Understanding of hydrogen desorption mechanism from defect point of view

Widespread adoption of hydrogen for vehicle applications critically depends on the discovery of solid materials to store hydrogenwith high volumetric and gravimetric densities, as well as to extract and insert it at sufficiently rapid rates. Complex hydrides (such as NaAlH4, LiAlH4, LiBH4, Ca(BH4)2, Li4BN3H10 and LiNH2), light metal hydrides of MgH2 and AlH3 and ammonia borane (NH3BH3) have recently received considerable attention as potential hydrogen storage materials due to their high hydrogen gravimetric capacity as well as volumetric capacity.However, all of these materials suffer from unfavorable dehydrogenation thermodynamics and/or slow dehydrogenation kinetics, limiting their practical applications. Yet, the harsh and ill-controlled dehydrogenation process might be overcome eventually by carefully controlling and catalyzing the hydrogen release process, calling for an in-depth understanding of dehydrogenation mechanisms. Dehydrogenation in the abovementioned materials involves the bond-breaking process and migration of constituent atoms, which can be accomplished through the creation and subsequent diffusion of some native point defects. Furthermore, native point defects may also serve as nucleation sites for new phases that are formed in the process of dehydrogenation reactions. Thus a deep understanding of the nature of point defects is necessary. Computational studies based on density functional theory (DFT) play an important role in defect-related properties and may help to understand the details of the hydrogen release process at the atomic level. In this perspective, we provide information about point defects to understand the dehydrogenation properties of hydrogen storage materials with a band gap. The defect properties in metallic hydrogen storage materials [1] are not within the scope of this brief review. Prof. Van de Walle’s group from the University of California, Santa Barbara, has done some pioneering computational work on point defects in NaAlH4 [2,3]. Since this work, the properties of defects, especially H-related defects, in Na3AlH6 [4], LiAlH4 [5], LiBH4 [6], Li4BN3H10 [6–8], LiNH2 [9,10], MgH2 [11,12] and AlH3 [13] have been investigated computationally. One key quantity in defect physics is the formation energy for the defect. According to the formation energies of defects, it is possible to find out which kinds of defects are predominant in lattices. The above-mentioned hydrogen storage materials are classified as insulators with wide band gaps; thus, the point defects in the lattice may be in charged states and the formation energies are Fermilevel dependent. In first-principles calculations, the defect is usually surrounded by a few dozen to a few hundred atoms, i.e. the so-called supercell approach. For a defect with a charge q (including its sign), its formation energy Ef is given by [14]

Widespread adoption of hydrogen for vehicle applications critically depends on the discovery of solid materials to store hydrogen with high volumetric and gravimetric densities, as well as to extract and insert it at sufficiently rapid rates. Complex hydrides (such as NaAlH 4 , LiAlH 4 , LiBH 4 , Ca(BH 4 ) 2 , Li 4 BN 3 H 10 and LiNH 2 ), light metal hydrides of MgH 2 and AlH 3 and ammonia borane (NH 3 BH 3 ) have recently received considerable attention as potential hydrogen storage materials due to their high hydrogen gravimetric capacity as well as volumetric capacity. However, all of these materials suffer from unfavorable dehydrogenation thermodynamics and/or slow dehydrogenation kinetics, limiting their practical applications. Yet, the harsh and ill-controlled dehydrogenation process might be overcome eventually by carefully controlling and catalyzing the hydrogen release process, calling for an in-depth understanding of dehydrogenation mechanisms.
Dehydrogenation in the abovementioned materials involves the bond-breaking process and migration of constituent atoms, which can be accomplished through the creation and subsequent diffusion of some native point defects. Furthermore, native point defects may also serve as nucleation sites for new phases that are formed in the process of dehydrogenation reactions. Thus a deep understanding of the nature of point defects is necessary. Computational studies based on density functional theory (DFT) play an important role in defect-related properties and may help to understand the details of the hydrogen release process at the atomic level. In this perspective, we provide information about point defects to understand the dehydrogenation properties of hydrogen storage materials with a band gap. The defect properties in metallic hydrogen storage materials [1] are not within the scope of this brief review. Prof. Van de Walle's group from the University of California, Santa Barbara, has done some pioneering computational work on point defects in NaAlH 4 [2,3]. Since this work, the properties of defects, especially H-related defects, in Na 3 AlH 6 [4], LiAlH 4 [5], LiBH 4 [6], Li 4 BN 3 H 10 [6][7][8], LiNH 2 [9,10], MgH 2 [11,12] and AlH 3 [13] have been investigated computationally.
One key quantity in defect physics is the formation energy for the defect. According to the formation energies of defects, it is possible to find out which kinds of defects are predominant in lattices. The above-mentioned hydrogen storage materials are classified as insulators with wide band gaps; thus, the point defects in the lattice may be in charged states and the formation energies are Fermilevel dependent. In first-principles calculations, the defect is usually surrounded by a few dozen to a few hundred atoms, i.e. the so-called supercell approach. For a defect with a charge q (including its sign), its formation energy E f is given by [14] where E(X q ) and E(bulk) denote the total energies of the supercell containing a defect X in the charged state q and of the defect-free supercell, respectively. n i indicates the number of atoms of type i (host atoms or impurities) that have been added to (n i > 0) or removed from (n i < 0) the supercell upon defect creation, and μ i are the corresponding chemical potentials of these species, which represent the energy of the reservoirs with which atoms are being exchanged. ε F is the chemical potential of the electrons or Fermi level, which accounts for exchanging electrons with an electron reservoir. ε F is conventionally taken with respect to the valence band maximum (VBM) E V of the perfect lattice and can vary from VBM to conduction band minimum (CBM). The last term of equation (1) is a correction term that accounts for the electrostatic interactions between supercells [14].
As shown in equation (1), the formation energies of defects depend on the atomic chemical potentials and the Fermi level. The chemicals reflect the reservoirs for atoms that are involved in creating the defect. The dehydrogenation conditions under which the defects are created uniquely define the relevant reservoirs. So the chemical potentials can be determined by local equilibrium conditions in the dehydrogenation. For example, the first step of desorption of NaAlH 4 is: (2) The local equilibrium conditions for dehydrogenation of NaAlH 4 could be (i) equilibrium with NaAlH 4 , Na 3 AlH 6 and Al; (ii) equilibrium with NaAlH 4 , Na 3 AlH 6 and H 2 ; or (iii) equilibrium with NaAlH 4 , Al and H 2 [3]. For equilibrium condition (i), we have: where the chemical potentials of μ NaAlH4 and μ Na3AlH6 are the free energy of a formula unit of NaAlH 4 and Na 3 AlH 6 , respectively. μ 0 Al is the chemical potential of Al in its standard state. By solving equations (3)-(5) the chemical potentials of Na, Al and H under condition (i) can be obtained. More details on the discussion of chemical potentials can be found in [3,8].
First-principles calculations have indicated that H-related point defects in NaAlH 4 [2], LiAlH 4 [5], LiBH 4 [6], Li 4 BN 3 H 10 [6,7], LiNH 2 [9,10], MgH 2 [11,12], AlH 3 [13] and NH 3 BH 3 are charged, and the neutral H vacancy or interstitial is higher in energy than the corresponding charged one. Figure 1 shows the formation energy of H-related defects in NaAlH 4 , and Table 1 summarizes the predominant H-related defects in some novel hydrogen storage materials. The formation energies of charged H-related defects in the above materials are Fermi-level dependent; thus hydrogen release kinetics may be adjusted by adding electrically active impurities to shift the equilibrium position of the Fermi level determined by the  [13] V + H and V − H H-rich condition condition of charge neutrality [2,14]. It has been found that [2] adding Zr impurity to NaAlH 4 has a tiny effect on the formation energies of H-related defects, but incorporation of a small fraction of Ti in it can result in shifting the Fermi level from its equilibrium value in intrinsic material by 0.44 eV. Correspondingly, this shift makes the formation energy of the predominant H-related point defect of H − i decrease by 0.44 eV. This finding might account for the observation of the improved hydrogen release kinetics in experiment when Ti-based species were balled with NaAlH 4 . The role of impurities in LiBH 4 , Li 4 BN 3 H 10 , MgH 2 and NaMgH 3 on hydrogen release has also been addressed in the literature [6,15].
Another important issue related to hydrogen release is the diffusion of defects, especially for the diffusion barrier. The diffusion barriers for the predominant defects can be directly obtained by the nudged elastic band (NEB) method or the dimer method. For NaAlH 4 [2]. The substantially low value of the diffusion barriers indicates that defect migration is not a rate-limiting step for self-diffusion of hydrogen, and that the formation of