and Ni.[79] In this case, the two nearest octahedral sites are separated by a gap of ≈2.5–2.9 Å. In bcc metals, e.g., Fe,[80] V,[81] Nb,[82] Ta,[81a] Mo,[83] and W,[84] when the hydrogen concentration does not exceed the solid-solution phase, hydrogen occupies tetrahedral sites with a separation of 1.01–1.17 Å. Due to short distances, significant hydrogen tunneling is observed and this plays a role in the diffusion coefficient and thus results in a lower activation barrier for diffusion in bcc metals as compared to fcc metals.[27],69,72,85] Similar to bcc metals, hydrogen occupies tetrahedral sites in hexagonal metals like Ti,[86] Zr,[87] Sc,[88] Y,[89] and La alloys;[90] however, due to the degree of high concentration range of solid solutions especially in rare earth metals ≈20% for La and Y at room temperature, the next-nearest neighboring pairs of hydrogen condense strongly to form a bridging metal atom (in c-axis direction).[27] In general hydrogen diffusion is not the rate-limiting step in metal hydrides. However, considering the Mg/MgH₂ systems, theoretical calculations have predicted that the diffusion of hydrogen in MgH₂ can be locally limited by the presence of impurities including doping elements such as Ti, Fe, and Mn often used as catalysts.[91] In contrast, applying external strain has been postulated to enhance hydrogen diffusion within the MgH₂ lattice.[92] Diffusion in amorphous phase/grain boundary is often faster,[93] owing to the broad distribution of site energies. The diffusion of interstitial hydrogen in amorphous materials can be determined from the equation[94]

$$D = D_0 e^{\left(\frac{-Q^o}{RT}\right)} \frac{\partial}{\partial c} \left\{ (1-c)^2 e^{\left(\frac{\mu}{RT}\right)} \right\} \quad (12)$$

where c is the hydrogen concentration for a distribution of site energies, D₀ is a prefactor, and Q° is the average activation energy which is the difference between the constant saddle point energy and the average value of the site energy, R is the universal gas constant, and μ is the chemical potential of hydrogen.

In ionic hydrides, the models above are not valid and the release or uptake of hydrogen is often driven by initial and significant mass transport of hydride ionic species and not
<figure><img src="image_14.png" /><figcaption>**Figure 11.** Illustration of ionic diffusion in complex hydrides. a) Progressive evolution of crystal with zero defects up to a superionic conductor; and b) models of ionic diffusion.[571] In complex hydrides superionic conduction is expected to occur after the melting of the hydride.</figcaption></figure>

hydrogen itself (Figures 10c and 11). For example, in the case of NaAlH₄, the diffusion of AlH₄⁻ and/or AlH₃[67b] and their recombination at the surface of the NaAlH₄ particles into Al and H₂ is believed to control the desorption rate.[67a,95] Similarly, in the case of borohydrides, translational diffusion of BH₄⁻ was found to be dominant upon hydrogen release.[96] As a consequence, improving the hydrogen sorption kinetics of ionic hydrides strongly depends upon means to improve ionic conduction. As discussed above, diffusion of mobile species in a crystalline solid needs to overcome energetic barriers across two local energy minima. For ionic solids, the ionic conductivity (σ) can be expressed as[97]

$$ \sigma = N(\text{Ze})\mu_{\text{M}} \quad (13) $$

where N is the concentration of mobile ions (cationic or anionic), Ze is the ionic charge, and μM is the rate of diffusion of the mobile ion per unit of electric field, i.e., the ion mobility. Both the ion mobility and concentration of the charge carriers are temperature dependent following

$$ N = N_0 \exp(-E_f/k_B T) \quad (14) $$

$$ \mu_M = \mu_0 \exp(-E_m/k_B T) \quad (15) $$

where N₀ and μ₀ are the pre-exponential factors, and Eₜ and Eₘ the energies of formation and migration of the mobile charge carriers, respectively. The activation energy Eₐ for the diffusion of ions in solids is thus the sum of Eₜ and Eₘ.

Improving diffusion in ionic compounds depends highly upon the amount of interstitials, vacancies and partially occupied lattice sites (Figure 11a), as these provide minimum energy diffusion pathways.[98] Interstitials and vacancies sites can be created by aliovalent substitution of cations/anions. However, this approach often reaches a maximum diffusion at the optimum aliovalent substitution. Beyond this optimum, the strong lattice distortion induced by the excessive aliovalent concentration induces a slower ionic diffusion. For example, Ca

aliovalent doping in LiBH₄ showed an increase in Li-ion conductivity up to a maximum of 4 mol% Ca(BH₄)₂ doping, then the conductivity decreases.[99]

Both ion valency and size can influence ionic conductivity.[100] Indeed, a decrease in ionic diffusion can be observed when electrostatic interaction between the mobile ionic specie and the ion forming the structural skeleton is increased. Similarly, increasing the valence of the mobile ionic species as well as ionic size reduces mobility. Figure 12 aims to summarize these effects for borohydrides, although the trends are less obvious than those observed for solid-state battery electrolytes[100] owing to the limited number of available data. If the diffusing ion is too small it may remain trapped within large electrostatic wells of the larger ions forming the skeleton. On the other hand, if the diffusing ionic species are too large, they will experience large bottlenecks when diffusing between the smaller skeleton ions. Therefore, achieving fast hydrogen kinetics in ionic hydrides requires a careful balance of often diverging requirements.

A hopping mechanism is generally used to describe the diffusion of ionic species (Figure 11). In this mechanism, ions can jump from one site to a neighboring site through vacancies such as Schottky defects or via interstitial defects (Frenkel defects) as described in Figure 11b. However, it should be noted that diffusion mechanisms may take additional forms depending upon the material structure and the temperature. For example, in the superionic state a caterpillar mechanism (Figure 11b), whereby the mobile ion can jump to a neighboring site even if occupied, may take place as well as continuous diffusion at high temperatures.[101]

Diffusion coefficients of 5 × 10⁻¹⁶ m² s⁻¹ at 75 °C and 7 × 10⁻¹⁴ m² s⁻¹ at 200 °C have been reported for the BH₄⁻ and AlH₄⁻ anions, respectively.[102] Li⁺ conductivity in LiAlH₄ was found to be of 2.9 × 10⁻⁶ S cm⁻¹ at 107 °C.[102] At room temperature the tetrahedral BH₄⁻ unit reorients at a fast rate,[103] and a slow Li⁺ hopping motion has been suggested to occur.[104] At higher temperatures (>120 °C), upon the structural phase transition of LiBH₄ from orthorhombic to hexagonal the diffusion of Li⁺ is very fast[104,105] and this is accompanied with a rapid reorientational motion of the BH₄⁻ ions (at a rate of
<figure><img src="image_15.png" /><figcaption>**Figure 12.** Ionic diffusion in complex hydrides. Summary of reported values of a) ionic conductivity of Li⁺, Na⁺, and Mg²⁺ in respective borohydrides, and b) the associated activation energy for diffusion as function of the cation valency at 25 °C. Activation energy for the ionic diffusion of BH₄⁻ as function of c) the cation valency and d) the cation radii of borohydrides. The ions should not be too small or too large.</figcaption></figure>

≈10¹³ s⁻¹)[¹⁰⁶] suggesting that the fast rotation of the anion is facilitating the translational diffusion of Li⁺. Upon melting of LiBH₄ at 450 °C (before hydrogen release), both translational diffusion of BH₄⁻ (with a diffusion coefficient of 10⁻⁹ m² s⁻¹)[⁹⁶] and Li⁺ is believed to occur. The fast diffusion of Li observed in LiBH₄ and its wide electrochemical window has triggered significant work toward the use of borohydrides as solid state ionic conductors for batteries.[¹⁰⁷] For example, a partial substitution of the BH₄⁻ ion with I⁻ has shown to lead to a Li⁺ conductivity of ≈10⁻⁴ S cm⁻¹ at 60 °C.[¹⁰⁸] Similarly, lithium amide-borohydrides showed Li⁺ conductivity of 10⁻⁴ to 10⁻³ S cm⁻¹ at room temperature.[¹⁰⁹] Similar improvements in conductivity have been reported for NaBH₄-based electrolytes, and in particular NaCB₉H₁₀ showed Na⁺ conductivity of ≈10⁻² S cm⁻¹ at the ambient.[¹⁰⁷f] However, it should be noted that the improvements of Li⁺ or Na⁺ conductivity did not effectively translate into better overall hydrogen storage properties of complex hydrides, because ionic diffusion is only one parameter among the multi-interrelated factors controlling the storage mechanism of hydrogen in complex hydrides.

## 3.2. Current Understanding of the Basic Thermodynamics Rules Controlling the Temperature for Hydrogen Sorption

The Gibbs free energy difference between products and reactants (ΔG) is the driving force of a reaction and it has to be negative for the reaction to occur (Figure 13d). Therefore, controlling ΔG is crucial for the thermodynamics of the reaction hydrogen with sorption materials. The definition of Gibbs free energy is G = H - TS. At a fixed temperature (T)

$$ \Delta G = \Delta H - T\Delta S \quad (16) $$

where ΔH is the enthalpy change of the reaction, while ΔS is the entropy change of the system. In the case of hydrogen/metal systems, ΔS mainly results from the entropy (ΔS<sub>gas</sub>) of the hydrogen gas which is lost upon hydrogen absorption in the material. Other entropy effects (ΔS<sub>e</sub>) as summarized below are relatively small in comparison to ΔS<sub>gas</sub>[²³]

$$ \Delta S_e = \Delta S_H^{\text{vib}} + \Delta S_{\text{host}}^{\text{vib}} + \Delta S^{\text{el}} + \Delta S^{\text{conf}} \quad (17) $$

where ΔS<sub>H</sub><sup>vib</sup> corresponds to the vibrational entropy of the hydrogen atoms in the hydrides, ΔS<sub>host</sub><sup>vib</sup> is related to the vibrational spectrum of the host material, ΔS<sup>el</sup> is due to the difference in the electronic heat capacity between the materials and its hydrided state, and ΔS<sup>conf</sup> corresponds to configurational entropy of the vacancies in the hydride

Under reference state (25 °C, P°)

$$ \Delta G^{\circ} = \Delta H^{\circ} - T\Delta S^{\circ} \quad (18) $$

also

$$ \Delta G^{\circ} = -RT \ln K_{\text{eq}} \quad (19) $$

and for a reaction

$$ \Delta G = \Delta G^{\circ} + RT \ln Q_{r} \quad (20) $$

where K<sub>eq</sub> is the equilibrium constant and Q<sub>r</sub> the reaction quotient.
<figure><img src="image_16.png" /><figcaption>**Figure 13.** Basic tool box for designing hydrogen storage materials. a) Ease of reaction path via H⁻, H, or H⁺ for the generation of hydrogen,[22] b) enthalpy of formation of various metal hydrides; c) metal hydrogen bond strength of potential catalysts; and d) general principles for adjusting the properties of hydrides. Both thermodynamic and kinetics should be improved at the same time. The hydride path is potentially better to release hydrogen,[22] although it may lead to irreversibility. Combining unstable and stable hydrides can lead to intermediate enthalpies of reaction with hydrogen.</figcaption></figure>

For a hydriding reaction, assuming that the rest of the reactants are solids (which is the case for most metal hydrides and complex hydrides), the pressure of hydrogen is the only factor that needs to be taken into account and thus $K_{eq}$ and $Q_r$ take the form

$$K_{eq} = \frac{P^{\circ}}{P^{\text{eq}}} \qquad (21)$$

$$Q_r = \frac{P^{\circ}}{P_{\text{op}}} \qquad (22)$$

where $P_{op}$ is the operating pressure.

Therefore, the Van't Hoff plot for the reaction can be derived from (17)–(19) as

$$\ln \frac{P^{\text{eq}}}{P^{\circ}} = \frac{\Delta H^{\circ}}{RT} - \frac{\Delta S^{\circ}}{R} \qquad (23)$$

and by associating (19)–(22), the Gibbs free energy of the hydriding reaction is

$$\Delta G = RT \ln \frac{P^{\text{eq}}}{P_{\text{op}}} \qquad (24)$$

where $P^{\text{eq}}$ is the equilibrium pressure. Therefore, when $P_{op} > P^{\text{eq}}$, $\Delta G < 0$, the hydriding reaction is thermodynamically favored; and the reverse applies for the dehydriding reaction.

The relation between the equilibrium pressure and operating temperature is represented by the pressure composition temperature (PCT) curves as shown in **Figure 14a**. A hysteresis is often observed experimentally between the equilibrium pressure of absorption $P_{\text{abs}}^{\text{eq}}$ and desorption $P_{\text{des}}^{\text{eq}}$ in metallic systems (Figure 14a) and this can be expressed as an energy difference following

$$\Delta G_{\text{hyst}} = RT \ln(P_{\text{abs}}^{\text{eq}}/P_{\text{des}}^{\text{eq}}) \qquad (25)$$

The hysteresis observed in metallic systems is often explained by the occurrence of irreversible plastic deformations upon hydrogen absorption and was initially observed in the Pd–H system.[110] In some compounds, such as the AB₅, the hysteresis may decrease after a few cycles, owing to a reduction of $P_{\text{abs}}^{\text{eq}}$ and this has been attributed to the introduction of defects as a result of the significant lattice expansion experienced by the alloy upon hydrogen insertion.[111] In complex hydrides similar hysteresis have also been observed, e.g., in LiBH₄ + MgH₂,[112] and NaBH₄ + MgH₂ systems.[113] It should be noted that sloppy equilibrium plateaux are often assigned to inhomogeneity in the alloy composition.[114] However, this can also be observed in pure metals and thus the hypothesis is that the slope is the result of constraints generated during the insertion/deinsertion of the hydrogen atoms in the metal.[115] In nanosized metal hydrides, similar sloppy plateau pressures have been interpreted along a broader particle size distribution, a structural heterogeneity of the nanoparticle, or interfacial
<figure><img src="image_17.png" /><figcaption>**Figure 14.** Current approaches for adjusting the thermodynamics of hydrides. a) Pressure composition temperature (PCT) plot (left) and the corresponding van't Hoff plot (right) for hydrogen storage reactions in metals and complex hydrides; b) generalized enthalpy diagram illustrating the destabilization of hydrides through: a) alloy formation with less stable hydrides, e.g., Fe, Ni (green); reaction of a hydride, e.g., LiH, MgH₂, NaBH₄, NaAlH₄, with a corresponding protide, e.g., LiNH₂, NH₃ (orange); stabilization of the decomposition products (blue); and nanosizing (pink). To enable room temperature hydrogen storage, ΔH should be brought down while ΔS remains constant.</figcaption></figure>

stress at the stabilizer/nanoparticle interface.[116] Such a sloppy plateau complicates the utilization of hydrides in real systems because of the inherent variation of the hydrogen released pressure In addition, the strong stress (≈several GPa) undergone by the alloys absorbing hydrogen results in a decrepitation and thus a pulverization of the alloy particles.[117]

From the equilibrium plateau pressure measured at different temperatures, a corresponding van't Hoff plot can be generated based on the PCT (Figure 14a). At a given operating pressure (P<sub>op</sub>), there is a corresponding temperature (T<sub>op</sub>) for hydriding and dehydriding reactions. The operating conditions (P''<sub>op</sub> and T''<sub>op</sub>) can be tuned by altering the thermodynamics, i.e., the enthalpy and entropy of the reactions, which are related by the Gibbs free energy (24). In equilibrium, ΔG equals to zero, and the decomposition temperature (T<sub>dec</sub>) at 0.1 MPa can be expressed as

$$T_{\text{dec}} = \frac{\Delta H^{\circ}_{\text{dec}}}{\Delta S^{\circ}_{\text{dec}}} \qquad (26)$$

Some of the desorbing temperatures of common hydrides are tabulated in several references[26,118] and the most interesting ones, of high hydrogen capacity, require high temperatures for releasing hydrogen (Table 2). Ideally, high capacity hydrides should be capable of absorbing and releasing hydrogen at near ambient temperature. Assuming a constant ΔS (130.7 J mol<sup>-1</sup> K<sup>-1</sup> of H₂, for most metal hydrides),[26] the enthalpy of decomposition and thus the strength of the metal-

hydrogen (M-H) bond is the main indicator of the thermal stability of the hydrides. If the enthalpy is high, a low equilibrium pressure and high decomposition temperature are expected (Figure 13). To achieve an equilibrium pressure of 0.1 MPa and room temperature, the enthalpy of decomposition should be close to 19.6 kJ mol<sup>-1</sup>. One approach to achieve this is through the combination of elements with opposite bonding strength to hydrogen (Figures 13b and 14b). For example, the combination of 3/2MgH₂ + 1/2MgCu₂[119] and MgH₂/Al[120] has been shown to lead to some improvement in the thermodynamics, e.g., a decreasing in the enthalpy of reaction from 75 kJ mol<sup>-1</sup> H₂ for pristine Mg/MgH₂ to 70 kJ mol<sup>-1</sup> H₂ once Mg is milled with Al.[121] However, this occurs with a concomitant decrease in hydrogen storage capacity.

One framework to understand the strength of the M-H bonds is by considering the standard enthalpy of formation (ΔH°<sub>f</sub>), which is related to the enthalpy of atomization of the M (ΔH°<sub>atom</sub>), enthalpy of ionization of M (ΔH°<sub>ioniz</sub>), the crystal or lattice energy (ΔH°<sub>cryst</sub>), and thus according to the Born-Haber cycle[22]

$$\Delta H^{\circ}_{f} = \Delta H^{\circ}_{\text{atom}} + \Delta H^{\circ}_{\text{ioniz}} + n(1.51 \text{ eV}) - \Delta H^{\circ}_{\text{cryst}} \qquad (27)$$

where n corresponds to the stoichiometry of the hydride (MH<sub>n</sub>).

From this equation it can be deduced that only stable hydrides are formed for (ΔH°<sub>cryst</sub>) significantly larger than ΔH°<sub>atom</sub> + ΔH°<sub>ioniz</sub>, which is the case for ionic hydrides of relatively low melting point.
Since $\Delta H^{\circ}_{\text{atom}}$ and $\Delta H^{\circ}_{\text{ioniz}}$ are correlated to the standard redox potential ($E^{\circ}$) of hydrides and a correlation exists between the decomposition temperature ($T_{\text{dec}}$) of hydrides and the standard redox potential of binary hydrides, Grochala and Edwards proposed the use of such a correlation to predict $T_{\text{dec}}$ of ternary hydrides.[22] Accordingly, in a similar manner $\Delta H$ can be used to design hydride materials of targeted $T_{\text{dec}}$ (Figure 7), through an appropriate choice and combination of elements forming hydrides with intermediate $E^{\circ}$. However, the experimental hydrogen behavior of the ternary hydrides predicted by this assumption is often difficult to control owing to the multiple decomposition paths of hydrides and their reaction intermediates.[122] As compared to binary metal hydrides, controlling the thermodynamics of complex hydrides is much more difficult due to their dehydrogenation along multistep paths with additional side reactions generating byproducts.[123] The most common complex hydrides of interest for high capacity hydrogen storage are alanates ($[\text{AlH}_4]^{-1}$), borohydrides ($[\text{BH}_4]^{-1}$) and amides ($[\text{NH}_2]^{-1}$). The full decomposition of alanates is a two or three-step process depending on the metallic cation involved. For example (Figure 10a), the decomposition of $\text{LiAlH}_4$ and $\text{NaAlH}_4$ involve the intermediate phase $[\text{AlH}_6]^{3-}$, while $\text{Mg(AlH}_4)_2$ decomposes directly to $\text{MgH}_2$ and $\text{Al}$.[124] Similarly, the decomposition of borohydrides can be assumed to be a two-step reaction, including the formation of the corresponding metal hydride and boron in the first instance and then further decomposition of metal hydride (Figure 10b).

As for amides, the decomposition process is even more intricate.[125] The direct decomposition of amides alone does not lead to the release of hydrogen, but ammonia or nitrogen instead.[16a,125,126] Hydrogen released from amides requires the aid of metal hydrides[118a]

$$ \text{LiNH}_2 + 2\text{LiH} \leftrightarrow \text{Li}_2\text{NH} + \text{LiH} + \text{H}_2 \quad (28) $$

$$ \text{Li}_2\text{NH} + \text{LiH} \leftrightarrow \text{Li}_3\text{N} + \text{H}_2 \quad (29) $$

Mixture of amide and metal hydride of different metals may also result in different reaction paths, and thus complexity further escalates. Hence to date, it is very difficult to translate an successful approach from one hydride to the others.

It is often assumed that the hydrogen sorption behavior of complex hydrides mainly depends upon the $\Delta H$ of the reaction although significant changes in $\Delta S$ may also occur, in addition to the related changes in $\Delta S$ associated with the transition of hydrogen from the gas phase to the hydride phases. In fact, the entropy of complex hydrides strongly varies during hydrogen desorption/cycling and is a function of the decomposition path as highlighted below[127]

$$ \text{LiBH}_4 \rightarrow \text{Li} + \text{B} + 2\text{H}_2 \quad \Delta S = 219 \text{ Jmol}^{-1} \text{K}^{-1} \text{H}_2 \quad (30) $$

$$ \text{LiBH}_4 \rightarrow \text{LiH} + \text{B} + 3/2\text{H}_2 \quad \Delta S = 150 \text{ Jmol}^{-1} \text{K}^{-1} \text{H}_2 \quad (31) $$

It has thus become crucial to develop tools not only to control $\Delta H$ but also the $\Delta S$ of the hydrogen reaction with sorption materials (see Section 8.1.3).

### 3.3. General Principles to Design Hydrogen Storage Materials

By taking into account the above considerations, Figure 13 summarizes the basic “tool box” to design hydrogen storage materials. First of all, a form of hydrogen, i.e., as atomic (H), hydride (H⁻), protide (H⁺), or a combination (Figures 13a and 14) believed to lead to a low reversible energy path is chosen. This choice is further dictated by the requirement of high gravimetric and volumetric hydrogen density of the hydride material, and this not only further limits the list of potential candidates but also brings an additional level of complexity in achieving low hydrogen release and uptake temperatures. Indeed, as discussed above, the elements capable of binding large amounts of hydrogen often form stable hydrides and thus release their hydrogen at high temperatures (Figures 4 and 7). To overcome such a drawback, the common strategy is then to adjust the enthalpy of the reaction by combining the chosen hydride with elements forming unstable hydrides (Figure 13b). If the rate of hydrogen release and uptake is not adequate, the search for a catalyst begins by assessing the potential of various transition metals that may help the dissociative absorption and/or recombination of hydrogen during desorption. Figure 13c summarizes some potential catalysts. The hydrogen dissociation rate of the catalyst should be high enough and its binding strength to hydrogen relatively low to facilitate hydrogen release/uptake from the hydride material. Hence, the difficulty in adjusting the properties of hydrogen storage materials is in the requirement of combining strategies for improving both thermodynamics and kinetics simultaneously (Figure 13d), without “losing” the hydrogen storage capacity. Surprisingly, these basic principles/concepts, developed in the very early stages of the field of hydrogen storage research, for the design of hydrogen storage materials still continue to form the basis of the more advanced and recent approaches further discussed below.

### 4. Metallurgical Synthetic Approaches to form Metal Alloys and Alternative Phases to Store Hydrogen

With the discovery of hydrogen sorption in Pd[128] and the early alloys to effectively store hydrogen at the ambient, initial research efforts have focused on metallurgical approaches to design better hydrogen storage alloys. Hence, a vast majority of the early research efforts has been devoted toward a fundamental understanding of the behavior of hydrogen in metals. To date, this knowledge is more than relevant as alloys such as those based on $\text{LaNi}_5$, $\text{TiFe}$ and $\text{Ti BCC}$ have potential to open new applications in the storage of renewable energy. This section is thus reviewing current understanding along the modification of the hydrogen storage properties of metal alloys as well as conventional methods to prepare such alloys.

#### 4.1. Modifying the Thermodynamics of Metal Hydrides

Many metals are known to form hydrides (Figure 4), hence by an alloying approach several metal alloy families have been found to effectively store hydrogen with altered properties as compared to
pure metals. These alloys include intermetallics, e.g., AB, AB₂, AB₅, A₂B, and solid solutions.[26] Solid solutions are formed by dissolving one or more minor elements (solutes) into a hydrogen absorbing primary element (solvent). Examples of solid solutions include Ti and V-based solid solution hydrides such as Ti–V–Fe (Ti₄₃.₅V₄₉.₀Fe₇.₅),[118b] V–Ti–Cr,[129] and V–Ti–Ni,[130] and unlike stoichiometric or near-stoichiometric intermetallic compounds, solid solutions can also have nonstoichiometric arrangements.

Intermetallic compounds correspond to an intermediate phase between A and B. Alloying two elements increases the degrees of freedom in the final materials and leads to various structural defects including vacancies and antiphase boundaries that may benefit the overall hydrogen storage properties. In these families, A is a transition metal or a rare earth with a strong affinity for hydrogen and thus of high thermal stability, e.g., YH₂, LaH₂, TiH₂, and ZrH₂. B corresponds to transition metals forming unstable hydrides, e.g., Cr, Mn, Fe, Co, and Ni absorbing hydrogen at very high hydrogen pressures only (Figures 4 and 13). Intermetallics synthesized through the association of A and B usually form alloys with intermediate hydrogen sorption properties. This is typically the case for a range of intermetallics including TiFe, ZrV₂, and alloys of the type CaCu₅. For example, LaNi₅ has an enthalpy of –32 kJ mol⁻¹ H₂,[131] while La forms a stable hydride (LaH₂) with ΔH = –208 kJ mol⁻¹ H₂ and Ni forms an unstable hydride with ΔH = –6 kJ mol⁻¹ H₂.[11]

Table 2 lists elemental hydrides with their respective heat of formation, entropy values and decomposition temperature values. Using this list, one can calculate the enthalpy of formation (ΔHf) of ternary hydrides (AₓBᵧHₓ+y) by using the semiempirical model proposed by Miedema et al.[132a]

$$ \Delta H_f (A_x B_y H_{m+n}) = \Delta H(A_x H_n) + \Delta H(B_y H_m) - \Delta H(A_x B_y) \quad (32) $$

where ΔH(AₓHₙ) and ΔH(BᵧHₘ) are the heats of formation of hydride A and B and ΔH(AₓBᵧ) is the heat of formation of AₓBᵧ alloy. This model known as Miedema's rule of reverse stability was found to work well for ternary hydrides where A (with a strong affinity for hydrogen) is the minor phase as compared to B (with a low hydrogen affinity).

A more accurate model to determine the enthalpy of formation (ΔHf) of ternary hydrides is the semiempirical band-structure model of Griessen and Driessen[133]

$$ \Delta H_f (A_{yA} B_{yB} H_n) = \alpha \left( \frac{\gamma_A \Delta E_A^* + \gamma_B \Delta E_B^*}{\gamma_A + \gamma_B} \right) + \beta \quad (33) $$

where γi is related to the stoichiometric amount of the metal i in the ternary alloy (A_{yA}B_{yB}H_n), ΔEᵢ* corresponds to the band structure energy parameter of the scaled density of state of the metal i after alloying, and α = 29.62 kJ eV⁻¹ mol⁻¹ H, β = –135.0 kJ mol⁻¹ H.

When both metals have the same d-band width but different density of states at the Fermi energy, (33) can be simplified as

$$ \Delta H^{\circ}(A_{1-y}B_{y}H_{n}) = (1-y)\Delta\tilde{H}^{\circ}(AH_{n}) + y\Delta\tilde{H}^{\circ}(BH_{n}) \\ - \alpha ny(1-y)(\emptyset_{A} - \emptyset_{B})(n_{A} - n_{B}) \frac{1}{(1-y)n_{A} + yn_{B}} \quad (34) $$

where ΔH̃°(iHₙ) is the integral enthalpy of formation for one mole of binary metal hydride i, i.e., ΔH°(MHₙ) = ΔH̃°(iHₙ)/n, Øi is the work function and ni the value of density of states at the Fermi energy of the metal i.

It should however be noted that as compared to the Miedema's model, which was designed to describe a wide class of alloys, the band-structure model is limited to metal hydrides where all hydrogen atoms are surrounded by a fraction of A and B atoms. Therefore this model does not apply to disordered alloys or intermetallic compounds such as AB₂ Laves phases which can contain several types of interstitial sites including A₂B₂, AB₃, and B₄ sites. More recent approaches building on data mining and machine learning techniques (see Section 9.2.3), e.g., artificial neural network and principal component analysis, have also emerged with the aim of providing more accurate methods for the prediction of the enthalpy of AB intermetallics.[134]

## 4.2. Tuning the Hydrogen Thermodynamics of Alloys by Partial Elemental Substitution

Partial substitution provides a means to adjust the thermodynamic properties of hydrides, and the approach has been widely used to design hydrides with capability for hydrogen storage at the ambient and with capacity beyond the hydrogen storage capacity of LaNi₅ and FeTi, i.e., >1.2–1.9 mass% H₂.[135] For example, Reilly and Wiswall[136] reported in 1968 that alloying nickel with magnesium decreased the enthalpy of reaction with hydrogen to –64.5 kJ mol⁻¹ H₂ instead of the –75 kJ mol⁻¹ H₂ of the Mg/MgH₂ system. This meant a temperature of absorption of ≈ 240 instead of 400 °C.[46a] However, this was at the expense of the hydrogen storage capacity, i.e., it was reduced from 7.6 to 3.6 mass% H₂. By further tuning of the thermodynamics via the addition of Ca, the (Mg₁-xCaₓ)₂Ni alloy displayed an enthalpy of –37 kJ mol⁻¹ H₂ and thus hydrogen desorption at room temperature but with a storage capacity of 1.2 mass% H₂.[137] This demonstrates the potential of the approach but also highlights the difficulty of modifying the properties of hydrogen materials by a simple partial substitution.

For some hydrides, a correlation between elemental substitution and Peq has been evidenced and this may help to some extent with the adjustment of the thermodynamic properties. Taking LaNi₅ as an example, a partial substitution of the Ni leads to a decrease of Peq and a reduction of the storage capacity. When Ni is partially substituted with Al, this leads to a 50% loss of hydrogen storage capacity owing to the larger size of the Al atoms (Figure 15b). Partial substitution with Co (>0.25%) can lead to two Peq as the result of the formation of an intermediate γ phase.[138] The evolution of Peq has also been shown to be linearly correlated to the composition of the alloy and thus the evolution of cell volume upon partial substitution of La or Ni (Figure 15d).[42a] From this variation, one can thus tune Peq to the desired value.

Contrary to MgH₂ and LaNi₅, a small amount of substitution of Mn in TiMn₁.₅ was found to cause no loss of capacity with a notable shift in the plateau pressure.[139] For TiMn₁.₅, Peq was shifted up by substituting Mn with Cu, Fe, and Co; while V and Ni led to a reduction of Peq. The potential of partial substitution
<figure><img src="image_18.png" /><figcaption>**Figure 15.** Tuning the thermodynamic of intermetallics with LaNi₅ as an example. a) Phase diagram of the La-Ni system; b) evolution of the equilibrium plateau pressure of LaNi₅ upon partial alloying with various elements; c) crystalline structure of LaNi₅ structure showing the position of the hydrogen atoms (the c-axis has been doubled in length for clarity); and d) linear dependence of the equilibrium plateau pressure of LaNi₅ as function of the evolution of the cell volume upon a partial substitution of La by R or Ni by M in LaNi₅. Reproduced with permission.[42a] Copyright 2001, Springer-Verlag.</figcaption></figure>

may thus be an effective means to adjust the plateau pressure of hydrides. However, this may also pose problems during large scale manufacturing, as a small amount of impurities (0.1 by stoichiometric composition) can cause shifts in plateau pressure outside the desired range.

### 4.3. Methods for Making Hydrogen Storage Alloys and Intermetallics

The first step to make an alloy of a given composition is to check the phase diagram of the elements to be alloyed in order to determine possible miscibility issues, phase formations, and ultimately determine the synthetic conditions, i.e., temperature, pressure, and heat treatment. For binary alloys the phase diagrams are relatively well known.[140] However, the phase diagram of ternary or higher intermetallic compounds are scarce.[141] Therefore in this case, a trial and error approach is the only option to achieve the targeted composition and phase. The phase diagram of La-Ni is shown on Figure 15a, as an example. The La atom can be substituted by many elements, e.g., 4f elements such as Ce, Nd, and Pr owing to the existence of many AB₅ compounds.[142] The possibilities to substitute Ni are more complicated. Ni can be fully substituted by Cu, Co, or Pt, but it is also possible to substitute Ni in LaNi₅ with low levels of Al,[143] Mn,[144] or Fe,[145] for example.

The level of substitution varies strongly depending on preparation method and the duration and temperature of the heat treatment process. For example, in LaNi₅-xMnₓ, x can be 0.99 if the alloy is annealed at 1175 °C and 2 for annealing at 800 °C.[144,146] The solidification of the alloy phase can occur congruently; however, this is not always the case (e.g., FeTi exhibits a peritectic solidification). It should be noted that depending on the level of deviation from the stoichiometric composition, the amount of impurities (e.g., C, N, O), and the volatility of the melted elements, additional vacancies and structural defects may be created within the alloy lattice. With LaNi₅ this was not observed. However, in FeTi, an excess of one of the atoms (Fe or Ti) was found to lead to a substitution on the deficient sublattice.[147]

The common procedure to prepare alloys is through melting methods including radiofrequency induction and arc melting. Once a homogenous melt is formed, the cooling rate of the melt should be carefully controlled to avoid any phase segregation. For a small sample size this can be achieved by quenching the melt (e.g., at a cooling rate >10 °C s⁻¹) to limit any shift in elemental concentration across the alloy. When the solidification of the melt is peritectic, additional annealing in high vacuum or an inert atmosphere (to avoid oxidation) is often required to facilitate the homogenization of the alloy and remove any microsegmentation and structural defects. Homogenization is beneficial for the hydrogen sorption properties because it often leads to a “flat” equilibrium plateau pressure without significant hysteresis.[26]
<figure><img src="image_19.png" /><figcaption>Figure 16. Illustration of mechanical milling tools. Summary of milling conditions for a) attritor mill; b) spex mill; and c) planetary mill.</figcaption></figure>

## 5. Mechanical Milling for Making New Hydrogen Storage Materials

Following the development of mechanical alloying by high energy ball milling technique around 1966 at the Paul D. Mercia Research laboratory to produce nickel-based superalloys for gas turbines,[148] mechanical milling/alloying emerged in the late 1980s as an approach for the synthesis of hydride materials. In particular, Ivanov et al. were among the first to report the synthesis of Mg-based alloys with Ni, Fe, Co, or Ce for hydrogen storage.[149] The nanostructuration by ball milling of Mg₂Ni[150] and then MgH₂[151] both leading to significantly enhanced hydrogen kinetics further established mechanical milling as a significant tool for the modification of the properties of hydride materials.[152] Mechanical milling involves a very simple experimental setup and operating conditions, and the possibility of producing large quantities of solvent free material further contributed to the widespread uptake of this technique across research laboratories.[153] Various terminologies are commonly used to describe the different type of processes that can be achieved through mechanical milling: i) mechanical alloying describes the processing of mixtures of metallic powders to form alloys,[154] ii) mechanical milling refers to the refinement of powder, iii) mechanochemical synthesis is related to the acceleration of solid state synthetic reactions upon milling between chemical powders, iv) cryomilling refers to milling carried out at liquid nitrogen temperature—usually to facilitate powder refinement and further reduction in particle size, v) reactive milling involves the milling of powders in the presence of chemically reactive solids, liquids, or gases (e.g., hydrogen, oxygen, nitrogen, air, etc.)—for example, the milling of Mg under hydrogen atmosphere to produce MgH₂.[155]

During the milling process, the material is subjected to different stresses including compression, deformation, attrition, shear, and impact resulting from the milling tools

(Figures 16 and 17).[153,156] This generates new surfaces and cracks but also defects and dislocations resulting in a reduction of crystallite sizes, a higher number of grain boundaries, and a high surface area. All these are supposed to facilitate the kinetics of hydrogen release/uptake and diffusion in hydrides.[118b,157] Depending on the ductility of the powder, agglomeration of particles may also occur upon ball impact by cold-welding, mechanical interlocking resulting from rough surfaces or chemical reactions between particles.

### 5.1. Fundamentals of the Mechanical Milling Process

The process of mechanical milling is simple and conducted by placing the powder(s) to be refined, alloyed, or reacted in a milling vial with balls that are free to impact the powder.[157a,158] The vial and/or balls are shaken by various means (Figure 16) and the powder is milled upon repeated and intermittent impact of the balls until a steady state between cold-welding and fracturing is achieved. Instead of balls, cubes, rods, ballpebs, and cylpebs have been tried, but the spherical shape of the milling balls offers a higher gridding surface area and favors lower wear.[159] Several theoretical models have been proposed in an attempt to describe the milling process;[160] however, in a first approximation, the intensity involved in the milling process can be simplified as the product of the kinetic energy (ΔE) released in a single impact of a ball and the total number of impacts per unit of time ($N_i$)

$$I = \varepsilon \Delta E N_i \qquad (35)$$

where ε is a factor comprised between 0 (elastic collision—no energy transferred) and 1 (perfect inelastic collision), and ΔE = 1/2m<sub>b</sub>v<sub>b</sub><sup>2</sup>, with m<sub>b</sub> and v<sub>b</sub> corresponding to the mass of the ball and its velocity. Accordingly, the overall specific energy (W<sub>s</sub>) involved in milling the material is given by[161]
<figure><img src="image_20.png" /><figcaption>**Figure 17.** Principles of the mechanical milling process. a) Particle fracturing upon ball-powder-ball collision; b) deformation characteristics of ductile and brittle hydride materials upon mechanical milling; c) evolution of particle/grain size for brittle materials; and d) flow chart showing the different paths for using mechanical milling for preparing hydrides.</figcaption></figure>

$$W_s = \frac{1}{M} lt \qquad (36)$$

where M is the mass of material milled and $t = (t_b/t_{imp})\tau$. $t_b$ corresponds to the ball impulse duration, $t_{imp}$ the impulse period and $\tau$ the grinding duration.

It can thus be concluded from the equation above that $W_s$ depends only on intrinsic parameters of the milling process including the mass of the grinding balls, the rotation frequency of the vials, the ball-to-powder ratio, and the milling time.

The initial particle size of the starting material does not matter because upon milling the particle size is rapidly reduced to a few µm within the first few minutes of milling as it follows an exponential decay (Figure 17c).[162] However, the ductility or brittleness of the powder to be milled will significantly affect the end results. Upon milling, ductile particles will tend to be flattened and cold welded and thus no particle size reduction occurs but some alloying at the atomic level may happen (Figure 17b). In contrast, brittle particles are extensively fragmented; however, at very small particle sizes cold-welding occurs because the particles then behave in a ductile manner. A mixture of ductile–brittle leads to the embedment of the brittle particles within the ductile matrix.

The main types of high energy mills found in many laboratories are summarized in Figure 16 as well as the velocity of the balls, kinetic energy per impact, shock frequency and power involved. Spex mills are much more energetic than attritor mills. Hence, for the same end result much shorter milling times are required with a high energy shaker mill. However, a higher energy of milling also involves an increase in the milling vial temperature owing to the impact of the balls. It has been estimated that only 5–15% of the energy is effectively

used toward milling, while the remaining energy is dissipated as heat.[163] Higher temperatures may facilitate atomic/ionic diffusion and thus lead to a better homogenization of the final product; but these may also trigger the premature decomposition of the milled powder and thus be detrimental.

The overall increase in macroscopic temperatures of the mill has been found to be moderate, e.g., ≈50 °C for a Spex mill,[164] but at the microscopic level upon ball-to-ball or ball-to-wall collision temperatures have been predicted to be as high as 350 °C.[164b,165] However, by observing the phase transformation occurring on some materials during milling, it was estimated that the temperature locally raised to at least 1000 °C during the processing of γ-AlO(OH) on a planetary mill, for example.[166] To minimize this, mechanical milling can be carried out with the aid of a cooling medium including air, water, or liquid nitrogen.[167] However, in practice, long milling runs are often divided into shorter milling periods paused with cooling breaks to dissipate any excessive heat. This is particularly important when milling complex hydrides with low decomposition temperatures (Tables 2 and 3).[122,168]

Various grinding materials can also be used and the most common milling tools are made of hardened steel, chromium steel, tungsten carbide, agate, silicon nitride, and sintered corundum. However, the density of the milling balls should be high enough to create sufficient impact on the powder to be milled. Furthermore, when possible the milling tools should be made of the same material as the powder to avoid cross contamination otherwise the milling tools should be harder than the powder to mill to avoid excessive cross-contamination. Indeed, a drawback of mechanical milling is the contamination emerging from the milling tools. In particular, using stainless steel in a high energy shaker mill can introduce a large amount
<table><caption>**Table 3.** Summary of selected hydrides modified or synthesized by mechanical milling. Notations are as follow: Reactant melting point (R<sub>mp</sub>), hydrogen pressure (P<sub>H</sub>) used during milling, total milling time (T<sub>t</sub>), milling speed (M<sub>s</sub>), ball to-powder weight ratio (B:P), ball diameter (B<sub>d</sub>), and process control agent (PCA).</caption><thead><tr><th>Compounds</th><th>Device</th><th>Reactants</th><th>R<sub>mp</sub> [°C]</th><th>P<sub>H</sub> [MPa]</th><th>M<sub>t</sub> [h]</th><th>M<sub>s</sub> [rpm]</th><th>B:P</th><th>B<sub>d</sub> [mm]</th><th>PCA</th><th>Side product</th><th>Ref.</th></tr></thead><tbody><tr><td rowspan="15">Binary hydride</td><td>MgH<sub>2</sub></td><td>planetary</td><td>Mg</td><td>650</td><td>0.066</td><td>3</td><td>—</td><td>3</td><td>—</td><td>—</td><td>[466]</td></tr><tr><td>MgH<sub>2</sub></td><td>Planetary</td><td>Mg</td><td>650</td><td>2.4</td><td>47.5</td><td>—</td><td>12</td><td>—</td><td>—</td><td>[467]</td></tr><tr><td>MgH<sub>2</sub></td><td>Fritsch P6</td><td>Mg</td><td>650</td><td>1–9</td><td>18</td><td>500</td><td>10:1</td><td>10</td><td>—</td><td>[468]</td></tr><tr><td>MgH<sub>2</sub></td><td>Fritsch P4</td><td>Mg</td><td>650</td><td>8</td><td>4</td><td>400, 800</td><td>60:1</td><td>12</td><td>—</td><td>[155]</td></tr><tr><td>MgH<sub>2</sub></td><td>Spex 8000</td><td>Mg</td><td>650</td><td>1</td><td>3</td><td>—</td><td>2:1</td><td>—</td><td>—</td><td>[469]</td></tr><tr><td>MgH<sub>2</sub></td><td>Fritsch P4</td><td>Mg</td><td>650</td><td>7</td><td>4</td><td>400, 800</td><td>60:1</td><td>12</td><td>—</td><td>[470]</td></tr><tr><td>MgH<sub>2</sub></td><td>Planetary</td><td>Mg</td><td>650</td><td>Ar</td><td>35</td><td>250</td><td>20:1</td><td>10</td><td>Steric acid, toluene</td><td>[471]</td></tr><tr><td>MgH<sub>2</sub></td><td>Uni ball mill II</td><td>Mg</td><td>650</td><td>0.5</td><td>50</td><td>196</td><td>44:1</td><td>—</td><td>Graphite</td><td>MgO</td><td>[472]</td></tr><tr><td>MgH<sub>2</sub></td><td>Spex 8000</td><td>MgH<sub>2</sub></td><td>327</td><td>Ar</td><td>10</td><td>—</td><td>8:1</td><td>—</td><td>MgCl<sub>2</sub></td><td>MgO</td><td>[473]</td></tr><tr><td>MgH<sub>2</sub></td><td>Fritsch P5</td><td>MgH<sub>2</sub></td><td>327</td><td>Ar</td><td>100</td><td>—</td><td>10:1</td><td>10</td><td>MgO</td><td>—</td><td>[174a]</td></tr><tr><td>MgH<sub>2</sub></td><td>Spex 8000</td><td>MgH<sub>2</sub></td><td>327</td><td>Ar(0.6)</td><td>10</td><td>—</td><td>10:1</td><td>—</td><td>Nb<sub>2</sub>O<sub>5</sub></td><td>Mg<sub>x</sub>Nb<sub>y</sub>O<sub>z</sub>, MgO</td><td>[474]</td></tr><tr><td>TiH<sub>1.9</sub></td><td>Planetary</td><td>Ti</td><td>1668</td><td>2.4</td><td>5.5</td><td>—</td><td>—</td><td>12</td><td>—</td><td>—</td><td>[467]</td></tr><tr><td>TiH<sub>1.8</sub></td><td>Spex 8000</td><td>Ti</td><td>1668</td><td>0.2</td><td>3</td><td>—</td><td>—</td><td>—</td><td>—</td><td>—</td><td>[475]</td></tr><tr><td>TiH<sub>2</sub></td><td>Fritsch P4</td><td>Ti</td><td>1668</td><td>8</td><td>0.16</td><td>400, 800</td><td>60:1</td><td>12</td><td>—</td><td>—</td><td>[155]</td></tr><tr><td>TiH<sub>2</sub></td><td>Fritsch P4</td><td>Ti</td><td>1668</td><td>30</td><td>20</td><td>400</td><td>50:1</td><td>15</td><td>—</td><td>—</td><td>[476]</td></tr><tr><td>TiH<sub>2</sub></td><td>Fritsch P4</td><td>Ti</td><td>1668</td><td>7</td><td>4</td><td>400, 800</td><td>37.5:1</td><td>12</td><td>—</td><td>—</td><td>[470]</td></tr><tr><td>TiH<sub>2</sub></td><td>Fritsch P5</td><td>Ti</td><td>1668</td><td>—</td><td>10</td><td>300</td><td>10:1</td><td>10</td><td>Toluene</td><td>TiH<sub>x</sub></td><td>[477]</td></tr><tr><td>VH<sub>x</sub></td><td>Fritsch P5</td><td>V</td><td>1910</td><td>1</td><td>0.17</td><td>400</td><td>30:1</td><td>7</td><td>—</td><td>—</td><td>[478]</td></tr><tr><td rowspan="21">Ternary hydride</td><td>ZrNiH<sub>3</sub></td><td>Fritsch P5</td><td>ZrNi</td><td>—</td><td>2</td><td>3</td><td>—</td><td>30:1</td><td>10</td><td>—</td><td>—</td><td>[220]</td></tr><tr><td>ZrNiH</td><td>Fritsch P7</td><td>ZrNi</td><td>—</td><td>0.1</td><td>0.08</td><td>400</td><td>30:1</td><td>7</td><td>—</td><td>—</td><td>[479]</td></tr><tr><td>TiNiH<sub>3</sub></td><td>Rod-mill</td><td>Ti+Ni</td><td>1668, 1445</td><td>0.1</td><td>200</td><td>—</td><td>30:1</td><td>10</td><td>—</td><td>—</td><td>[480]</td></tr><tr><td>TiVH<sub>2.8</sub></td><td>Fritsch P5</td><td>TiV or Ti+V</td><td>1668, 1910</td><td>0.4</td><td>100</td><td>—</td><td>20:1</td><td>10</td><td>—</td><td>—</td><td>[481]</td></tr><tr><td>LaNi<sub>5</sub>H<sub>0.15</sub></td><td>Fritsch P7</td><td>LaNi<sub>5</sub></td><td>480</td><td>1</td><td>0.08</td><td>400</td><td>30:1</td><td>7</td><td>—</td><td>—</td><td>[482]</td></tr><tr><td>Mg<sub>2</sub>FeH<sub>6</sub></td><td>Fritsch P5</td><td>2Mg+Fe</td><td>650, 1538</td><td>1</td><td>20</td><td>325</td><td>4:1</td><td>10</td><td>—</td><td>MgH<sub>2</sub>, Fe</td><td>[483]</td></tr><tr><td>Mg<sub>2</sub>FeH<sub>6</sub></td><td>Spex 8000</td><td>2MgH<sub>2</sub>+Fe</td><td>327, 1538</td><td>0.1</td><td>60</td><td>—</td><td>10:1</td><td>10</td><td>—</td><td>Mg, MgO, Fe</td><td>[484]</td></tr><tr><td>Mg<sub>2</sub>CoH<sub>5</sub></td><td>BX254E</td><td>2MgH<sub>2</sub>+Co</td><td>327, 1495</td><td>0.1</td><td>10</td><td>—</td><td>8:1</td><td>4</td><td>—</td><td>—</td><td>[485]</td></tr><tr><td>Mg<sub>2</sub>CoH<sub>5</sub></td><td>Fritsch P6</td><td>2MgH<sub>2</sub>+Co</td><td>327, 1495</td><td>0.5</td><td>5</td><td>400</td><td>80:1</td><td>—</td><td>—</td><td>—</td><td>[486]</td></tr><tr><td>Mg<sub>2</sub>NiH<sub>4</sub></td><td>—</td><td>2MgH<sub>2</sub>+Ni</td><td>327, 1445</td><td>—</td><td>60</td><td>—</td><td>—</td><td>—</td><td>Graphite</td><td>—</td><td>[487]</td></tr><tr><td>Mg<sub>2</sub>NiH<sub>4</sub></td><td>Fritsch P5</td><td>2MgH<sub>2</sub>+Ni</td><td>327, 1445</td><td>0.5</td><td>22</td><td>325</td><td>4:1</td><td>10</td><td>—</td><td>Mg, MgH<sub>2</sub>, Ni</td><td>[488]</td></tr><tr><td>Mg<sub>2</sub>NiH<sub>4</sub></td><td>Fritsch P4</td><td>2MgH<sub>2</sub>+Ni</td><td>327, 1445</td><td>7.5</td><td>12</td><td>400</td><td>60:1</td><td>12</td><td>—</td><td>MgO</td><td>[489]</td></tr><tr><td>Mg<sub>2</sub>NiH<sub>1.8</sub></td><td>Fritsch P7</td><td>Mg<sub>2</sub>Ni</td><td>—</td><td>1</td><td>80</td><td>400</td><td>30:1</td><td>7</td><td>—</td><td>—</td><td>[490]</td></tr><tr><td>Mg<sub>0.8</sub>Ti<sub>0.2</sub></td><td>Spex 8000</td><td>Mg, Ti</td><td>327, 1668</td><td>Ar</td><td>16</td><td>—</td><td>—</td><td>—</td><td>Graphite</td><td>—</td><td>[491]</td></tr><tr><td>Mg<sub>0.75</sub>Ti<sub>0.25</sub></td><td>Spex 8000</td><td>Mg, Ti</td><td>327, 1668</td><td>Ar</td><td>12</td><td>—</td><td>—</td><td>—</td><td>Stearic acid</td><td>—</td><td>[491]</td></tr><tr><td>Al<sub>3</sub>Mg<sub>2</sub></td><td>Planetary</td><td>Al, Mg</td><td>660, 327</td><td>Ar</td><td>10</td><td>—</td><td>—</td><td>15</td><td>Stearic acid</td><td>Al, Mg</td><td>[492]</td></tr><tr><td>γ-Mg<sub>17</sub>Al<sub>12</sub></td><td>Retsch PM400</td><td>Mg, Al</td><td>327, 660</td><td>Ar</td><td>40</td><td>300</td><td>10:1</td><td>—</td><td>Methanol</td><td>MgH<sub>2</sub>, MgO, Mg<sub>28</sub>Al<sub>45</sub></td><td>[493]</td></tr><tr><td>Al<sub>3</sub>Zr</td><td>Fritsch P4</td><td>Al, Zr</td><td>660, 1855</td><td>Ar</td><td>5</td><td>300</td><td>20:1</td><td>10</td><td>Ethanol</td><td>Al, ZrH<sub>2</sub>, Zr, ZrC</td><td>[494]</td></tr><tr><td>Ti<sub>3</sub>Al</td><td>—</td><td>Ti, Al</td><td>1668, 660</td><td>Ar</td><td>300</td><td>—</td><td>—</td><td>—</td><td>Methanol, benzene</td><td>—</td><td>[495]</td></tr><tr><td>Ti<sub>x</sub>Al<sub>1-x</sub></td><td>Planetary</td><td>Ti, Al</td><td>1668, 660</td><td>Ar</td><td>100</td><td>—</td><td>10:1</td><td>10</td><td>Hexane</td><td>TiC, TiH<sub>2</sub>, Al<sub>2</sub>Ti<sub>4</sub>C<sub>2</sub></td><td>[496]</td></tr><tr><td>Sr<sub>2</sub>AlH<sub>7</sub></td><td>Fritsch P5</td><td>SrH<sub>2</sub>, Al</td><td>1050, 660</td><td>Ar</td><td>10</td><td>150</td><td>20:1</td><td>10</td><td>Stearic acid</td><td>Sr<sub>8</sub>Al<sub>7</sub>, Sr</td><td>[497]</td></tr><tr><td rowspan="3">Quaternary hydride</td><td>(Mg<sub>0.8</sub>Ti<sub>0.2</sub>)<sub>0.95</sub>Ni<sub>0.05</sub></td><td>Spex 8000</td><td>Mg, Ti, Ni</td><td>327, 1668, 1445</td><td>Ar</td><td>24</td><td>—</td><td>—</td><td>—</td><td>Graphite</td><td>—</td><td>[491]</td></tr><tr><td>(Mg<sub>0.65</sub>Ti<sub>0.35</sub>)<sub>0.95</sub>Ni<sub>0.05</sub></td><td>Spex 8000</td><td>Mg, Ti, Ni</td><td>327, 1668, 1445</td><td>Ar</td><td>28</td><td>—</td><td>—</td><td>—</td><td>Stearic acid</td><td>—</td><td>[491]</td></tr><tr><td>Mg<sub>1.9</sub>Al<sub>0.1</sub>Ni</td><td>Spex 8000</td><td>Mg, Al, Ni</td><td>327, 660, 1445</td><td>Ar</td><td>40</td><td>—</td><td>10:1</td><td>12</td><td>CCl<sub>4</sub></td><td>—</td><td>[498]</td></tr></tbody></table>
<table><caption>Table 3. Continued.</caption><thead><tr><th>Compounds</th><th>Device</th><th>Reactants</th><th>R<sub>mp</sub> [°C]</th><th>P<sub>H</sub> [MPa]</th><th>M<sub>t</sub> [h]</th><th>M<sub>s</sub> [rpm]</th><th>B:P</th><th>B<sub>d</sub> [mm]</th><th>PCA</th><th>Side product</th><th>Ref.</th></tr></thead><tbody><tr><td>Mg<sub>88.5</sub>Ni<sub>1.5</sub>La<sub>γ</sub></td><td>Planetary</td><td>Mg, Ni, La</td><td>327, 1445, 920</td><td>Ar(0.1)</td><td>10</td><td>200</td><td>30:1</td><td>—</td><td>Graphite</td><td>Mg, MgH<sub>2</sub>,<br>Mg<sub>2</sub>NiH<sub>0.3</sub>,<br>Mg<sub>2</sub>NiH<sub>4</sub>,LaH<sub>2</sub></td><td>[499]</td></tr><tr><td>(TiNi)<sub>0.7</sub>Mg<sub>0.3</sub></td><td>Planetary</td><td>Ti<sub>50</sub>Ni<sub>50</sub>, Mg</td><td>—, 327</td><td>Ar</td><td>40</td><td>400</td><td>4.4:1</td><td>20</td><td>Ethanol</td><td>Fe<sub>2</sub>TiO<sub>4</sub></td><td>[500]</td></tr><tr><td>Alanate</td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td></tr><tr><td>LiAlH<sub>4</sub></td><td>Fritsch P5</td><td>LiH+Al</td><td>692, 660</td><td>1</td><td>24</td><td>400</td><td>—</td><td>9.5</td><td>—</td><td>Li<sub>3</sub>AlH<sub>6</sub></td><td>[501]</td></tr><tr><td>Li<sub>3</sub>AlH<sub>6</sub></td><td>—</td><td>2LiH+LiAlH<sub>4</sub></td><td>692, 150</td><td>—</td><td>5</td><td>—</td><td>30:1-42:1</td><td>—</td><td>—</td><td>—</td><td>[296b]</td></tr><tr><td>LiMg(AlH<sub>4</sub>)<sub>3</sub></td><td>Retsch</td><td>3LiAlH<sub>4</sub>+MgCl<sub>2</sub></td><td>150, 714</td><td>—</td><td>3</td><td>—</td><td>11.3:1</td><td>—</td><td>—</td><td>LiCl</td><td>[502]</td></tr><tr><td>NaAlH<sub>4</sub></td><td>Fritsch P7</td><td>NaH+Al</td><td>300, 660</td><td>8.3</td><td>4</td><td>500</td><td>50:1</td><td>—</td><td>—</td><td>NaCl,Al</td><td>[503]</td></tr><tr><td>NaAlH<sub>4</sub></td><td>Fritsch P5</td><td>NaH+Al</td><td>300, 660</td><td>0.6</td><td>120</td><td>350</td><td>50:1</td><td>10</td><td>—</td><td>Na<sub>3</sub>AlH<sub>6</sub>,NaH,Al</td><td>[504]</td></tr><tr><td>NaAlH<sub>4</sub></td><td>Fritsch P5</td><td>NaH+Al</td><td>300, 660</td><td>1.2</td><td>240</td><td>230</td><td>10:1</td><td>10</td><td>—</td><td>NaCl,Al</td><td>[504]</td></tr><tr><td>NaAlH<sub>4</sub></td><td>Planetary</td><td>NaH+Al</td><td>300, 660</td><td>3</td><td>60</td><td>350</td><td>30:1</td><td>8</td><td>—</td><td>Na<sub>3</sub>AlH<sub>6</sub>,Al</td><td>[505]</td></tr><tr><td>Na<sub>3</sub>AlH<sub>6</sub></td><td>Planetary</td><td>NaH+Al</td><td>300, 660</td><td>0.5</td><td>30</td><td>350</td><td>30:1</td><td>8</td><td>—</td><td>Al</td><td>[505]</td></tr><tr><td>Na<sub>3</sub>AlH<sub>6</sub></td><td>Fritsch P4</td><td>3NaH+Al</td><td>300, 660</td><td>10</td><td>8</td><td>400</td><td>90:1</td><td>15</td><td>—</td><td>NaCl,NaH,Al</td><td>[225c]</td></tr><tr><td>Na<sub>3</sub>AlH<sub>6</sub></td><td>Spex 8000</td><td>2NaH+NaAlH<sub>4</sub></td><td>300, 183</td><td>0.2</td><td>20</td><td>—</td><td>10:1</td><td>12.7,14.3</td><td>—</td><td>Al</td><td>[506]</td></tr><tr><td>Na<sub>2</sub>LiAlH<sub>6</sub></td><td>Spex 8000</td><td>NaH+LiH+NaAlH<sub>4</sub></td><td>300, 692, 183</td><td>0.2</td><td>20</td><td>—</td><td>10:1</td><td>12.7,14.3</td><td>—</td><td>Al</td><td>[506]</td></tr><tr><td>Na<sub>2</sub>LiAlH<sub>6</sub></td><td>Fritsch P4</td><td>2NaH+LiH+Al</td><td>300, 692, 660</td><td>10</td><td>4</td><td>400</td><td>90:1</td><td>15</td><td>—</td><td>NaCl</td><td>[225c]</td></tr><tr><td>NaLi<sub>2</sub>AlH<sub>6</sub></td><td>Fritsch P4</td><td>NaH+2LiH+Al</td><td>300, 692, 660</td><td>10</td><td>4</td><td>400</td><td>90:1</td><td>15</td><td>—</td><td>NaCl,NaAlH<sub>4</sub>,Al</td><td>[225c]</td></tr><tr><td>KAlH<sub>4</sub></td><td>Fritsch P7</td><td>KH+Al</td><td>417, 660</td><td>10</td><td>10</td><td>400</td><td>—</td><td>—</td><td>—</td><td>KH,Al</td><td>[502]</td></tr><tr><td>KAlH<sub>4</sub></td><td>Fritsch P7</td><td>KH+Al</td><td>417, 660</td><td>>17.5</td><td>0.5</td><td>—</td><td>—</td><td>10,5</td><td>—</td><td>Al</td><td>[507]</td></tr><tr><td>K<sub>2</sub>LiAlH<sub>6</sub></td><td>SPEX8000</td><td>2KH+LiAlH<sub>4</sub></td><td>417, 150</td><td>—</td><td>0.5</td><td>—</td><td>27:1</td><td>—</td><td>—</td><td>—</td><td>[225a]</td></tr><tr><td>K<sub>2</sub>NaAlH<sub>6</sub></td><td>Fritsch P4</td><td>2KH+NaAlH<sub>4</sub></td><td>417, 183</td><td>1</td><td>25</td><td>150</td><td>40:1</td><td>—</td><td>—</td><td>—</td><td>[508]</td></tr><tr><td>Mg(AlH<sub>4</sub>)<sub>2</sub></td><td>Magneto</td><td>2NaAlH<sub>4</sub>+MgCl<sub>2</sub></td><td>183, 714</td><td>0.6(Ar)</td><td>40</td><td>225</td><td>40:1</td><td>—</td><td>—</td><td>NaCl,Al</td><td>[509]</td></tr><tr><td>Mg(AlH<sub>4</sub>)<sub>2</sub></td><td>Fritsch P7</td><td>2AlH<sub>3</sub>+MgH<sub>2</sub></td><td>150, 327</td><td>0.1</td><td>4</td><td>400</td><td>175:1</td><td>10</td><td>—</td><td>AlH<sub>3</sub>,MgH<sub>2</sub>,Al</td><td>[510]</td></tr><tr><td>Mg(AlH<sub>4</sub>)<sub>2</sub></td><td>Spex 8000</td><td>2NaAlH<sub>4</sub>+MgCl<sub>2</sub></td><td>183, 714</td><td>—</td><td>1</td><td>—</td><td>35:1</td><td>7.9</td><td>—</td><td>NaCl</td><td>[511]</td></tr><tr><td>Mg(AlH<sub>4</sub>)<sub>2</sub></td><td>Planetary</td><td>2NaAlH<sub>4</sub>+MgCl<sub>2</sub></td><td>183, 714</td><td>—</td><td>3</td><td>500</td><td>60:1</td><td>—</td><td>—</td><td>NaCl</td><td>[512]</td></tr><tr><td>Ca(AlH<sub>4</sub>)<sub>2</sub></td><td>Planetary</td><td>2NaAlH<sub>4</sub>+CaCl<sub>2</sub></td><td>183, 772</td><td>1</td><td>48</td><td>200</td><td>40:1</td><td>—</td><td>—</td><td>CaH<sub>2</sub>,Al,NaCl</td><td>[513]</td></tr><tr><td>CaAlH<sub>5</sub></td><td>Fritsch P7</td><td>CaH<sub>2</sub>+AlH<sub>3</sub></td><td>816, 150</td><td>0.1</td><td>3</td><td>400</td><td>80:1</td><td>10</td><td>—</td><td>CaH<sub>2</sub>,Al</td><td>[510]</td></tr><tr><td>Borohydride</td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td></tr><tr><td>KBH<sub>4</sub></td><td>Fritsch P7</td><td>2MgH<sub>2</sub>+KBO<sub>2</sub></td><td>327, 947</td><td>—</td><td>2</td><td>490</td><td>—</td><td>13</td><td>—</td><td>MgO</td><td>[514]</td></tr><tr><td>NaBH<sub>4</sub></td><td>—</td><td>4MgH<sub>2</sub>+Na<sub>2</sub>B<sub>4</sub>O<sub>7</sub></td><td>327, 743</td><td>—</td><td>2</td><td>2750</td><td>—</td><td>13</td><td>—</td><td>MgO</td><td>[515]</td></tr><tr><td>NaBH<sub>4</sub></td><td>QM-3C</td><td>2MgH<sub>2</sub>+NaBO<sub>2</sub></td><td>327, 966</td><td>—</td><td>4</td><td>—</td><td>20:1</td><td>—</td><td>—</td><td>MgO</td><td>[393]</td></tr><tr><td>Mg(BH<sub>4</sub>)<sub>2</sub></td><td>Fritsch P7</td><td>5MgCl<sub>2</sub>+4NaBH<sub>4</sub></td><td>714, ≈500</td><td>0.1(Ar)</td><td>2</td><td>—</td><td>—</td><td>7</td><td>—</td><td>NaCl</td><td>[516]</td></tr><tr><td>Zn(BH<sub>4</sub>)<sub>2</sub></td><td>Spex 8000</td><td>ZnCl<sub>2</sub>+2NaBH<sub>4</sub></td><td>290, ≈500</td><td>—</td><td>0.5</td><td>—</td><td>35:1</td><td>7.9</td><td>—</td><td>NaCl</td><td>[517]</td></tr><tr><td>Mn(BH<sub>4</sub>)<sub>2</sub></td><td>Planetary</td><td>3LiBH<sub>4</sub>+MnCl<sub>2</sub></td><td>275, 654</td><td>—</td><td>6</td><td>175</td><td>32:1</td><td>12</td><td>—</td><td>LiCl</td><td>[518]</td></tr><tr><td>Ca(BH<sub>4</sub>)<sub>2</sub></td><td>Spex 8000</td><td>CaH<sub>2</sub>+MgB<sub>2</sub></td><td>816, 830</td><td>—</td><td>1.5</td><td>—</td><td>10:1</td><td>—</td><td>—</td><td>MgH<sub>2</sub></td><td>[262a]</td></tr><tr><td>K<sub>2</sub>Mn(BH<sub>4</sub>)<sub>4</sub></td><td>Planetary</td><td>2KBH<sub>4</sub>+MnCl<sub>2</sub></td><td>585, 654</td><td>—</td><td>6</td><td>175</td><td>32:1</td><td>14</td><td>—</td><td>KMnCl<sub>3</sub></td><td>[231a]</td></tr><tr><td>LiZn<sub>2</sub>(BH<sub>4</sub>)<sub>5</sub></td><td>—</td><td>2LiBH<sub>4</sub>+ZnCl<sub>2</sub></td><td>275, 290</td><td>Ar</td><td>2</td><td>—</td><td>—</td><td>—</td><td>—</td><td>LiCl</td><td>[229]</td></tr><tr><td>NaZn<sub>2</sub>(BH<sub>4</sub>)<sub>5</sub></td><td>—</td><td>3NaBH<sub>4</sub>+ZnCl<sub>2</sub></td><td>≈500, 290</td><td>Ar</td><td>2</td><td>—</td><td>—</td><td>—</td><td>—</td><td>NaCl</td><td>[229]</td></tr><tr><td>NaZn(BH<sub>4</sub>)<sub>3</sub></td><td>—</td><td>4NaBH<sub>4</sub>+ZnCl<sub>2</sub></td><td>≈500, 290</td><td>Ar</td><td>2</td><td>—</td><td>—</td><td>—</td><td>—</td><td>NaCl</td><td>[229]</td></tr><tr><td>Mg<sub>8</sub>Mn<sub>(1-x)</sub>(BH<sub>4</sub>)<sub>2</sub></td><td>Fritsch P7</td><td>LiBH<sub>4</sub>,MgCl<sub>2</sub>,MnCl<sub>2</sub></td><td>275, 714, 354</td><td>—</td><td>5.83</td><td>600</td><td>25:1</td><td>15</td><td>—</td><td>LiCl,Li<sub>2</sub>MgCl<sub>4</sub></td><td>[519]</td></tr><tr><td>Li<sub>3</sub>MgZn<sub>5</sub>(BH<sub>4</sub>)<sub>15</sub></td><td>Fritsch P7</td><td>LiBH<sub>4</sub>,MgCl<sub>2</sub>,ZnCl<sub>2</sub></td><td>275, 714, 290</td><td>Ar</td><td>5.83</td><td>600</td><td>25:1</td><td>15,12,10</td><td>—</td><td>Li<sub>2</sub>MgCl<sub>4</sub></td><td>[230]</td></tr><tr><td>Li<sub>3</sub>MnZn<sub>5</sub>(BH<sub>4</sub>)<sub>15</sub></td><td>Fritsch P7</td><td>LiBH<sub>4</sub>,MnCl<sub>2</sub>,ZnCl<sub>2</sub></td><td>275, 654, 290</td><td>Ar</td><td>3</td><td>600</td><td>25:1</td><td>15,12,10</td><td>—</td><td>Li<sub>2</sub>MnCl<sub>4</sub></td><td>[230]</td></tr><tr><td>Amide</td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td></tr><tr><td>LiNH<sub>2</sub></td><td>Fritsch P7</td><td>LiH</td><td>692</td><td>0.4(NH<sub>3</sub>)</td><td>2</td><td>400</td><td>—</td><td>—</td><td>—</td><td>—</td><td>[520]</td></tr><tr><td>NaNH<sub>2</sub></td><td>Fritsch P7</td><td>NaH</td><td>300</td><td>0.4(NH<sub>3</sub>)</td><td>1</td><td>400</td><td>—</td><td>—</td><td>—</td><td>—</td><td>[520]</td></tr><tr><td>Mg(NH<sub>2</sub>)<sub>2</sub></td><td>Fritsch P7</td><td>MgH<sub>2</sub></td><td>327</td><td>0.4(NH<sub>3</sub>)</td><td>13</td><td>400</td><td>—</td><td>—</td><td>—</td><td>—</td><td>[520]</td></tr><tr><td>Ca(NH<sub>2</sub>)<sub>2</sub></td><td>RM-10</td><td>CaH<sub>2</sub></td><td>816</td><td>0.5(NH<sub>3</sub>)</td><td>18</td><td>—</td><td>—</td><td>7</td><td>—</td><td>CaH<sub>2</sub></td><td>[521]</td></tr><tr><td>Ca(NH<sub>2</sub>)<sub>2</sub></td><td>Planetary</td><td>Ca</td><td>842</td><td>—</td><td>24</td><td>—</td><td>30:1</td><td>—</td><td>—</td><td>—</td><td>[522]</td></tr></tbody></table>
<table><caption>**Table 3. Continued.**</caption><thead><tr><th>Compounds</th><th>Device</th><th>Reactants</th><th>R<sub>mp</sub> [°C]</th><th>P<sub>H</sub> [MPa]</th><th>M<sub>t</sub> [h]</th><th>M<sub>s</sub> [rpm]</th><th>B:P</th><th>B<sub>d</sub> [mm]</th><th>PCA</th><th>Side product</th><th>Ref.</th></tr></thead><tbody><tr><td rowspan="6">Imide</td><td>Li<sub>2</sub>NH</td><td>Fritsch P7</td><td>LiNH<sub>2</sub>+LiH</td><td>390, 692</td><td>1</td><td>2</td><td>400</td><td>—</td><td>7</td><td>—</td><td>—</td><td>[221c]</td></tr><tr><td>CaNH</td><td>Fritsch P7</td><td>Ca(NH<sub>2</sub>)<sub>2</sub>+3CaH<sub>2</sub></td><td>—,816</td><td>1</td><td>2</td><td>400</td><td>—</td><td>7</td><td>—</td><td>CaH<sub>2</sub></td><td>[521]</td></tr><tr><td>Ca<sub>2</sub>NH</td><td>—</td><td>Ca<sub>3</sub>N<sub>2</sub>+CaH<sub>2</sub></td><td>—,816</td><td>—</td><td>—</td><td>—</td><td>—</td><td>—</td><td>—</td><td>—</td><td>[283]</td></tr><tr><td>MgNH</td><td>—</td><td>Mg(NH<sub>2</sub>)<sub>2</sub>+MgH<sub>2</sub></td><td>—,714</td><td>—</td><td>72</td><td>—</td><td>26:1</td><td>10</td><td>—</td><td>—</td><td>[523]</td></tr><tr><td>Li<sub>2</sub>Ca(NH<sub>2</sub>)<sub>2</sub></td><td>Fritsch P7</td><td>2LiNH<sub>2</sub>+CaH<sub>2</sub></td><td>390, 816</td><td>—</td><td>1</td><td>400</td><td>—</td><td>—</td><td>—</td><td>—</td><td>[524]</td></tr><tr><td>Li<sub>2</sub>Mg(NH<sub>2</sub>)<sub>2</sub></td><td>Planetary</td><td>LiNH<sub>2</sub>+2MgH<sub>2</sub></td><td>390, 714</td><td>Ar</td><td>36</td><td>—</td><td>—</td><td>—</td><td>—</td><td>—</td><td>[525]</td></tr><tr><td rowspan="6">Metal amidoborane</td><td>LiNH<sub>2</sub>BH<sub>3</sub></td><td>Fritsch P7</td><td>LiH+NH<sub>3</sub>BH<sub>3</sub></td><td>692, 104</td><td>—</td><td>2</td><td>200</td><td>40:1</td><td>10</td><td>—</td><td>—</td><td>[526]</td></tr><tr><td>NaNH<sub>2</sub>BH<sub>3</sub></td><td>Planetary</td><td>NaH+NH<sub>3</sub>BH<sub>3</sub></td><td>300, 104</td><td>—</td><td>—</td><td>200</td><td>—</td><td>—</td><td>—</td><td>—</td><td>[236]</td></tr><tr><td>Ca(NH<sub>2</sub>BH<sub>3</sub>)<sub>2</sub></td><td>Fritsch P7</td><td>CaH<sub>2</sub>+2NH<sub>3</sub>BH<sub>3</sub></td><td>816, 104</td><td>0.1 (He)</td><td>1</td><td>200</td><td>—</td><td>—</td><td>—</td><td>—</td><td>[527]</td></tr><tr><td>LiNa(NH<sub>2</sub>BH<sub>3</sub>)<sub>2</sub></td><td>—</td><td>LiH+NaH+2NH<sub>3</sub>BH<sub>3</sub></td><td>692, 300, 104</td><td>—</td><td>—</td><td>—</td><td>—</td><td>—</td><td>—</td><td>—</td><td>[237]</td></tr><tr><td>Na<sub>2</sub>Mg(NH<sub>2</sub>BH<sub>3</sub>)<sub>2</sub></td><td>—</td><td>NaH,MgH<sub>2</sub>,NH<sub>3</sub>BH<sub>3</sub></td><td>300, 714, 104</td><td>—</td><td>—</td><td>—</td><td>—</td><td>—</td><td>—</td><td>—</td><td>[234a]</td></tr><tr><td>KAl(NH<sub>2</sub>BH<sub>3</sub>)<sub>4</sub></td><td>Fritsch P6</td><td>KAlH<sub>4</sub>+ NH<sub>3</sub>BH<sub>3</sub></td><td>—,104</td><td>12</td><td>12</td><td>—</td><td>30:1</td><td>10</td><td>—</td><td>KAlH<sub>4</sub>,NH<sub>3</sub>BH<sub>3</sub></td><td>[235e]</td></tr></tbody></table>

of Fe contamination (>5 wt%) into the milled material.[164a] Using less energetic milling machines or “softer milling” tools can help reduce the level of contamination from the milling tools. For example, Ares et al. reported a significant reduction of contamination by ball milling MgH₂ on a planetary mill with zirconia milling tools (Zr contamination ≈0.5 mass%) instead of stainless steel vials and balls (Fe contamination ≈ 4 mass%).[169]

Contamination is also influenced by the length of the milling process—longer milling times leading to higher levels of contamination—and the size of the balls used and remaining empty space in the vial. If the size of individual balls is too large as compared to the capacity of the milling vial, then the interaction between ball/powder is going to be low and this leads to a low milling efficiency. Similarly, the vial should contain sufficient powder so there is enough ball/powder interaction. If the vial is overfilled, then there is not enough space for the balls to move. As a rule of thumb the vial should be filled at 50% to enable an appropriate milling of the powder and minimize direct ball to ball/wall impact leading to some partial fragmentation of the milling tools and a lower milling efficiency. In this configuration, the ball to powder mass ratio is often taken as 10:1 in vials of small capacity (≈50 mL). A larger ball to powder mass ratio of 50:1 or 100:1 may be used in a large capacity attritor mill. The ball to powder mass ratio has a significant effect on the milling time. The higher this ratio the shorter the time (Figure 17c).[170]

To limit cold-welding and facilitate further powder refinement, a process control agent can also be added during milling. This process control agent can be a solid, liquid or gas and it is believed to reduce the surface energy created at freshly made surfaces, hence limiting particle agglomeration and thus cold-welding upon impact. Several process control agents in the form of hydrocarbons, surfactants, and carbon materials have been used in the range of 1–5 mass% of the total powder charge (Table 3).[171] However, these compounds may decompose during milling and bring oxygen and/or carbon to the milled powders resulting in the formation of oxides and car-

bides. Hence the choice of the process control agent depends upon the acceptable purity of the final material. In the case of materials for hydrogen storage, this choice is critical since this may lead to serious deterioration of material's properties. Excessive surface contamination of the hydride powder may delay the hydrogen release/uptake at the core of the particles. In this case, hydrogen cycling at high temperatures is required to overcome the oxide or carbon layers. For example, in the case of MgH₂,[172] the formation of an oxide passivation layer of ≈4 nm delays the hydrogen release by ≈50 °C. Using carbon materials may also facilitate the refinement of MgH₂ particles but in this case this will lead to the release of CO (a poison for PEM fuel cells) and C–H contaminants in the hydrogen stream.[173] We reported several years ago the use of oxides and in particular MgO as effective process control agents leading to very fast hydrogen sorption kinetics comparable to the best catalysts reported so far with pure hydrogen release only.[174] It should be noted that no universal process control agent exists and this intrinsically depends upon the chemical nature of the material to be milled. Figure 17d summarizes the main steps toward the synthesis/modification of materials for hydrogen storage. Furthermore, in the case of hydrogen storage materials, mechanical milling must be carried out under a very well controlled atmosphere (inert argon or hydrogen) to minimize the formation of oxide or nitrides.

## 5.2. Modification of the Properties of Hydrides via Mechanical Milling

By using mechanical milling on materials for hydrogen storage, one of the aims is to reduce the hydrogen-metal bond strength by making new alloys, nanostructures and/or creating fresh surfaces and a large amount of grain boundaries to improve the hydrogen kinetics.[157a] An underlying assumption is that mechanical alloying can also lead to the formation of metastable phases, i.e., a nonequilibrium state resulting from the “energizing and rapid quenching of the nonequilibrium phase”
upon repeated and severe plastic deformations (Figure 17b).[175]
With the reports on the drastic improvements of the hydrogen kinetics of the Mg/MgH₂ systems[150,151b,176] and the demonstrated hydrogen reversibility in NaAlH₄ with Ti doping,[63,65a,177] mechanical milling has also been revealed to be an effective technique to finely disperse/react transition metals and other elements or alloys that are believed to act as catalysts or for the synthesis of novel hydride materials.[46c,52a,151a,178]

### 5.2.1. Nanostructured Hydrides

Nanocrystalline materials, characterized by a crystal size <100 nm, can be formed by techniques including vapor condensation,[179] sputtering,[180] thermal decomposition,[181] chemical precipitation,[182] and so on.[183] But the advantage of mechanical milling is in the possibility of producing bulk quantities of material. However, from a commercial perspective, mechanical milling remains an expensive production method and alternatives in the forms of high pressure torsions, equal channel angular pressing, and cold rolling have been proposed.[184]

Nanocrystalline materials are typically described by a mixture of nanodomains composed of lattice atoms with some long range order and boundary atoms with a random interatomic arrangement separating individual nanograins. In nanocrystalline materials, the number of nanocrystalline domain/grain boundaries interfaces can be as large as $6 \times 10^{25} \text{ m}^{-3}$ for grain sizes of 10 nm, and by assuming that the grains are spherical or cubic, the volume fraction of nanocrystalline material (C) can be determined as[185]

$$C = 3\delta/d \qquad (37)$$

where $\delta$ and $d$ correspond to the average grain thickness and diameter, respectively.

This means that the volume fraction of atoms in the grains boundary can be as high as 50% for 5 nm grains. Such an evolution in crystallinity can be determined directly by transmission electron microscopy imaging or indirectly from the broadening of the diffraction peaks width as observed by X-ray diffraction analysis. Assuming appropriate correction from instrumental and strain effects, the average grain size can also be calculated with the Scherrer equation from the diffraction pattern.[186]

The severe plastic deformation occurring during high energy milling generally leads to the formation of nanocrystalline structures. This has been observed in many materials (e.g., metals, intermetallics, and alloys), and thus can be considered as a general trend. The exact mechanism leading to nanostructuration remains unclear, but it is expected that the disintegration of the coarse grains is the result of an increase in the density of dislocations, vacancies and stacking faults within the heavily strained regions formed under severe plastic deformation.[187] As these relax along the shear bands, nanocrystalline grains randomly oriented and free form dislocations are formed. The energy transferred from the mechanical milling process to the material can then be considered to be stored in the form of atomic disorder and/or grain boundaries and follow

$$\Delta G(\text{milling}) = \Delta G(\text{disorder}) + \Delta G(\text{grain boundaries}) \qquad (38)$$

It should be noted that the formation of nanostructures has mainly been observed in metallic hydrides. For complex hydrides such a nanostructuration has not been observed owing to the low hardness/ductility of the powder (Table 4).[188] During the process, many parameters can influence the properties of the final products including the milling temperature, milling time, grinding ball diameter, ball-to-powder ratio, and process control agent. Therefore, the determination of the achievable grain size is not straightforward.[153,189] However, Li et al. found that in most of the cases, the rate of refinement of the grain size roughly shows a power-law relationship with milling processing time (Figure 17c)[190]

$$d = Kt^{-2/3} \qquad (39)$$

where $d$ is the grain size, $t$ is the milling time, and $K$ is a constant.

Mohamed also proposed a dislocation model whereby the minimum achievable grain size can be related to the hardness, melting temperature, and stacking fault energy of the material to be milled.[191] The minimum grain size achievable was found to scale with the inverse of the melting temperature of fcc metals including Al, Ag, Cu, and Ni. However, no similar trend was found for fcc metals with a high melting point and bcc, hcp metals.[192] Nanocrystalline materials display higher mechanical strength, hardness, and extremely high diffusion rates as compared to their microcrystalline counterparts,[164b,185,193] and these are believed to lead to unique physical, mechanical, and magnetic properties. In the case of metals for hydrogen storage, the lower atomic packing density at the grain boundaries can facilitate the hydrogen solubility[194] and diffusion,[27m] and thus provides additional hydrogen “paths” for the rapid hydrogenation of materials. This is particularly the case for Pd, Ni, and MgH₂.[151a,195] It is, however, important to mention that upon the first hydrogen cycling, the nanocrystalline structure will undergo a rapid restructuration and crystalline growth owing to the annihilation of defects,[46c,174a,196] the crystalline structure evolution upon hydrogen uptake and release (e.g., 30% expansion/contraction upon Mg/MgH₂ conversion), and the driving force to reduce interfacial energy at grain boundaries. The latter is proportional to the grain boundary area and thus is relatively high in nanocrystalline materials, i.e., $\approx 4 \text{ J m}^{-2}$ instead of the $1 \text{ J m}^{-2}$ observed in polycrystalline materials.[197]

Typically, upon mechanical milling crystallite sizes can be reduced from several µm to $\approx 10-50 \text{ nm}$.[46c,157a,174] This crystallite size will increase back to $>80 \text{ nm}$ upon hydrogen cycling,[198] and this can affect the hydrogen kinetics of hydrides. For example, upon hydrogen cycling, nanocrystalline MgH₂ shows kinetics in the order of 30 and 120 min for hydrogen release and uptake, respectively.[151a,174a,199] However, as the grain size grows from 9 to $>100 \text{ nm}$, hydrogen kinetics have been observed to slow down by a factor 2 to 5 after 100 cycles.[200]

Strategies to maintain the nanocrystalline structure of metal hydrides remains unclear, although it has been claimed that the addition of particles in the form of TiH₂,[155,201] MgO,[174] or Nb₂O₅[202] may reduce crystalline growth upon hydrogen cycling. Hence in addition to catalytic activity, reduced forms of Nb₂O₅ embedded within the MgH₂/Mg domains may also limit the restructuration of grain boundaries.[202c] In contrast, MgO cannot catalyze the
<table><caption>**Table 4.** Summary of the hardness of selected elements/compounds relevant for the synthesis of hydrides.</caption><thead><tr><th colspan="4">Elements/compounds: hardness [Mohs]<sup>a)</sup></th></tr></thead><tbody><tr><td>Al: 2.2–2.9<sup>[529]</sup></td><td>Al<sub>2</sub>O<sub>3</sub>: 8–9<sup>[530]</sup></td><td>N: –</td><td>NH<sub>4</sub>Cl: 2<sup>[531]</sup></td></tr><tr><td>Ag: 2.7<sup>[529]</sup></td><td>AgCl: 2.5<sup>[531]</sup></td><td>Na: 0.5<sup>[532]</sup></td><td>NaF: 2–2.5<sup>[531]</sup></td></tr><tr><td>Au: 2.5<sup>[529]</sup></td><td>–</td><td></td><td>NaCl: 2<sup>[531]</sup></td></tr><tr><td>C: 0.5–10<sup>[528,533]</sup></td><td>Graphite: 0.5–1<sup>[533a]</sup></td><td>Nb: 6<sup>[534]</sup></td><td>Nb<sub>2</sub>O<sub>5</sub>: 5–6<sup>[535]</sup></td></tr><tr><td>Ce: 2.5<sup>[532]</sup></td><td>CeO<sub>2</sub>: 6<sup>[536]</sup></td><td>Ni: 3.5–5<sup>[529]</sup></td><td>NiO: 5.5<sup>[531]</sup></td></tr><tr><td>Ca: 1.75<sup>[532]</sup></td><td>CaF<sub>2</sub>: 4<sup>[531]</sup></td><td>Os: 7<sup>[529]</sup></td><td>–</td></tr><tr><td>Cd: 2<sup>[529]</sup></td><td>–</td><td>Pb: 1.5<sup>[529]</sup></td><td>–</td></tr><tr><td>Cu: 3<sup>[529]</sup></td><td>CuO: 3.5<sup>[530]</sup></td><td>Pd: 4.5–5<sup>[537]</sup></td><td>–</td></tr><tr><td>–</td><td>Cu<sub>2</sub>O: 3.5–4<sup>[529]</sup></td><td>Pt: 4–4.5<sup>[537]</sup></td><td>–</td></tr><tr><td>Co: 5<sup>[529]</sup></td><td>–</td><td>Ta: 7<sup>[529]</sup></td><td>–</td></tr><tr><td>Cr: 4.5–8<sup>[529]</sup></td><td>Cr<sub>2</sub>O<sub>3</sub>: 9<sup>[531]</sup></td><td>Ti: 6<sup>[529]</sup></td><td>TiH<sub>2</sub>: 7.8<sup>[533a]</sup></td></tr><tr><td rowspan="5">Fe: 3.5–4.5<sup>[529]</sup></td><td>Fe<sub>2</sub>O<sub>3</sub>: 6<sup>[531]</sup></td><td></td><td>TiO<sub>2</sub> (anatase): 5.5–6<sup>[538]</sup></td></tr><tr><td>Fe<sub>3</sub>O<sub>4</sub>: 6<sup>[531]</sup></td><td></td><td>TiB: 9<sup>[531]</sup></td></tr><tr><td>FeTiO<sub>3</sub>: 5–6<sup>[539]</sup></td><td></td><td>TiC: 8–9<sup>[529]</sup></td></tr><tr><td>FeAl<sub>2</sub>O<sub>4</sub>: 7.5–8<sup>[531]</sup></td><td>Sb: 3<sup>[532]</sup></td><td>–</td></tr><tr><td>KCl: 2<sup>[540]</sup></td><td>Si: 7<sup>[529]</sup></td><td>SiO<sub>2</sub>: 6.5–7<sup>[531]</sup></td></tr><tr><td rowspan="5">La: 2.5<sup>[541]</sup></td><td>LaNi<sub>5</sub>: 3.5<sup>[541]</sup></td><td></td><td>SiC: 9<sup>[542]</sup></td></tr><tr><td>La<sub>2</sub>O<sub>3</sub>: 8<sup>[536]</sup></td><td>Sn: 2<sup>[529]</sup></td><td>SnO<sub>2</sub>: 6–7<sup>[531]</sup></td></tr><tr><td>LaNi<sub>3.5</sub>CoAl<sub>0.5</sub>: 5–6<sup>[543a]</sup></td><td>Sr: 1.5<sup>[532]</sup></td><td>SrTiO<sub>3</sub>: 6–6.5<sup>[544]</sup></td></tr><tr><td>LaNi<sub>4.5</sub>Al<sub>0.5</sub>: ≈6<sup>[543a]</sup></td><td>V: 7<sup>[532]</sup></td><td>V<sub>2</sub>C<sub>3</sub>: 8<sup>[529]</sup></td></tr><tr><td>LaNi<sub>4</sub>Co: ≈6<sup>[543]</sup></td><td></td><td>VC: 9–10<sup>[529]</sup></td></tr><tr><td rowspan="2">Li: 0.6<sup>[532]</sup></td><td>LiH: 3.5<sup>[545]</sup></td><td>W: 7<sup>[529]</sup></td><td>WC: 8–9<sup>[531]</sup></td></tr><tr><td>LiNH<sub>2</sub>: ≈6.5<sup>[546]</sup></td><td>Zn: 2.5<sup>[529]</sup></td><td>ZnO: 4<sup>[547]</sup></td></tr><tr><td rowspan="4">Mg: 2.5<sup>[529]</sup></td><td>MgH<sub>2</sub>: 4<sup>[198b]</sup></td><td>Zr: 7–7.5<sup>[548]</sup></td><td>ZrFe<sub>2</sub>: ≈8.3<sup>[546]</sup></td></tr><tr><td></td><td></td><td>ZrO<sub>2</sub>: 8–8.5<sup>[549]</sup></td></tr><tr><td></td><td></td><td>Zr<sub>0.2</sub>La<sub>0.8</sub>Ni<sub>4.5</sub>Al<sub>0.5</sub>: 5.56<sup>[543a]</sup></td></tr><tr><td>MgO: 6<sup>[529]</sup></td><td></td><td></td></tr><tr><td rowspan="4">Mn: 7<sup>[529]</sup></td><td>MgSiO<sub>3</sub>: 5–6<sup>[531]</sup></td><td></td><td></td></tr><tr><td>MnO<sub>2</sub>: 6–6.5<sup>[531]</sup></td><td>Mischmetal</td><td>MmNi<sub>4.6</sub>Mn<sub>0.4</sub>: 5–5.5<sup>[550a]</sup></td></tr><tr><td>Mn<sub>2</sub>O<sub>3</sub>: 6.5<sup>[529]</sup></td><td></td><td>MmNi<sub>4.6</sub>Mn<sub>0.8</sub>: ≈5.5<sup>[550a]</sup></td></tr><tr><td>Mn<sub>3</sub>O<sub>4</sub>: 5–5.5<sup>[531]</sup></td><td>–</td><td>–</td></tr></tbody></table>

dissociation/recombination of hydrogen at the H₂/Mg/MgH₂ interfaces and thus its role is purely limited to that of grain boundaries stabilization. As observed in several metallic nanocrystalline structures,[203] the activation energy for the grain growth is comparable to the activation energy associated with lattice diffusion and isothermal grain growth can be modeled following the equation

$$d^g - d_0^g = Kt \qquad (40)$$

where $d_0$ is the initial grain size and $d$ is the grain size at time $t$, and at a constant temperature, $K$ is a constant, and $g$ (value varying from 2 to 10)[185] the grain growth exponent.

It is also noteworthy that nanostructuration does not seem to affect the thermodynamic properties of metal hydrides.[196]

Considering the high excess energy at grain boundaries, a reduction of enthalpy greater than 30% should be possible assuming Mg particles smaller than 5 nm.[27e] However, this has not been observed experimentally because of recrystallization occurring upon hydrogen cycling. It has also been proposed that an excess in vibrational and configurational entropy at grain boundary could prevent grain growth,[197] but this has not been observed in hydride systems.

The structure inside the grain is often assumed to be the same as that of the initial micrometer-sized material, although it may contain defects such as dislocations, twin boundaries and stacking faults.[204] The structure of the grain boundaries has been the subject of many hypotheses in particular to support the enhanced hydrogen kinetics observed in the Mg/MgH₂ systems and alloys. For example, it is often assumed that the atomic structure in grain boundaries is more random without long range order,[205] and this would support models predicting faster hydrogen diffusion.[27m] However, it remains unclear if the structure of grain boundaries in nanocrystalline materials significantly differs from that of coarse-grained materials.[206]

### 5.2.2. Catalysts Dispersion for Improved Hydrogen Sorption Kinetics

Assuming the right hardness of the powders involved, mechanical milling is a very effective method for the refinement of powders and thus the dispersion of catalytic phases at the surface of hydride materials. This is particularly the case with MgH₂ which upon milling not only undergoes nanostructuration, but also can be “decorated” with various catalysts including transition metals (e.g., Ni, Ti, Fe, Nb, V),[52b,58–59,178b,207] their oxides,[202a,208] or intermetallics (e.g., FeTi, LaNi₅).[209] Typically a few at% of catalyst is sufficient. However, it should be noted that depending on the respective hardness of the powders to be milled and the conditions of milling, the catalyst may not only end up at the surface of the hydride but also be fully embedded within the hydride particles.[202c,207a,210] The embedded catalyst is believed to provide additional paths for hydrogen diffusion within the nanocrystalline structure.[202c,211] Other archetypical examples are complex hydrides and in particular NaAlH₄ for which ball milling has been revealed to be very effective for the dispersion of Ti-based[212] and other additives (e.g., Ce, Pr, Mn, Sc, Pr, Zr and Carbons)[213] that are believed to act as catalysts. However, in this case, not only the catalyst is finely “confined” within the hydride phase but it is also partially or fully reacted owing to the strong reducing character of complex hydrides.

<footnote>a) Converted to Mohs hardness value. [528]</footnote>
For example, upon milling TiCl₃ with NaAlH₄, a metathesis reaction occurs

$$ (1-x)\text{NaAlH}_4 + x\text{TiCl}_3 \rightarrow (1-4x)\text{NaAlH}_4 + 3x\text{NaCl} + x\text{Ti} + 3x\text{Al} + 6x\text{H}_2 \quad (41) $$

Two hypotheses have been suggested concerning the chemical state of Ti species when used as catalyst in alanates: i) Ti° is formed during ball milling and the active catalyst is a Ti/ Al alloy, and ii) Ti is substituted within the lattice of the bulk hydride.[177,214] Leon et al. identified the reduction of TiCl₃ to Ti° by ball milling NaAlH₄ with TiCl₃ and Ti was in a distorted state.[214,215] Balde et al. found that the majority (≈70%) of Ti occupied interstitial sites in NaAlH₄ and the rest of the Ti remained at the surface of Al.[216] TiAl₃ clusters were also observed by Weidenthaler et al. upon heating at 225 °C a mixture of Ti-doped NaAlH₄.[216,217] To date, the mechanism for the substitution of Na by Ti within the lattice of hydride remains unclear, but Weidenthaler et al. assumed that the replacement of Al atoms seems more plausible since the radius of the Ti⁺ cation is smaller than that of Na⁺.[217] The exact role of Ti in facilitating the release and uptake of hydrogen in NaAlH₄ is also controversial. First-principles calculations suggest that the dehydrogenation of Ti-doped NaAlH₄ is rate limited by the diffusion of AlH₄⁻ and AlH₃ vacancies and Ti has been suggested to facilitate the release of hydrogen from these diffusing species.[67a] Ti may also facilitate: i) the catalytic scission of the H₂ bond at the surface, ii) the catalytic breaking of H-AlH₃ bonds, and/or iii) the nucleation of the Al phase resulting from strong Ti-Al bonds and the formation of a coherent TiAl₃ phase.[67b,218]

### 5.2.3. Synthesis of Novel Hydrides

Mechanical milling has traditionally been used to prepare alloys and intermetallic compounds of the desired nanostructure and phase.[150,219] Early investigations also explored the formation of hydrides under low hydrogen pressures.[27i, 149a,220] The development of advanced milling tools including high pressure vials equipped with pressure and temperature sensors has facilitated the emergence of new approaches to mechanochemically synthesize hydride materials, including amides, alanates, borohydrides, and reactive composite hydrides.[189a]

Metal amides, such as KNH₂, LiNH₂, NaNH₂, Mg(NH₂)₂, and Ca(NH₂)₂ can be prepared by ball milling alkaline earth metal hydrides in gaseous ammonia to form the corresponding amides, following the general reaction path[189a,221]

$$ \mathrm{MH}_n + n\mathrm{NH}_3 \rightarrow \mathrm{M}(\mathrm{NH}_2)_n + n\mathrm{H}_2 \quad (42) $$

As a common practice, NH₃ is loaded at ≈0.4–0.5 MPa in the milling vials containing the desired metal hydride and the vials are refilled approximately every 2 h to ensure the progression of the reaction. An unavoidable problem with the practical application of metal amides for hydrogen storage is the simultaneous release of NH₃ and H₂.[213b,222] To limit this, metal amides are usually milled with the corresponding

metal hydride to form imides or nitrides,[213b,222,223] and this ultimately modifies the thermodynamic path for hydrogen release through H⁺ (amides) and H⁻ (hydrides) interreactions.[224]

The important family of alanates (e.g., LiAlH₄, NaAlH₄, Na₃AlH₆, KAlH₄, Mg(AlH₄)₂, CaAlH₅ and Ba₂AlH₇) can also be synthesized by ball milling metal hydrides with Al or AlH₃ under hydrogen pressure.[152] Li₃AlH₆ and Na₃AlH₆ have also been prepared by milling LiH/NaH with LiAlH₄/NaAlH₄ under hydrogen (Table 3). A few examples on the synthesis of bimetallic alanates have also been reported. These include Na₂LiAlH₆, NaLi₂AlH₆, K₂LiAlH₆, K₂NaAlH₆, and LiMg(AlH₄)₃.[225] For example, Na₂LiAlH₆ and NaLi₂AlH₆ were prepared by directly milling NaH, LiH, and Al or NaAlH₄ under hydrogen pressure (Table 3). Similarly, K₂LiAlH₆, K₂NaAlH₆ were synthesized from KH with LiAlH₄ and NaAlH₄, respectively (Table 3). In contrast, the synthesis of LiMg(AlH₄)₃ was only achieved by an indirect route involving the ball milling of Mg(AlH₄)₂ with LiCl.[225a,226]

The direct synthesis of borohydrides from their elements by ball milling is difficult to achieve owing to the high hydrogen pressure needed and the difficulty of breaking the B–B or B–metal bond in favor of the formation B–H. LiBH₄ and NaBH₄ have been successfully obtained by ball milling Li, B, or Na, B under 15 MPa hydrogen pressure (Table 3).[227] By using diborane (B₂H₆) as the starting reactant the synthetic route is easier and quantitative[228]

$$ \mathrm{MH}_n + n/2\mathrm{B}_2\mathrm{H}_6 \rightarrow \mathrm{M}(\mathrm{BH}_4)_n \quad (43) $$

where M is the metal and n is the valence.

This approach has been effective for the synthesis of LiBH₄, NaBH₄, Ca(BH₄)₂, Mg(BH₄)₂, and Y(BH₄)₃.[227b,228] However, it is not generally applicable, since the formation of the borohydride phase competes with the formation of the corresponding boride. For example, in the case of TiH₂ ball milled with B₂H₆, the final solid product is not Ti(BH₄)₂ but TiB₂, which is thermodynamically more stable.[227a] Bimetallic borohydrides have also been prepared by ball milling metal borohydrides with alkali metal chlorides or other borohydrides (Table 3). For example, Ravnsbæk et al. reported the synthesis of LiZn₂(BH₄)₅, NaZn₂(BH₄)₅, NaZn(BH₄)₃, by ball milling different ratios of LiBH₄, NaBH₄ with ZnCl₂, respectively. These new bimetallic borohydrides revealed remarkable structural diversity and low decomposition temperatures.[229] Similarly, the trimetallic borohydride Li₃MZn₅(BH₄)₁₅ (M = Mg, Mn) was synthesized by ball milling a LiBH₄, M(BH₄)₂ and ZnCl₂ mixture.[230]

A metathesis reaction approach is a common way to synthesize metal borohydrides and alanates by mechanochemical reaction. Usually an n-valent metal chloride or fluoride (NXₙ) is reacted with alkaline metal borohydrides (M(BH₄)) or alanates (M(AlH₄)) to produce a new complex metal borohydride or alanates.[231] The reaction can be simplified as

$$ \mathrm{NX}_n + n\mathrm{M}(\mathrm{BH}_4)/\mathrm{nM}(\mathrm{AlH}_4) \rightarrow \mathrm{N}(\mathrm{BH}_4)_n/\mathrm{N}(\mathrm{AlH}_4)_n + n\mathrm{MX} \quad (44) $$

The advantage of performing such a synthesis by ball milling is that the reaction is done at the solid state and thus does not involve solvents that may form intermediate complexes. However, a purification step involving some solvents may be
required at the end of the process to remove the halide and intermediate products.

Mixed hydrides, e.g., metal borohydrides/alanates/amides have also been synthesized mechanochemically. By tuning the ratio in the initial mixtures, different borohydride amides or alanate amides can be formed (Table 3). For example, Noritake et al. reported the synthesis of Li₂(BH₄)NH₂ and Li₄(BH₄) (NH₂)₃ by varying the ratio of LiNH₂:LiBH₄ from 1:1 to 3:1.[232] Chua et al. found that LiAlN₂H₂ and Li₂AlNH₂ were formed by the interaction of LiAlH₄ with 2NaNH₂ and 2LiAlH₄ with NaNH₂, respectively.[233]

Compositional alteration of ammonia borane (AB) by mechanochemical milling with metal hydrides to form metal amidoboranes has also been investigated by mechanochemical synthesis, with the aim of partially substituting its protic Hδ⁺ with a more electropositive metal[234]

$$M_1H_n + nNH_3BH_3 \rightarrow M_1(NH_2BH_3)_n + nH_2 \quad (45)$$

$$M_1H_n + M_2H_m + (m+n)NH_3BH_3 \rightarrow M_1M_2(NH_2BH_3)_{m+n} + 2(m+n)H_2 \quad (46)$$

where n and m are the valence states of M₁ and M₂, respectively. M₁ can be Li, Na, K, Ca, and Sr. In this case, M₁M₂ is LiNa, KAl, and Na₂Mg.[235]

Xiong et al. synthesized LiNH₂BH₃ and NaNH₂BH₃ by ball milling stoichiometric amounts of AB with LiH and NaH, respectively.[236] The synthesized LiNH₂BH₃ and NaNH₂BH₃ released ≈11 and ≈7.4 mass% H₂ at ≈91 °C, respectively. Unlike AB, which releases borazine with H₂ at the same time, there was no emission of borazine found from LiNH₂BH₃ and NaNH₂BH₃.[236] Na[Li(NH₂BH₃)₂] was found to decompose exothermally between 75 and 110 °C with ammonia as a byproduct.[237] A similar approach was also successfully applied to hydrazine borane N₂H₄BH₃, a derivative of AB,[238] to synthesize alkali hydrazinidoboranes of formulae MN₂H₃BH₃ (with M = Li, Na, K).[239] The latter hydrazinidoborane derivatives were found to have much better dehydrogenation properties in terms of H₂ purity, onset temperatures of dehydrogenation and safety. LiNH₂BH₃·NH₃BH₃ as an AB derivative has also been synthesized by ball milling under H₂ pressure. It showed better dehydrogenation kinetics over LiNH₂BH₃ and AB alone. The first-principles study showed that the enhanced dehydrogenation property was via the combination of hydridic Hδ⁺(B) from LiNH₂BH₃ and protonic Hδ⁺(N) from the adjacent AB.[240]

### 5.2.4. Formation of Other Novel Phases by Mechanical Milling

The intense plastic deformation induced in high energy milling can sometimes result in the introduction of disorder to the extent of forming an amorphous structure.[164] The mechanism for amorphization is similar to the solid state amorphization reaction proposed by Schwarz and Johnson.[241] It is therefore often observed that upon mechanical milling partial or complete amorphization of hydride materials can occur. For example, Wang et al. synthesized amorphous Mg–FeTi₁.₂ with

enhanced hydrogen absorption kinetics.[242] An amorphization can also be observed for complex hydrides. For example, ball milling MgB under 10 MPa hydrogen pressure led to amorphous Mg(BH₄)₂.[243] Its decomposition behavior was similar to crystalline Mg(BH₄)₂. The possibility of rapid energization and quenching during the milling process upon ball to powder impact can potentially also lead to the formation of metastable phases,[164a] i.e., phases which have a higher free energy than the equilibrium phase thermodynamically stable. For example, the high pressure metastable γ-MgH₂ phase is commonly obtained in a low amount (≈2–6%) by ball milling MgH₂.[46a,52b,174a] Xiao et al. reported a higher amount (≈18%) of γ-MgH₂ by milling MgH₂ with LiCl.[244] γ-MgH₂ is believed to lead to improved hydrogen diffusions and better thermodynamic properties and thus a lower temperature at which hydrogen can be released. However, the rapid restructuration of ball milled MgH₂ upon the first hydrogen cycling has never enabled the characterization of its hydrogen properties. Recently, we reported the electrochemical synthesis of a γ/β-MgH₂ mixture with ≈30% γ-MgH₂, which enabled for the first time the experimental determination of the hydrogen properties of γ-MgH₂.[245] γ-MgH₂ was found to display a reduced enthalpy of 57.7 ± 5.3 kJ mol⁻¹ H₂ instead of the 75 kJ mol⁻¹ H₂ of β-MgH₂; and thus was amenable to fully release hydrogen at 200 °C.

It should also be noted that mechanical milling has also shown some potential toward the formation of solid solutions[246] and the alloying of immiscible systems. For example, the solubility of Ti in Mg is very small (<0.1 at%). By mechanical alloying, under nonequilibrium conditions,[246a] this solubility can be extended to 12.5 at% due to a favorable size factor (9%) and an isomorphous structure.[247]

The dissolution of the alloying atom through mechanical milling can also modify the lattice parameters of the host materials. For example, Cd, In, Li, Al, Zn, Ag, Ga, Ti, and Sn have been found to reduce the Mg lattice parameters by forming a solid solution. In contrast, Sc and Pd were found to expand the lattice of Mg.[247,248] The formation of Mg-based solid solutions can enhance its hydrogenation and dehydrogenation. For example, the bcc solid solution (Nb, Mg)Hₓ (x < 2) was predicted to lead to lower hydrogen desorption temperatures (300 instead of 400 °C) as compared to pristine MgH₂.[249] In contrast, the formation of Mg(Li) solid solution did not lead to any improvement in the hydrogen properties of Mg because of the formation of the segregated and stable LiH phase upon hydrogenation.[248]

## 6. Metastable and Amorphous Phases: Storing Hydrogen through Thermodynamic Instability

A metastable state is a thermodynamically stable state with a local free energy minimum which does not correspond to the thermodynamic equilibrium state. Metastable phases can be crystalline and amorphous and are formed by various synthetic techniques including melt spinning,[250] atomization process,[251] ball milling,[44] and high pressure torsion.[252] For hydrogen storage materials, it is expected that metastable hydrides can display lower hydrogen desorption temperatures owing to their thermodynamic instability. For example,
nanocrystalline TiFe when ball milled with 4 wt% Ti was reported to lead to the formation of the nanocrystalline Fe₄₆Ti₅₄ metastable phase.[151b,253] This not only decreased Peq but also flattened the plateau.

Some weakly bound chemical hydrides, e.g., AlH₃, NH₃BH₃, and LiAlH₄, are also considered as metastable or kinetically stabilized hydrides.[254] These compounds have a Peq above 0.1 MPa at the ambient and thus only decompose very slowly at room temperature due to kinetic limitations. However, for these compounds hydrogen reversibility only occurs under harsh conditions of temperature and pressure, for example, 8.9 × 10³ MPa hydrogen pressure and 600 °C are required to recombine hydrogen with Al.[255] We found that after nano- sizing these materials, some hydrogen reversibility was pos- sible. Hence, for LiAlH₄ nanoconfined in a mesoporous carbon, hydrogen uptake and release was achieved at 300 °C and under 7 MPa hydrogen pressure.[256] Similarly, upon nano- sizing NH₃BH₃, hydrogen reversibility was achieved at 200 °C and under 6 MPa hydrogen pressure.[24c]

Amorphous alloys such as TiCu and Ti₂Cu may display larger hydrogen storage capacity compared to their crystalline counterparts.[257] For example, TiCu amorphous alloys were reported to absorb about 35% more hydrogen than crystal- line TiCu.[257] This is because in the crystalline structure, the number of interstitial sites within the lattice limit the intro- duction of hydrogen.[258] However, no evidence of a plateau pressure exists for amorphous hydrides owing to the large distribution of interstitial site energies (i.e., a large variation of sites with different coordination numbers and geometrical configurations) available for hydrogen.[27m] The lack of equi- librium plateau is a drawback for application, since hydrogen cannot be delivered at a constant pressure from the amorphous alloy. In some cases, the absorption of hydrogen can also result in an amorphization of the alloy, e.g., Zr–Rh systems, and hydrogen induced amorphization has been observed on many intermetallic compounds.[259] The empirical rule of Aoki et al. can be used to determine if an alloy can undergo hydrogen amorphization. For a Goldschmidt radii Rₐ/R<sub>B</sub> > 1.37 (with R corresponding to the ionic radius of A and B), the alloy AB should undergo hydrogen amorphization as a result of its low crystal lattice stability.[260] Most amorphous alloys are formed by rapid quenching (cooling rate of the melt ≈ 10⁴ °C s⁻¹). By mechanical milling, an amorphous phase can also be generated and in the case of nanocrystalline materials, it is believed that the structure at grain boundaries is amorphous – hence, the higher hydrogen solubility and the faster hydrogen kinetics,[151] assuming that the temperature of hydrogenation does not exceed the recrystallization temperature of the metal hydride.

## 7. Improving the Hydrogen Properties of Light Hydrides via Chemical Destabilization

Light metal hydrides such as LiH and MgH₂, and light metal complex hydrides such as amides (e.g., LiNH₂), alanates (e.g., NaAlH₄), and borohydrides (e.g., LiBH₄), have been at the center of many investigations owing to their high volumetric and gravimetric hydrogen storage capacities. The challenges remain their high thermodynamic barrier and sluggish

kinetics. The chemical bonding in light metal hydrides is pre- dominately covalent, ionic, or a mixture. This means that light hydrides are thermodynamically stable in comparison to tran- sition metal hydrides having a delocalized metallic bonding within their structures. Covalent, or ionic bonds also result in large activation barriers for atomic motion, and this results in slow hydrogen sorption kinetics which further restricts the hydrogen reversibility. Following the early examples of thermo- dynamic destabilization of metal hydrides (Figure 14b),[119,136] similar ideas have emerged to improve the thermodynamics of complex hydrides but this often assumes that the entropy of reaction for these hydrides remains constant which is not the case.[261]

### 7.1. The Approach of Reactive Hydride Composites (RHC)

Vajo et al.[112,261] were among the first to demonstrate the con- cept of RHC based on the idea that the reactions between two or more hydrides (MHₙ and M₁Hₘ) can lead to the formation of an unstable hydride phase and/or a more stable product (MM₁) during dehydrogenation and thus facilitate the release of hydrogen (Figure 14b)

$$MH_n + M_1H_m \leftrightarrow MM_1 + \frac{n+m}{2}H_2 \quad (47)$$

The enthalpy of formation of composite hydrides is usually exothermic, and this leads to an overall reduction of the dehy- drogenation enthalpy for the composite hydride, while a high hydrogen storage capacity is maintained. The dehydrogena- tion enthalpy (ΔH°deh) in this case is equal to the difference in the standard enthalpy of formation between the products and starting hydrides

$$\Delta H^{\circ}_{deh} = \frac{2}{n+m} (\Delta_{f}H^{\circ}_{MM_{1}} - \Delta_{f}H^{\circ}_{MH_{n}} - \Delta_{f}H^{\circ}_{M_{1}H_{m}}) \quad (48)$$

Similarly, the entropy of dehydrogenation (ΔS°deh) is given by the difference between standard entropy of the products and starting hydrides

$$\Delta S^{\circ}_{deh} = \Delta S^{\circ}_{H_{2}} + \frac{2}{n+m} (\Delta S^{\circ}_{MM_{1}} - \Delta S^{\circ}_{MH_{n}} - \Delta S^{\circ}_{M_{1}H_{m}}) \quad (49)$$

The enthalpy of formation and entropy of the decomposition reaction of some common hydrides that can be used for synthe- sizing RHC systems are summarized in Table 2.

The LiBH₄ and MgH₂ mixture is an archetypical example of RHC that has been widely investigated owing to the high hydrogen capacity of both hydrides, i.e., 18.5 and 7.6 mass% H₂, respectively. The composite LiBH₄/MgH₂ is effectively less stable than the individual hydrides (Figure 18a) and follows the decomposition paths

$$\mathrm{LiBH}_{4}+1/2\mathrm{MgH}_{2} \leftrightarrow \mathrm{LiH}+1/2\mathrm{MgB}_{2}+2\mathrm{H}_{2} \quad (50)$$

$$\mathrm{LiH}+1/2\mathrm{MgB}_{2} \leftrightarrow \mathrm{Li}+1/2\mathrm{MgB}_{2}+1/2\mathrm{H}_{2} \quad (51)$$
<figure><img src="image_21.png" /><figcaption>**Figure 18.** Thermodynamics of RHC systems with LiBH₄/MgH₂ as an example. a) Evolution of the ΔH and ΔS as function of the reaction path; and b) phase boundary diagram of LiBH₄ + 1/2MgH₂ determined from theoretical calculations based on the enthalpy and entropy obtained by Vajo et al.[112] ΔH and ΔS were determined by using the NIST Chemistry WebBook and NIST-JANAF thermochemical tables. The calculated values of ΔS reported here do not correspond to the experimental values reported in equations (30) and (31).</figcaption></figure>

For Figure 18a, it can be seen that the desorption enthalpy of (50), 33 kJ mol⁻¹ H₂, is reduced as compared to that of pristine LiBH₄ (49 kJ mol⁻¹ H₂). In contrast, the entropy of (50), 112 J K⁻¹ mol⁻¹ H₂, was found to be higher than that of pristine LiBH₄ (97 J K⁻¹ mol⁻¹ H₂) but lower than the entropy for MgH₂ (131 J K⁻¹mol⁻¹ H₂). Achieving hydrogen reversibly for borohydrides was also found to be easier under this approach. Indeed, direct rehydrogenation of mixtures of MgB₂ and the binary hydrides (LiH, NaH, or CaH) led to the formation of LiBH₄, NaBH₄, and Ca(BH₄)₂.[262] The kinetic barrier for the hydrogenation reaction starting from MgB₂ is believed to be much lower than that starting from elemental boron.[262b] Additionally, the reactivity of boron in MgB₂ is much higher than elemental boron, possibly due to: a) the different strength of boron bonding (B-B = 297 ± 21 kJ mol⁻¹, B-H = 340 kJ mol⁻¹),[263] b) the different bond lengths (1.77-1.83 Å for icosahedra of elemental boron,[264] and 1.78 Å for B-B and 2.5 Å for Mg-B of MgB₂[265]) and c) the different coordination numbers in B-B, B-H, Mg-B-H.

However, as compared to the alloying of metals, reactive hydride mixtures involving complex hydrides are much more difficult to control owing to the possibility of multiple decomposition paths. These further depend upon the reaction conditions and the ratio between the two hydrides,[243] and such a complexity can lead to very different destabilization effects. For example, it was found that the 2LiBH₄/MgH₂ mixture led to production of MgB₂ (reaction (50)) under hydrogen atmosphere according to the boundary diagram (Figure 18b), and this facilitates the destabilization of LiBH₄ while stabilizing the dehydrogenated state. Through this mixture, the enthalpy of the reaction was reduced by 25 kJ mol⁻¹ H₂, as compared to pure LiBH₄,[112] and the dehydrogenation temperature was reduced to 225 °C at a hydrogen equilibrium pressure of 0.1 MPa. This is significantly lower than the decomposition temperature of the single hydrides (Table 2).[112] In contrast, the same 2LiBH₄/MgH₂ mixture decomposed under vacuum led to formation of elemental Mg and B following (50)

$$ \mathrm{LiBH_4} + \mathrm{1/2MgH_2} \leftrightarrow \mathrm{Li} + \mathrm{1/2Mg} + \mathrm{B} + \mathrm{5/2H_2} \quad (52) $$

Additionally, enriching MgH₂ (up to a mass ratio of 1:4) in the LiBH₄/MgH₂ system was found to result in the forma-

tion of Li-Mg alloys, e.g., the Mg₀.816Li₀.184 and the Mg₀.70Li₀.30 alloys, following the reaction paths (53) and (54)[266]

$$ 0.3\mathrm{LiBH_4} + \mathrm{MgH_2} \rightarrow 0.3\mathrm{LiBH_4} + \mathrm{Mg} + \mathrm{H_2} \quad (53) $$

$$ 0.3\mathrm{LiBH_4} + \mathrm{Mg} + \mathrm{H_2} \rightarrow 0.78\mathrm{Mg}_{0.816}\mathrm{Li}_{0.184} + 0.52\mathrm{Mg}_{0.70}\mathrm{Li}_{0.30} + 0.3\mathrm{B} + 1.60\mathrm{H_2} \quad (54) $$

This also resulted in partial hydrogen reversibility of the system under 10 MPa hydrogen pressure and at 400 °C.[266] Hence, the formation of MgB₂ is not necessary for enabling hydrogen reversibility.

The most favorable conditions for hydrogen reversibility can be determined from the phase boundary diagram of the LiBH₄ + 1/2MgH₂ system (Figure 18b). Under sufficient hydrogen back pressure, direct decomposition of LiBH₄ leading to formation of elemental boron can be avoided, and thus it is possible to alter the reaction path by adjusting the hydrogen pressure upon the decomposition of the hydride mixture.[267]

Although the concept of RHC can lead to some enhancement in the dehydrogenation reaction, this does not address the kinetic issues which significantly limit hydrogen reversibility under moderate conditions owing to slow ionic diffusion even at high temperatures.[268] Hydrogen absorption in a metal involves multiple steps: physisorption, dissociation, chemisorption, diffusion, phase transition, hydride formation of nucleation and growth (Figure 8a). The reverse processes occur during the dehydrogenation reaction and the rate of hydrogen sorption is governed by the slowest reaction step. Understanding hydrogen or ionic diffusion in complex hydrides involves very different mechanisms that may evolve during the hydrogen uptake/release because of the changes in crystalline and amorphous structures occurring in reactive hydride mixtures. In this case, fast hydrogen desorption kinetics also requires both the mobility of hydrogen carrying species (e.g., the BH₄⁻ unit)[269] and the metal counter ions. The finite hydrogen flux (jH,i) in the material is thus limited by the gradient in chemical potential (∇μH,i) for the ions (i) across the different hydride phases as follows[270]

$$ j_{H,i} = -L_i \nabla \mu_{H,i} \quad (55) $$
where $L_i$ is the parameter related to the transport mechanism within the material, such as the concentration of the hydrogen gas and temperature.[271]

For borohydrides, the chemical potential ($μ_{H,i}$) is proportional to $ln(P)$. Thus the diffusion of hydrogen is linearly correlated to $ln(P/P^{eq})$ at the reaction interface.[272] In contrast, hydrogen desorption from $MgH_2$ varies linearly with $(P - P^{eq})^a$, with $0.5 < a < 1$, assuming that hydrogen diffusion through the $MgH_2$ particles is the main limiting step.[178b] Hence, owing to the variations in hydrogen desorption regimes, different decomposition steps corresponding to the decomposition of $MgH_2$ first and then of $LiBH_4$ can lead to a single or double plateau pressure.

Experimentally, the $E_a$ for the dehydrogenation of the $LiBH_4/MgH_2$ system determined by the Kissinger method was found to be of $221 \pm 5$ and $306 \pm 10$ kJ mol$^{-1}$ for the successive dehydrogenation of $MgH_2$ and $LiBH_4$, respectively.[273] These values are higher than the $E_a$ reported for pristine $MgH_2$ ($156$ kJ mol$^{-1}$)[270] and $LiBH_4$ ($59 \pm 2$ kJ mol$^{-1}$),[274] alone. Accordingly, kinetic enhancement through the RHC system approach can be limited. A similar observation was made for the $NaBH_4/MgH_2$ composite. In this case, $E_a$ was found to be of $100 \pm 4$ and $133 \pm 13$ kJ mol$^{-1}$, for the decomposition of $MgH_2$ and pristine $NaBD_4$, respectively.[270]

The kinetics of RHC systems is controlled by various activation barriers and also the microstructure of the system and the particles/particles interactions as illustrated in Figure 10.[275] Modeling the reactions kinetics through the Johnson-Mehl-Avrami (JMA) approach revealed that the absorption and desorption reactions both follow a 3D growth of contracting volume with a constant and decreasing interface velocity.[55b-d] This means that during the de/rehydrogenation process, the growth of new phases starts from the surface and propagates inward toward the particles' center. Therefore under this assumption, the nucleation of $MgB_2$ can be a major impediment for the progress of the reactions in the case of the $LiBH_4/MgH_2$ system, for example.[268] Indeed, experimentally an incubation period is often observed after the decomposition of $MgH_2$ into metallic Mg, before the start of the second step, i.e., decomposition of $LiBH_4$.[270] Therefore, parameters including phase distribution, particles/particles interaction, grain boundaries, and crystallite size should be properly controlled to maximize the mobility of atoms/ions and promote the reactivity of the complex hydrides. For hydrogen absorption, the mobility of Mg, B and Li, as well as hydrogen diffusion upon the decomposition of $MgB_2$ and LiH govern the regeneration of the initial borohydride/hydride mixture. It has been claimed that hydrogen reaction at the $MgB_2$ interface is the rate-limiting step during absorption, rather than the diffusion of hydrogen and ions participating in the regeneration of $LiBH_4$ and $MgH_2$.[268] This again suggests that the microstructure of $MgB_2$ is important for the rehydrogenation process.

Addition of so called “catalysts” including transition metal halides, is a popular approach to further enhance the desorption kinetics of RHC, leading in some cases to improved hydrogen reversibility.[275a] For instance, it has been claimed that addition of Ti additives can act as nucleation seeds and provide effective growth sites for the $MgB_2$.[276] The kinetics of RHC can

be further improved via nanoconfinement (Figure 6c). For example, $2LiBH_4-MgH_2$ confined within a resorcinol formaldehyde carbon aerogel scaffold showed a reduced $E_a$ of $89 \pm 2$ and $279 \pm 5$ kJ mol$^{-1}$ for the $MgH_2$ and $LiBH_4$ dehydrogenations, respectively.[273] This is attributed to the reduced diffusion distance for hydrogen.

## 7.2. Metathesis Reaction and Partial Substitution to Trigger Low Temperature Desorption

Reducing the energetic barriers to allow the decomposition of hydrides under milder conditions involves weakening of M-H bonds and intrahydrogen pairing (H-M-H). An alternative method to enhance the dehydrogenation and reversibility of complex hydrides is through the addition of a second metal to partially substitute for the central metal cation (M$^+$) in the complex hydride.[277]

$$ (1+m/n)M(BH_4)_n + M_1X_m (X = Cl, Br, I, etc.) \rightarrow M_1M(BH_4)_{n+m} + m/nMX_n \quad (56) $$

This can be achieved by taking advantage of significant electronegativity differences between hydrides forming elements and thus the partial substitution of M by a more electronegative element $M_1$ (Figure 14b). Archetypical examples of such an approach are based on borohydrides for which a correlation between the metal electronegativity and the hydrogen decomposition temperature has been established (Figure 19).[252b] In this case, the modification of the hydride and the partial substitution of M are usually performed via mechanical milling with the desired metal halide (see section 5.2.3). Examples of this include bimetallic borohydrides incorporating Sc and Zn within the Group I metal borohydrides (Figure 19a), because $Sc(BH_4)_3$ and $Zn(BH_4)_2$ are thermodynamically unstable.[278] By ball milling of $LiBH_4$ and $ScCl_3$, $LiSc(BH_4)_4$ with a hydrogen storage capacity of 14.5 mass% was easily prepared.[279] This new compound showed low temperature hydrogen release; however, no reversibility was achieved at 250 °C under 12 MPa hydrogen pressure.[279d] The release of hydrogen from $LiSc(BH_4)_4$ was found to follow multiple decomposition paths (57)–(60) and this does not facilitate hydrogen reversibility[279a]

$$ LiSc(BH_4)_4 \rightarrow 2/5LiBH_4 + 3/10Li_2B_{12}H_{12} + ScH_2 + 22/5H_2 - 128 \, ^{\circ}\mathrm{C} \quad (57) $$

$$ \rightarrow LiBH_4 + 1/10Li_2B_{12}H_{12} + ScB_2 + 29/5H_2 - 147 \, ^{\circ}\mathrm{C} \quad (58) $$

$$ \rightarrow Li_2B_{12}H_{12} + 2/3LiH + ScB_2 + 20/3H_2 - 170 \, ^{\circ}\mathrm{C} \quad (59) $$

$$ \rightarrow LiH + 4/5ScB_2 + 1/5ScB_{12} + 15H_2 \rightarrow 553 \, ^{\circ}\mathrm{C} \quad (60) $$

In addition the reaction leads to the release of gaseous $B_2H_6$, and the formation of the thermodynamically more stable intermediates $Li_2B_{12}H_{12}$ and LiH.

Mixed-cation borohydrides are of great interest due to: i) their high hydrogen content, ii) the facility to tune their hydrogen
<figure><img src="image_22.png" /><figcaption>**Figure 19.** Adjusting the thermodynamics of complex hydrides. Experimentally observed decomposition temperatures for selected a) metal borohydrides and bimetallic metal borohydrides; and b) metal borohydrides and amine borohydrides plotted as a function of the electronegativity, χp, of the metal cation. The blue dashed line represents the linear relationship between the decomposition temperatures and electronegativity of metal cation borohydrides.</figcaption></figure>

sorption properties by adjusting the molar ratios between the two reactants, and iii) the variety of crystalline structures that can be obtained.[229] However, this approach has several drawbacks: a) most dual-borohydride cations still release hydrogen above 100 °C, b) the synthesis of mixed-cation borohydrides often leads to undesired halide byproducts, which reduces the available hydrogen content, and c) their instability results in the formation of unwanted byproducts (e.g., B₂H₆) and thermodynamically stable products, e.g., dodecaborate ([B₁₂H₁₂]²⁻) and transition metal borides during thermal decomposition.[279d] This limits the conditions for rehydrogenation.[278] Nanoconfinement can provide an approach to overcome some of these problems owing to the shorter diffusion paths and the likelihood of minimizing the number of side reactions. For instance, nanoconfined eutectic composites of borohydrides have shown great improvements on the hydrogen release temperature,[280] with the possibility of modification of the decomposition

path.[281] This is often reflected by the suppression of the formation of volatile gases, i.e., B₂H₆, upon nanoconfinement, which may further improve the reversibility of the materials.

### 7.3. Destabilization of Hydrides through H⁺/H⁻ Interaction

Metal hydrides can be classified into three categories, based on the bonding between the metal and hydrogen: i) ionic, e.g., LiH where the hydrogen existed as a H⁻ ion, ii) covalent, and iii) metallic hydrides (Figure 4). In complex hydrides, such as LiBH₄, the hydrogen exists as a partially negatively charged Hδ⁻. In contrast, hydrogen exist as partially positively charged Hδ⁺ in covalent hydrides, such as LiNH₂.[282] A strategy to facilitate the release of hydrogen from hydrides is thus through the combination of ionic and covalent hydrides, which is equivalent to a Lewis acid/base reaction with Hδ⁻ acting as an electron donor, while Hδ⁺ acts as an electron acceptor with[282]

$$ \delta H^{-} + H^{\delta+} \rightarrow (1 + \delta)H_{2} \qquad (61) $$

$$ H^{-} + H^{-} \rightarrow H_{2} + 2e^{-} \qquad (62) $$

$$ H^{\delta+} + H^{\delta+} + 2\delta e^{-} \rightarrow H_{2} \qquad (63) $$

Reaction (61) is easier as compared to the reaction paths (62) and (63), which require high temperatures due to the repulsive potential between the two positively or negatively charged hydrogen ions and the additional need of electron transfer.

One of the first demonstrations of the H⁺/H⁻ concept was with the LiNH₂/LiH system.[283] By combining Hδ⁺ in LiNH₂ and H⁻ in LiH a dehydrogenation temperature as low as 150 °C was achieved along the decomposition paths[284]

$$ \mathrm{LiNH}_{2}+\mathrm{LiH} \rightarrow \mathrm{Li}_{2} \mathrm{NH}+\mathrm{H}_{2} \quad 150^{\circ} \mathrm{C} \quad 6.5 \mathrm{~mass}\% \mathrm{H}_{2} \qquad (64) $$

$$ \mathrm{LiNH}_{2}+2 \mathrm{LiH} \rightarrow \mathrm{Li}_{3} \mathrm{N}+2 \mathrm{H}_{2} \quad 320^{\circ} \mathrm{C} \quad 9.6 \mathrm{~mass}\% \mathrm{H}_{2} \qquad (65) $$

Further investigation by Hu and Ruckenstein[285] and Ichikawa et al.[221c] claimed the formation of intermediates (i.e., NH₃). Therefore, the destabilization effect could be from the interaction between Hδ⁺ in ammonia and H⁻ in LiH according to

$$ 2\mathrm{LiNH}_{2} \rightarrow \mathrm{Li}_{2} \mathrm{NH}+\mathrm{NH}_{3} \quad 374^{\circ} \mathrm{C} \quad 9.6 \mathrm{~mass}\% \mathrm{H}_{2} \qquad (66) $$

$$ \mathrm{LiH}+\mathrm{NH}_{3} \rightarrow \mathrm{LiNH}_{2}+\mathrm{H}_{2} \quad 100^{\circ} \mathrm{C} \quad 8.1 \mathrm{~mass}\% \mathrm{H}_{2} \qquad (67) $$

These initial results further inspired the exploration of mixing metal amides/imides with hydrides to achieve high capacity hydrogen storage materials. For example, the decomposition of 2LiNH₂/MgH₂ resulted in the formation of the Li₂Mg(NH)₂ intermediate, with a 4.2 mass% H₂ reversibility at 200 °C and 3 MPa hydrogen pressure.[286] It is noteworthy that Mg(NH₂)₂ and LiH mixtures with various ratios of amide to hydride can lead to formation of different products as following[287]
$$ \mathrm{Mg}(\mathrm{NH}_2)_2 + n\mathrm{LiH} \leftrightarrow \mathrm{Li}_2\mathrm{Mg}(\mathrm{NH}_2) + (n-2)\mathrm{LiH} + 2\mathrm{H}_2 \quad 140-250\,^{\circ}\mathrm{C} \quad 5.6 \text{ mass}\% \mathrm{H}_2 \quad (68) $$

$$ \mathrm{Mg}(\mathrm{NH}_2)_2 + 8/3\mathrm{LiH} \leftrightarrow 1/3\mathrm{Mg}_3\mathrm{N}_2 + 4/3\mathrm{Li}_2\mathrm{NH} + 8/3\mathrm{H}_2 \quad 140-280\,^{\circ}\mathrm{C} \quad 7 \text{ mass}\% \mathrm{H}_2 \quad (69) $$

$$ \mathrm{Mg}(\mathrm{NH}_2)_2 + 4\mathrm{LiH} \leftrightarrow 1/3\mathrm{Mg}_3\mathrm{N}_2 + 4/3\mathrm{Li}_3\mathrm{N} + 4\mathrm{H}_2 \quad 140-520\,^{\circ}\mathrm{C} \quad 9.1 \text{ mass}\% \mathrm{H}_2 \quad (70) $$

This concept of H⁺/H⁻ interaction has also been expanded to a large range of hydrides with the use of ball milling as the main technique to favor the intimate mixing between both hydrides. For example, different ratios between LiBH₄ and LiNH₂ led to quaternary hydride Li₃BN₂H₈,[288] and Li₄BN₃H₁₀.[289] Li₃BN₂H₈ has a theoretical hydrogen capacity of 11.9 mass% and this compound released 10 mass% H₂ with only a small amount of ammonia (2–3 mol%) above 250 °C. In this case, the destabilization effect was further evidenced by first-principles calculations, where the dehydrogenation enthalpy of the LiBH₄/2LiNH₂ mixture was reduced from 75 kJ mol⁻¹ H₂ for pure LiBH₄ to 23 kJ mol⁻¹ H₂. The same concept was also investigated for NaBH₄,[290] Mg(BH₄)₂.[291] and Ca(BH₄)₂,[292] and similar reaction mechanisms were observed. More recent investigations used CO(NH₂)₂—an organic carbonyl diamide—to modify LiBH₄ and produce LiBH₄·CO(NH₂)₂. This new system was able to release 8.13 mass% H₂ exothermically at temperatures below 250 °C.[293]

Mixing LiAlH₄ and LiNH₂ can also lead to the formation of Li–Al–N intermediates to release hydrogen at reduced temperatures. But once again, the reaction paths significantly depend upon the LiNH₂ to LiAlH₄ ratios[294]

$$ 2\mathrm{LiNH}_2 + \mathrm{LiAlH}_4 \leftrightarrow \mathrm{Li}_3\mathrm{AlN}_2 + 4\mathrm{H}_2 \quad 25-600\,^{\circ}\mathrm{C} \quad 9.6 \text{ mass}\% \mathrm{H}_2 \quad (71) $$

$$ 2\mathrm{LiNH}_2 + 2\mathrm{LiAlH}_4 \leftrightarrow 2\mathrm{Li}_2\mathrm{AlNH} + 5\mathrm{H}_2 \quad 50-700\,^{\circ}\mathrm{C} \quad 8.2 \text{ mass}\% \mathrm{H}_2 \quad (72) $$

$$ \mathrm{LiNH}_2 + 2\mathrm{LiAlH}_4 \leftrightarrow 2\mathrm{Al} + \mathrm{Li}_2\mathrm{NH} + \mathrm{LiH} + 4\mathrm{H}_2 \quad 85-320\,^{\circ}\mathrm{C} \quad 8.1 \text{ mass}\% \mathrm{H}_2 \quad (73) $$

To date, the problem with such an approach remains the low hydrogen reversibility. Only 3.5–4.0 mass% H₂ is usually reabsorbed in LiNH₂/LiAlH₄ mixtures at temperatures ≈380 °C and under 20 MPa hydrogen pressure. An additional drawback is the formation of the more stable Li₃AlH₆ as an intermediate, and because the rehydrogenation of Li₃AlH₆ to LiAlH₄ requires extreme conditions this further impedes hydrogen reversibility.[295] The decomposition of these systems toward LiH is thus a more thermodynamically favorable path.[296] The alternative is to directly mix LiNH₂ with Li₃AlH₆. In this case, a higher reversible hydrogen content of 7 mass% was observed at moderate conditions (300 °C, 10 MPa).[297]

An additional example of protic–hydridic-based hydrogen storage materials is ammonia borane (NH₃BH₃, AB). AB has a very high hydrogen content of 19.6 mass%. The existence of dihydrogen bonding in AB between protic (N–H) and hydridic

(B–H) hydrogens is believed to lead not only to enormously high volumetric and gravimetric density but also to aid the low temperature hydrogen release, e.g., 6.5 mass% of H₂ at 100 °C.[298] However, the use of AB as a hydrogen storage material is not without major drawbacks. The thermal decomposition of AB not only leads to the formation of various polymers, e.g., polyaminoborane ((NH₂BH₂)ₙ) and polyiminoborane ((NHBH)ₙ), but also volatile gases including ammonia (NH₃), diborane (B₂H₆), and borazine (B₃N₃H₆). Moreover, the full decomposition of AB only occurs at a temperature above 1000 °C, with formation of the thermodynamically stable boron nitride. This means that achieving full reversibility of AB is out of reach. Recent investigations focused on modifying the properties of AB via its confinement within a wide range of nanoscaffolds such as SBA-15,[299] carbons,[300] MOFs,[301] boron nitride,[302] and polypyrrole nanotubes.[303] We recently showed the possibility of achieving reversibility with a nanosized AB/Ni matrix at 200 °C and 6 MPa.[24c] This demonstrates that some paths may enable the use of AB derived compounds as reversible hydrogen storage materials.

Following the example of AB, attempts have also been made to synthesize novel compounds incorporating both hydridic and protic hydrogens, through the reaction of NH₃ with borohydrides, for example. The possibility to synthesize amine borohydride was reported many decades ago.[304] As illustrated by Figure 19b, this strategy leads to a destabilization of borohydrides involving metals of high electronegativity. In contrast, borohydrides with metallic cations of low electronegativity were found to be more stable. Borohydrides combined with NH₃ have shown hydrogen capacities reaching 17.8 mass%, with in some cases the protic and hydridic hydrogen enabling hydrogen release at near room temperature. For example, 8.9 mass% of hydrogen was released from Zn(BH₄)₂·2NH₃ and 16.8 mass% of hydrogen was released from Al(BH₄)₃·5NH₃.[305] However, some studies suggested that the thermal decomposition of amine borohydrides causes a loss of ammonia during the hydrogen release and this contributes to the irreversibility of these systems. For example, LiBH₄·NH₃[306] and Ca(BH₄)₂·nNH₃ (n = 1, 2, and 4)[306d,307] mainly release ammonia rather than hydrogen below 300 °C under dynamic vacuum. Furthermore, combinations of different amine borohydrides with borohydrides, such as Al(BH₄)₃·4NH₃/2LiBH₄,[305] LiMg(BH₄)₃(NH₃)/LiBH₄,[308] Ca(BH₄)₂·2NH₃/LiBH₄,[306d] and Ca(BH₄)₂·4NH₃/Mg(BH₄)₂[307a] can significantly increase the available amount of hydrogen while improving the issue of hydrogen purity.

## 8. The Nanosizing Approach

To enhance the effects of the basic concepts discussed above, intensive experimental work has been carried out in recent years on the hydrogen sorption properties of various hydride materials upon nanosizing, i.e., a reduction of the hydride particle size to a few nm (Figure 6c).[309] Upon nanosizing, improved thermodynamics, e.g., ΔH decreasing from 74.5 to 44.5–66 kJ mol⁻¹ H₂ was reported for Mg/MgH₂,[309a-d,310] faster hydrogen kinetics, e.g., Eₐ decreasing from 116 to 58 kJ mol⁻¹ H₂ for NaAlH₄ nanoparticles of 2–10 nm,[183,309q,311] and reduced hydrogen desorption temperatures[309f,312] were observed. For
<figure><img src="image_23.png" /><figcaption>**Figure 20.** Thermodynamics of metal hydrides at the nanoscale a) Schematic correlation between ΔH and ΔS of nanosized hydrides and the role of ΔG variation (Δ(ΔG)). The correlation between ΔH and ΔS depends upon the variations of ΔG as compared to the variation (Δ(ΔH)). If (Δ(ΔG)) < (Δ(ΔH)) a compensation effect is observed and an anticompensation effect is reached for the reverse, i.e., for (Δ(ΔG))' > (Δ(ΔH)); b) theoretically calculated evolution of ΔH for various metal/metal hydrides systems at the nanoscale;[322] c) theoretically evaluated desorption enthalpy for MgH₂ clusters with both HF and DFT methods as a function of the numbers of Mg atoms.[323c]</figcaption></figure>

example, LiH confined high surface area graphite released hydrogen at 350 °C instead of 680 °C.[309f] Mg nanoparticles with the size of 4 nm incorporated in CMK3 showed a remarkable decrease in hydrogen desorption temperature of 253 °C.[309b] Milder temperatures and pressures for achieving hydrogen reversibility[313] have also been reported for some complex hydrides and this was attributed to an alteration of reaction pathways upon nanosizing.[309i,314] For example, carbon nanotubes confined LiBH₄ showed a single decomposition at 140 °C, much lower than the temperatures required for bulk LiBH₄ (>400 °C) and some hydrogen reversibility.[313a,314b] Similarly, the hydrogen desorption temperature was found to be much lower for nanoconfined NaBH₄ (150 °C),[313a] LiAlH₄ (120 °C),[315] NaAlH₄ (≈180 °C)[309s] compared to their bulk counterparts.

## 8.1. Nanosizing for Better Thermodynamics

It is widely accepted that particles at the nanoscale display different physical/chemical properties as compared to their bulk counterparts as a result of their higher surface energy.[186,316] In this case, when considering a reaction with nanoscale materials, the Gibbs free energy (ΔG<sub>nano</sub>) can be expressed as

$$ \Delta G_{\text{nano}} = \Delta G_{\text{bulk}} + \Delta G_{\text{surf}} \quad (74) $$

where ΔG<sub>bulk</sub> is the reaction Gibbs free energy of the bulk material (i.e., the core of the nanoparticle) and ΔG<sub>surf</sub> corresponds

to the excess of surface energy owing to the nanodimension of the material and the excess energy arising from unsatisfied surface atoms bonds.[27e,316] When considering a hydrogenation reaction, the basic thermodynamic laws described above (section 3.2) still apply and

$$ \Delta G_{\text{nano}} = RT \ln \frac{P_{\text{nano}}^{\text{eq}}}{P} \quad (75) $$

where P<sub>nano</sub><sup>eq</sup> is the equilibrium pressure of the nanosized hydride.

Hence, by associating (74) and (75),

$$ \ln \frac{P_{\text{nano}}^{\text{eq}}}{P_{\text{bulk}}^{\text{eq}}} = \frac{\Delta G_{\text{surf}}}{RT} \quad (76) $$

Thus at a given temperature and assuming sufficient excess of surface energy, the equilibrium pressure for the hydriding reaction of nanoparticles will be higher than that for the bulk material. Hence in principle, the thermodynamics of the hydrogen absorption and desorption reactions could be tailored by adjusting the surface energy of nanosized hydrides (Figures 14a and 20). Further considering Equations (21) and (74), the van't Hoff equation for the hydrogen storage reactions at nanoscale takes the form

$$ \ln \frac{P_{\text{nano}}^{\text{eq}}}{P^{\circ}} = \frac{\Delta H^{\circ}}{RT} - \frac{\Delta S^{\circ}}{R} + \frac{\Delta G_{\text{surf}}}{RT} \quad (77) $$
In this equation, the surface Gibbs free energy ($ΔG_{surf}$) is a function of temperature[317] as observed experimentally[318]

$$ \Delta G_{surf} = \Delta H_{surf} - T\Delta S_{surf} \quad (78) $$

Combining (77) and (78) leads to a new van't Hoff equation for nanoscale materials

$$ \ln \frac{P_{\text{nano}}^{\text{eq}}}{P^{\circ}} = \frac{\Delta H^{\circ} + \Delta H_{\text{surf}}}{RT} - \frac{\Delta S^{\circ} + \Delta S_{\text{surf}}}{R} \quad (79) $$

Hence through nanosizing, additional factors could affect both the $ΔH$ and $ΔS$ of the hydriding reaction. Indeed, for individual particles both $ΔH$ and $ΔS$ have been observed to vary at the nanoscale.[319]

In the context of hydride materials, it is difficult to predict a general trend for $ΔH$ and $ΔS$ as function of particle size, since this is highly subjective to the particle size and the nature of the reactants and products. An additional factor to consider is the long-standing empirically proposed “enthalpy-entropy compensation effect” which has been observed for some hydrides;[320] although it has been proposed that it may be a phantom phenomenon.[27d,321]

### 8.1.1. Thermodynamics of Nanosized Metal Hydride Systems

By taking into consideration the above, the thermodynamics of nanosized hydrides could potentially be controlled by an additional surface free energy term expressed as[27e,316]

$$ \Delta G_{\text{surf}} = \frac{3V_M}{r} \left[ \gamma_{\text{MH}_2}(r) \left( \frac{V_{\text{MH}_2}}{V_M} \right)^{\frac{2}{3}} - \gamma_M(r) \right] + E_{\text{ads}} \quad (80) $$

where $V_M$ is the molar volume, $r$ is the particle radius and $γ$ the surface free energy per surface unit. The additional energy term ($E_{ads}$) is the result of the difference in surface energy between the nonhydried and hydried nanoparticle upon hydrogen binding.[27e,316]

Therefore, $ΔG_{surf}$ depends upon the particle size, its nanostructure, and the surface properties of the nanosized hydride as well as $E_{ads}$. The latter is hard to predict. If $ΔG_{surf}$ displays a positive value, the system will be destabilized. Otherwise, the hydride phase will be stabilized.

Theoretical calculations regarding the evolution of $ΔH$ have been carried out for various hydrides using the Wulff construction.[322] As shown in Figure 20b, the enthalpy variation introduced by nanosizing effects differs significantly for different M/MH systems and depending on the M/MH system a destabilization or stabilization may occur. However, according to these calculations, these effects only occur at extremely small particle sizes. Similar conclusions have been suggested by additional theoretical calculation on Mg/MgH$_2$ (Figure 20c).[323]

Although theoretical calculations suggest that the particle size needs to be very small (<5 nm) for any observable thermodynamic alteration (Figure 20b,c),[322] or even smaller for Mg/MgH$_2$ (<1 nm),[323a,c,d] reduction of enthalpy has been commonly reported experimentally for relatively large particles

(10–50 nm).[182a,309a,d,g,324] The divergence between experimental observations and theoretical calculations is possibly explained by the fact that theoretical simulations are often simplified and performed in vacuum without consideration of the chemical/physical environment that exists experimentally. Furthermore, the experimentally determined thermodynamic parameters are significantly affected by measurement methods and the experimental conditions used[126b] – hence, the significant disparities in experimental observations. It is noteworthy that particle size alone is not a determining factor for the reduction of $ΔH$ observed (Figure 21a). The experimentally determined $ΔH$ is greatly affected by the synthesis method which conditions the nanoparticle surface state, structure, and the atomic arrangements at the nanoparticle core.[126b] In an attempt to exclude parameters associated with synthetic variations, we investigated the evolution of thermodynamics as a function of the particle size reduction of the Mg/MgH$_2$ and LaNi$_5$/LaNi$_5$H$_6$ systems.[182a,309g] In both cases, the correlation of enthalpy and entropy was found to follow the linear “enthalpy-entropy compensation” effect as previously described for other chemical systems (Figure 21c,d). A similar trend was reported for Pd nanoparticles[325] and other hydrides.[309g,326] Such a compensation effect has also been observed experimentally for several chemical systems.[320]

Evolution of the entropy during hydrogen uptake and release mainly occurs through the conversion of hydrogen from the gaseous phase to its chemisorbed state in the hydride phase. This entropic value is often considered to be constant and equal to 130 J mol$^{-1}$ K$^{-1}$ for many hydrides.[27e] However, for nanomaterials, besides the dominant contribution of hydrogen, additional effects due to the higher degree of disorder in both the starting and final hydride material cannot be neglected and thus the effect of the additional $ΔS_{surf}$ factor in the van't Hoff Equation (79) must be considered.[316]

Owing to the complexity of determining entropic evolution by theoretical modeling, there have not been reports on possible enthalpy-entropy compensation effects by theoretical calculations. However, assuming that the experimentally observed “enthalpy-entropy compensation” effects are real, this does not favor lower hydrogen desorption temperatures ($T_{ibar} = \frac{\Delta H}{\Delta S}$) because any targeted reduction in temperature induced by $ΔH$ will lead to a similar reduction in $ΔS$. This in turn limits the increase in the equilibrium pressure. This correlation of $ΔH$ and $ΔS$ could be reversed through a distinct alteration of $ΔG$ as further discussed below (Figure 20a); however, strategies to do so have not been found yet.

### 8.1.2. Thermodynamics of Nanosized Complex Hydrides

The hydrogen absorption and desorption processes of complex hydrides are more complicated as compared to metal hydrides systems, since complex hydrides generally involve multistep reactions with additional side reactions generating byproducts.[123] Accordingly, it is very difficult to predict any thermodynamic evolution of the hydrogen properties of complex hydrides at the nanoscale.[16a] In addition, during hydrogen release, complex hydrides undergo significant entropic
<figure><img src="image_24.png" /><figcaption>**Figure 21.** Thermodynamic properties of metal hydrides at the nanoscale. a) Summary of reported values[182a,309a,b,d,326b] of ΔH for nanosized Mg/MgH₂ as function of Mg particle size; b) particle size dependent hydrogen desorption temperature of MgH₂ particles synthesized by chemical reduction;[182a] c) particle size dependence of ΔH and ΔS for Mg[182a] and LaNi₅[309g] when synthesized using the same method and similar conditions; d) correlation of ΔH and ΔS for Mg[182a] and LaNi₅.[309g]</figcaption></figure>

evolution owing to the number of products formed. Another typical characteristic of complex hydrides is their melting upon decomposition, e.g., LiAlH₄ melts at 150–175 °C and this does not simplify the behavior of such hydrides at the nanoscale.

There are various models that have been proposed to predict the melting temperature (T_nano) of nanoparticles.[327] For example, for spherical nanoparticles, Qi proposed that[328]

$$T_{\text{nano}} = T_{\text{bulk}} \left( 1 - \frac{2d}{D} \right) \qquad (81)$$

where T_bulk is the melting temperature of the bulk material, d is the diameter of the atom and D is the diameter of the nanoparticle.

Therefore, the melting point for nanoparticles is lower than that of the corresponding bulk material. For free standing complex hydride nanoparticles, this means that they would melt at lower temperatures and if this occurs before the release of hydrogen, the molten nanosized hydrides will simply behave like their bulk counterparts. One strategy to minimize this effect is to confine the nanoparticles within a stable inorganic scaffold.[186,329] By doing so, the diffusion of the decomposition products upon hydrogen release is further restricted and this facilitates hydrogen reversibility. Another approach is to suppress the melting behavior of complex hydrides through elemental doping.[330]

Assuming that the initial hydride nanoparticles retain their nanofeatures, the thermodynamics and kinetics of the hydrogen absorption and desorption reactions of nanosized

complex hydrides should deviate from the behavior of their bulk counterparts. In this case, the trends described above should apply to each reaction step involved in the dehydrogenation or re-hydrogenation process. Therefore, for individual reaction step at a given temperature a shift of the equilibrium pressure should be observed depending on the evolution of the nanoparticles and their disproportionation upon hydrogen release or uptake.

With NaAlH₄ as an example, it has been suggested from experimental observations that a shift of the equilibrium pressure occurs upon nanosizing and this leads to a destabilization of the hydride and the disappearance of the two step decomposition mechanism in favor of a single step (Figure 22a,b).[331] The theoretically calculated phase diagram and reaction map as function of particle size and temperature suggest similar assumptions.[332] In this case, the intermediate decomposition phase, Na₃AlH₆, would be destabilized upon nanosizing and this would result in a positive ΔGsurf value for the first decomposition step (NaAlH₄ → 1/3Na₃AlH₆ + 2/3Al + H₂) but a negative ΔGsurf value for the second step (1/3Na₃AlH₆ → NaH + Al + 1/2H₂)[331] leading to a single step decomposition from NaAlH₄ to NaH. The actual decomposition process of nanosized NaAlH₄ is probably more complex than this simplified assumption.[332a]

Similar thermodynamic alterations were observed for LiAlH₄ and LiBH₄ confined in carbon nanotubes (diameter 2 nm) with a single broad hydrogen release peak instead of multiple peaks corresponding to the multiple decomposition steps of the bulk materials.[314b] LiAlH₄ nanoparticles (2–16 nm) synthesized by