## 1 Introduction
However, because of the high chemical activity of the components and their surface oxidation in air hydrogen, interaction is limited. Activation of those alloys requires preliminary high-temperature (573-973 K) treatment [9; 21]. During the activation, the partial reduction of surface oxides, the formation of cracks and defects (dislocations, grain boundaries, violations of atomic packing) occur. The combination of these processes contributes to hydride formation [22; 23]. The key factors affecting activation parameters of HEAs are the valence state of surface atoms and the difference in the atomic sizes of constituting alloy elements [24]. In some cases, for example, for the arc-melted TiZrNbCrFeNi alloy with the C14 Laves phase structure, hydrogen absorption (up to 1.5 wt.%) at room temperature is observed without preliminary activation. However, desorption is possible only at 473 K; the reversible capacity does not exceed 1.1 wt.%. This is due to the formation of two phases, one of which is less active to accommodate hydrogen atoms [25]. The Ti-Zr-V-Cr-Ni alloy with a similar crystal structure demonstrates a higher reversible hydrogen sorption capacity (1.5 wt.%) and better cyclic stability, which indicates the future potential of developing this class of high entropy intermetallic-based materials for hydrogen storage [26]. Sometimes, during the formation of a hydride, the initial single-phase HEA transforms into a hydride consisting of several phases, which is probably determined by the elemental composition of the initial HEA. However, no general patterns can be traced in this case. It was found that, in addition to the main FCC phase, the tantalum-containing hydrides contain an additional BCT phase: 30.3 wt.% in TIVNbTaH\({}_{\rm x}\), 24.2 wt.% in TiVZr\({}_{0.15}\)NbTa\({}_{0.85}\)H\({}_{\rm x}\), 8.2 wt.% in TiVZr\({}_{0.5}\)NbTa\({}_{0.5}\)H\({}_{\rm x}\) and 2.2 wt.% in TiVZr\({}_{0.74}\)NbTa\({}_{0.26}\)H\({}_{\rm x}\)[27]. It has been shown that HEAs with a high content of 8-10 group (Fe, Co, Ni) metals tend to keep the initial C14 crystalline structure in their hydrides, while the volume of the unit cell can significantly increase with an almost unchanged ratio of c/a (for ZrTiVCrFeNi, the volumetric effect of hydride formation is 23.3% for an untreated alloy and 2.4% for a heat-treated alloy [28]). Mg-containing alloys tend to form hydrides with an FCC crystal structure; for example, BCC-HEA MgZrTiFe\({}_{0.5}\)Co\({}_{0.5}\)Ni\({}_{0.5}\) absorbs hydrogen in two stages: first, a solid solution with a BCC structure is formed, in which hydrogen occupies octahedral positions, and then the solid solution turns into a hydride with an FCC lattice, in which hydrogen atoms occupy tetrahedral positions [18]. Group 5 metals (V, Nb, Ta) with a bcc crystal structure have a high hydrogen permeability and are promising materials for membrane hydrogen separation [29; 30]. However, in most cases, individual metals are subjected to hydrogen embrittlement and passivation, which hinders their industrial use [31]. Nowadays, many works on developing HEA-based membranes whose properties will be superior to traditional Pd-based alloys are in progress. For example, at 673 K, the multicomponent Nb\({}_{40}\)Ti\({}_{18}\)Zr\({}_{12}\)Ni\({}_{25}\)Co\({}_{5}\) alloy has demonstrated a higher (3.8 \(\times\) 10\({}^{-8}\) mol H\({}_{2}\cdot\)m\({}^{-1}\cdot\)s\({}^{-1}\cdot\)Pa\({}^{-0.5}\)) permeability than the conventional Pd\({}_{0.23}\)Ag\({}_{0.77}\) (1.6 \(\times\) 10\({}^{-8}\) mol H\({}_{2}\) m\({}^{-1}\cdot\)s\({}^{-1}\cdot\)Pa\({}^{-0.5}\)) [32]. The thermochemical parameters of the interaction in metal-hydrogen systems are essential both for understanding the mechanism of hydrogen absorption and transport, and for practical application. For some HEAs, the enthalpy of hydride formation has been calculated on the basis of experimentally measured hydrogen absorption isotherms using the Van 't Hoff equation.