Hydrogen Storage Properties of Mg-Ni Alloys Processed by Fast Forging

: Fast forging of compacts made up of Mg and Ni powders is shown to be an e ﬀ ective method to induce severe plastic deformation with improved H 2 sorption properties. Here, using such processed samples, a comprehensive analysis of the sorption properties reveals that the ﬁrst hydrogenation sequence signiﬁcantly depends on the forging temperature, through di ﬀ erent microstructures. More in detail, no phase transformation occurs upon cold forging, while solid-state reaction leads to the formation of the Mg 2 Ni intermetallic compound upon forging above 400 ◦ C. Forging below the brittle-to-ductile transition (225–250 ◦ C) leads to faster H 2 uptake upon ﬁrst absorption owing to a more textured ﬁber along the c-axis and internal strains which promote hydrogen di ﬀ usion through the bulk material. Desorption kinetics remain slower with low-temperature forging, despite Ni recombining to form Mg 2 Ni during the ﬁrst desorption. After several cycles, a two-step behavior is observed with a fast absorption step occurring up to about 3 wt.%. Despite this limited uptake performance, the forging process can be considered as a straightforward, safe, and cost-e ﬃ cient process to produce large amounts of Mg-based alloys for hydrogen storage. In particular, such severe plastic deformation processes can be considered as reliable substitutes for ball-milling, which is highly e ﬃ cient but energy- and time-consuming.


Introduction
Among the materials relevant for solid hydrogen storage, magnesium hydride (MgH 2 ) stands out thanks to its high reversible hydrogen capacity and raw material availability [1]. Many studies have been devoted to magnesium hydrides and magnesium-based alloys to better understand and overcome their poor sorption kinetics which are detrimental for practical applications. With this aim, among processes that potentially improve hydrogen sorption properties, high-energy ball milling (HEBM) of MgH 2 powders-with or without the addition of transition metals-is broadly considered [2][3][4]. The resulting performance enhancement has been explained by a boosted hydrogen diffusion in the material resulting from the combination of reduced grain size, higher density of structural defects, and larger reactive surface area. However, the first hydrogenation of Mg, which is essential to produce the precursor MgH 2 , remains very slow and requires high temperature (>350 • C) and pressure up to 3 MPa [5]. Moreover, when large-scale production is desired, both the high pressure required for the first hydrogenation step and the HEBM process are prohibitive as they are time-and energy-consuming, and require significant manpower. Finally, handling such pyrophoric powders raises stringent safety concerns.
is larger than the Mg-Ni eutectic temperature (506 • C), which explains the formation of a Mg 2 Ni phase. The present paper reports on a detailed analysis of the sorption properties of fast forged Mg-Ni samples. In particular, we investigate the first H 2 absorption step which is found to strongly depend on the forging temperature. We also establish correlations between hydrogen sorption and solid microstructure. The sorption properties obtained by forging a pure magnesium sample are also presented in this paper to underline the key role of Ni addition.

Materials and Methods
Fast forging is performed at the laboratory scale using a homemade forge initially designed for hard magnets processing [28]. The device shown in [26] comprises a chamber under controlled atmosphere and a 150 kg hammer dropped from a height of 1.5 m on the upper piston of the chamber at a speed of 4.7 m/s. The sample, placed on an anvil in the chamber, is heated by means of a retractable high-frequency coil, which is pulled back just before forging (Figure 1a). The temperature is measured using an optical pyrometer with two wavelengths. Only one pass is applied to each sample. Samples were prepared by uniaxial compaction of powder using a pressure of 1 t/cm 2 . Magnesium powder was supplied by SFM SA in the form of atomized spherical particles of about 3-5 µm and nickel powder by Neco as particles of a few µm. All powders were handled in glove box under argon atmosphere. Mg-Ni pellets were prepared by manual grinding of these powders according to the composition Mg-10.45 at.% Ni. This hypoeutectic composition is close to the Mg-rich eutectic Mg-11.3 at.% Ni, but slightly lower in Ni in order to avoid the formation of the MgNi 2 phase. The expected hydrogen capacity of these samples is 6 wt.%, corresponding to 1.4 wt.% as Mg 2 NiH 4 +3.6 wt.% as MgH 2 when Ni is fully recombined with Mg, or 6 wt.% as MgH 2 , when Ni and Mg are not recombined.

Pure Magnesium Sample
A pure Mg sample prepared from atomized powder was forged at room temperature. Starting from a sample height of 20 mm, a thickness of 4.96 mm is achieved in one single pass, which corresponds to a reduction ratio of 75 %, and an equivalent von Mises εeq = 1.60. Figure 2 shows the XRD patterns acquired on polished surfaces (a) on the upper horizontal surface of the forged sample, and (b) on the vertical cross-section obtained by sawing the sample in two equal parts. The relative intensities of the diffraction peaks indicate a strong texture along the c-axis of the Mg phase for the horizontal surface. This texture was expected since the basal plane {0002} is the main slip plane of the hexagonal α-Mg phase. As reported by Jorge et al. [8], this texture corresponds to the best orientation for hydrogen sorption properties. While the atomized powder supplied by SFM SA consists of well-crystallized spherical particles, the XRD pattern performed after grinding of the forged sample into powder (not shown here) results in broader peak profiles, corresponding to crystallites size reduction (190 nm) and induced microstrain (0.14%). Each sample contained 3 pellets of 13 mm in diameter, which were introduced into a thin cylindrical stainless-steel tube of 20 mm high used to contain the material and to prevent its dispersion in the chamber during forging. The total mass of Mg or Mg-10.45 at.% Ni was of about 4.5 g per sample, depending on the composition. The forging process was applied under argon atmosphere at room temperature or at higher temperatures up to 530 • C. Samples were kept for 10 min at the target temperature to ensure thermal homogeneity. During forging, the height of the sample was reduced from 75% to 91%, depending on the composition and forging temperature. As shown in Figure 1b, the forged samples appear in the form of bulk materials surrounded by the steel sheath, which usually splits during forging.
The microstructure of the samples was investigated using a FE-SEM Zeiss Ultra Plus electronic microscope coupled with EDX analysis. The X-ray powder diffraction patterns were performed before Energies 2020, 13, 3509 4 of 14 and after hydrogenation using a Bruker D8 Endeavor diffractometer (Cu Kα radiation). Sample compositions, crystallite sizes, and lattice strain of the different phases were determined by Rietveld refinement using the Fullprof program [29]. Hydrogen absorption and desorption kinetics were measured using a volumetric Sievert device from HERA. In this device, the furnace, operated by a hydraulic cylinder, can be removed instantly. This makes it possible to cool down a sample quickly by cooling the sample holder with compressed air. The measurements were all made on a single piece of about 300-500 mg cut in the bulk sample and without any activation process. In this paper, the denomination "pressure cycle" is used for the successive absorption and desorption processes performed isothermally. More in detail, the pressure of hydrogen is set at 2 MPa for absorption and at 15 kPa for desorption. These two steps constitute a so-called pressure cycle, regardless of whether or not the sample absorbs hydrogen. The "reference of zero" for the uptake was reinitialized before each pressure cycle.

Pure Magnesium Sample
A pure Mg sample prepared from atomized powder was forged at room temperature. Starting from a sample height of 20 mm, a thickness of 4.96 mm is achieved in one single pass, which corresponds to a reduction ratio of 75 %, and an equivalent von Mises ε eq = 1.60. Figure 2 shows the XRD patterns acquired on polished surfaces (a) on the upper horizontal surface of the forged sample, and (b) on the vertical cross-section obtained by sawing the sample in two equal parts. The relative intensities of the diffraction peaks indicate a strong texture along the c-axis of the Mg phase for the horizontal surface. This texture was expected since the basal plane {0002} is the main slip plane of the hexagonal α-Mg phase. As reported by Jorge et al. [8], this texture corresponds to the best orientation for hydrogen sorption properties. While the atomized powder supplied by SFM SA consists of well-crystallized spherical particles, the XRD pattern performed after grinding of the forged sample into powder (not shown here) results in broader peak profiles, corresponding to crystallites size reduction (190 nm) and induced microstrain (0.14%).

Pure Magnesium Sample
A pure Mg sample prepared from atomized powder was forged at room temperature. Starting from a sample height of 20 mm, a thickness of 4.96 mm is achieved in one single pass, which corresponds to a reduction ratio of 75 %, and an equivalent von Mises εeq = 1.60. Figure 2 shows the XRD patterns acquired on polished surfaces (a) on the upper horizontal surface of the forged sample, and (b) on the vertical cross-section obtained by sawing the sample in two equal parts. The relative intensities of the diffraction peaks indicate a strong texture along the c-axis of the Mg phase for the horizontal surface. This texture was expected since the basal plane {0002} is the main slip plane of the hexagonal α-Mg phase. As reported by Jorge et al. [8], this texture corresponds to the best orientation for hydrogen sorption properties. While the atomized powder supplied by SFM SA consists of well-crystallized spherical particles, the XRD pattern performed after grinding of the forged sample into powder (not shown here) results in broader peak profiles, corresponding to crystallites size reduction (190 nm) and induced microstrain (0.14%). It was not possible to hydride this sample at 340 • C under 2 MPa of hydrogen pressure, even after 10 hydrogen pressure cycles (alternatively 2 MPa and 15 kPa). It was finally hydrided at 360 • C, but only at the third hydrogen pressure cycle ( Figure 3). A hydrogen uptake of 7.4 wt.% was achieved in 10 h, and the desorption reaction performed under 150 kPa was completed in 2 h at the same temperature of 360 • C. This difficulty in activating the forged sample, despite obtaining a favorable texture of the α-Mg phase, probably results from insufficient refinement of the microstructure with a single pass, but also from the fact that the sample obtained after forging was very dense and showed no porosity for the penetration of hydrogen. Even after activation, the following cycles performed at lower temperatures (not shown here) indicated very slow kinetics especially when desorption was performed at 310 • C.
It was not possible to hydride this sample at 340 °C under 2 MPa of hydrogen pressure, even after 10 hydrogen pressure cycles (alternatively 2 MPa and 15 kPa). It was finally hydrided at 360 °C, but only at the third hydrogen pressure cycle ( Figure 3). A hydrogen uptake of 7.4 wt % was achieved in 10 h, and the desorption reaction performed under 150 kPa was completed in 2 h at the same temperature of 360 °C. This difficulty in activating the forged sample, despite obtaining a favorable texture of the α-Mg phase, probably results from insufficient refinement of the microstructure with a single pass, but also from the fact that the sample obtained after forging was very dense and showed no porosity for the penetration of hydrogen. Even after activation, the following cycles performed at lower temperatures (not shown here) indicated very slow kinetics especially when desorption was performed at 310 °C.
In practice, when hydrogenation or desorption reactions are applied to a mass of several 100 g of magnesium, the experimental conditions in the tank tend to the thermodynamic equilibrium conditions [30]. Thus, when a magnesium hydride tank is used to supply a fuel cell with an outlet pressure of about 200 kPa, the temperature will immediately reach a value of about 310 °C. Then, cold forging of pure Mg, which leads to very slow kinetics at 310 °C, is not suitable for hydrogen storage.

Mg-Ni Eutectic Composition Forged at Room Temperature
The Mg-10.45 at.% Ni sample forged at room temperature (RT) led to the same reduction ratio as the sample of pure Mg (75%). The XRD pattern (not shown here) only reveals the presence of the initial Mg and Ni phases, without any traces of Mg2Ni and the size of the Mg crystallites calculated from the XRD pattern is similar to that of the pure Mg sample. Figure 4a shows the hydrogen sorption kinetics recorded at 310 °C. Conversely to the behavior of the pure Mg sample, it was possible to hydride this Ni-containing sample since the first application of hydrogen pressure, without any previous activation process and at a temperature as low as 310 °C. The first absorption did not present any incubation time. Fast initial kinetics were observed up to about 1 wt% of hydrogen, and then an acceleration of the reaction took place after about 1 h, to reach a hydrogen uptake of 5.4 wt.% in 5 h. Desorption was carried out in less than 1 h at 310 °C under 150 kPa of hydrogen pressure. However, it was not possible to complete the desorption reaction in these In practice, when hydrogenation or desorption reactions are applied to a mass of several 100 g of magnesium, the experimental conditions in the tank tend to the thermodynamic equilibrium conditions [30]. Thus, when a magnesium hydride tank is used to supply a fuel cell with an outlet pressure of about 200 kPa, the temperature will immediately reach a value of about 310 • C. Then, cold forging of pure Mg, which leads to very slow kinetics at 310 • C, is not suitable for hydrogen storage.

Mg-Ni Eutectic Composition Forged at Room Temperature
The Mg-10.45 at.% Ni sample forged at room temperature (RT) led to the same reduction ratio as the sample of pure Mg (75%). The XRD pattern (not shown here) only reveals the presence of the initial Mg and Ni phases, without any traces of Mg 2 Ni and the size of the Mg crystallites calculated from the XRD pattern is similar to that of the pure Mg sample. Figure 4a shows the hydrogen sorption kinetics recorded at 310 • C. Conversely to the behavior of the pure Mg sample, it was possible to hydride this Ni-containing sample since the first application of hydrogen pressure, without any previous activation process and at a temperature as low as 310 • C. The first absorption did not present any incubation time. Fast initial kinetics were observed up to about 1 wt.% of hydrogen, and then an acceleration of the reaction took place after about 1 h, to reach a hydrogen uptake of 5.4 wt.% in 5 h. Desorption was carried out in less than 1 h at 310 • C under 150 kPa of hydrogen pressure. However, it was not possible to complete the desorption reaction in these experimental conditions and an irreversible capacity of about 0.2 wt.% was observed at the end of the first and the second cycles. The following cycles, performed in the same experimental conditions, showed a very fast absorption reaction during the first minutes of reaction, followed by a strong reduction of the kinetics, which became even slower than during the first cycle (Figure 4b). Desorption reactions also exhibited a two-stage behavior, with about 1.5 wt.% desorbed within the 1st minute, and then slower kinetics.
Energies 2020, 13, x 6 of 14 experimental conditions and an irreversible capacity of about 0.2 wt.% was observed at the end of the first and the second cycles. The following cycles, performed in the same experimental conditions, showed a very fast absorption reaction during the first minutes of reaction, followed by a strong reduction of the kinetics, which became even slower than during the first cycle (Figure 4b).
Desorption reactions also exhibited a two-stage behavior, with about 1.5 wt.% desorbed within the 1st minute, and then slower kinetics.    at the end of the desorption step for the two first cycles, in Figure 4a. After three hydrogen cycles, the Mg2NiH0.3 phase appeared well crystallized and the magnesium phase was fully desorbed. In order to understand the two-stage behavior of the kinetics curves, the 4th absorption was stopped after 3 min only, by rapid cooling of the sample under hydrogen pressure. Most of the Mg2NiH0.3 phase was already transformed into Mg2NiH4 while less than one-third of the Mg was transformed into MgH2. The slow kinetics observed in the second stage of the reaction correspond to the hydrogenation of the remaining Mg, which seems not so well activated.

Mg-Ni Eutectic Composition Forged at Higher Temperatures
In order to promote a homogeneous distribution of thin Mg2Ni particles into the Mg matrix, 4 samples of the same Mg-10.45 at.% Ni hypoeutectic composition were forged at increasing temperatures up to 530 °C. The thickness and the corresponding reduction ratio are reported in Table 2 for each sample. As expected, the reduction ratio increased with the temperature and reached 91% at 530 °C. Figure 6 shows the XRD patterns acquired on the upper horizontal surface of the samples forged at increasing temperatures. The relative intensities of the diffraction peaks indicate a strong texture along the c-axis of the Mg phase, especially for the sample forged at 210 °C. This texture tended, however, to decrease when the forging temperature increased. The weight fractions calculated from the XRD patterns (Figure 7) are reported for each phase in Table 2. Up to a  In order to understand the two-stage behavior of the kinetics curves, the 4th absorption was stopped after 3 min only, by rapid cooling of the sample under hydrogen pressure. Most of the Mg 2 NiH 0.3 phase was already transformed into Mg 2 NiH 4 while less than one-third of the Mg was transformed into MgH 2 . The slow kinetics observed in the second stage of the reaction correspond to the hydrogenation of the remaining Mg, which seems not so well activated.

Mg-Ni Eutectic Composition Forged at Higher Temperatures
In order to promote a homogeneous distribution of thin Mg 2 Ni particles into the Mg matrix, 4 samples of the same Mg-10.45 at.% Ni hypoeutectic composition were forged at increasing temperatures up to 530 • C. The thickness and the corresponding reduction ratio are reported in Table 2 for each sample. As expected, the reduction ratio increased with the temperature and reached 91% at 530 • C.   Figure 6 shows the XRD patterns acquired on the upper horizontal surface of the samples forged at increasing temperatures. The relative intensities of the diffraction peaks indicate a strong texture along the c-axis of the Mg phase, especially for the sample forged at 210 • C. This texture tended, however, to decrease when the forging temperature increased. The weight fractions calculated from the XRD patterns (Figure 7) are reported for each phase in Table 2. Up to a temperature of 390 • C, the forging process did not produce significant amounts of Mg 2 Ni. The microstructure was very similar to the one of the sample forged at room temperature, with agglomerated Ni particles embedded in Mg matrix (Figure 8a,b). The size of these Ni agglomerates, approximately 2 µm in diameter, corresponds to the particle size of the Ni powder used as a precursor.
Energies 2020, 13, x 8 of 14 temperature of 390 °C, the forging process did not produce significant amounts of Mg2Ni. The microstructure was very similar to the one of the sample forged at room temperature, with agglomerated Ni particles embedded in Mg matrix (Figures 8a,b). The size of these Ni agglomerates, approximately 2 µm in diameter, corresponds to the particle size of the Ni powder used as a precursor.    (Table 3), the dispersion between different points of measurements being very low. The fraction of Mg 2 Ni reached 35 wt.% (Table 2), while the expected amount would be 43 wt.% if all the nickel was recombined with magnesium. At higher magnification (Figure 8d), the dark grey areas observed inside the Mg 2 Ni particles reveal a specific microstructure characteristic of eutectic formation, indicating that partial melting occurred locally. To determine whether the formation of Mg 2 Ni can be attributed to the effect of fast deformation, or simply to the fact that the alloy was exposed to a high temperature, a control experiment of static annealing was performed at 480 • C for 10 min, which is the delay applied to stabilize the temperature before forging. XRD analysis of this sample revealed no trace of Mg 2 Ni after annealing (supplementary work). The formation of this phase indeed results from the fast forging itself, and is due to the heat released during the mechanical deformation, which induces an overpass of the melting temperature of the eutectic (506 • C) for a short time, as already reported in [27].   Energies 2020, 13, x 9 of 14 nickel was recombined with magnesium. At higher magnification (Figure 8d), the dark grey areas observed inside the Mg2Ni particles reveal a specific microstructure characteristic of eutectic formation, indicating that partial melting occurred locally. To determine whether the formation of Mg2Ni can be attributed to the effect of fast deformation, or simply to the fact that the alloy was exposed to a high temperature, a control experiment of static annealing was performed at 480 °C for 10 min, which is the delay applied to stabilize the temperature before forging. XRD analysis of this sample revealed no trace of Mg2Ni after annealing (supplementary work). The formation of this phase indeed results from the fast forging itself, and is due to the heat released during the mechanical deformation, which induces an overpass of the melting temperature of the eutectic (506 °C) for a short time, as already reported in [27]. When forging is performed at 530 °C, above the melting temperature of the eutectic composition, Mg2Ni formation occurs while heating the sample, as confirmed by the XRD pattern recorded after annealing a sample at 530 °C for 10 min (supplementary work). After forging, both XRD (Table 2) and EDX analyses (Table 3) indicate a complete recombination of the Ni into Mg2Ni. The microstructure observed by SEM (Figure 8e) reveals well-faceted and homogenously dispersed Mg2Ni plates of about 10 µm length or more. This microstructure is completely different from what is observed at lower temperatures, as well as from the characteristic microstructure expected for a eutectic composition. Moreover, the Mg2Ni crystals appear fragmented (Figure 8f), indicating that When forging is performed at 530 • C, above the melting temperature of the eutectic composition, Mg 2 Ni formation occurs while heating the sample, as confirmed by the XRD pattern recorded after Energies 2020, 13, 3509 10 of 14 annealing a sample at 530 • C for 10 min (supplementary work). After forging, both XRD (Table 2) and EDX analyses (Table 3) indicate a complete recombination of the Ni into Mg 2 Ni. The microstructure observed by SEM (Figure 8e) reveals well-faceted and homogenously dispersed Mg 2 Ni plates of about 10 µm length or more. This microstructure is completely different from what is observed at lower temperatures, as well as from the characteristic microstructure expected for a eutectic composition. Moreover, the Mg 2 Ni crystals appear fragmented (Figure 8f), indicating that they were grown before applying the forging shock. Whatever the forging temperature, the EDX analyses indicate the presence of a small amount of Ni in the Mg phase and some Mg in the Ni phase ( Table 3). The accuracy of the EDX measurements does not allow us to state whether these elements are indeed present in solid solution. However, the cell parameters calculated from XRD do not indicate significant deviations from the pure metallic phases and these observations most probably result from the large probe area.
Hydrogen pressure cycles were performed on each sample at the same pressures (2 MPa and 150 kPa) but at 325 • C instead of 310 • C, because of the poor reactivity at 310 • C of the samples forged in temperature. During the first absorption (Figure 9a), the kinetic appeared very slow for the sample forged at 530 • C, and no absorption was observed for the one forged at 480 • C. On the other hand, the sample forged at 210 • C absorbed about 1 wt.% in 30 min. Low-temperature forging and even RT forging are then more effective for activation. This behavior was not expected since the samples forged at RT or low temperature did not contain Mg 2 Ni phase. However, as reported in Table 2, a decrease of the internal strains with the increase of the temperature was observed in the three phases. The brittle-to-ductile transition temperature of magnesium is of about 225-250 • C [31]. Thus, forging in the brittle state induces a higher amount of internal strains, and probably the formation of cracks and/or structural defects that promote the diffusion of hydrogen into the bulk material, while forging in the ductile state reduces porosity and does not induce cracks. The higher texture observed when forging at lower temperatures could also induce a faster diffusion of hydrogen into the material, as this texture corresponds to the best orientation for hydrogen sorption properties [8].
Absorption kinetics increased rapidly during the following pressure cycles (Figure 9b) to converge to the same curves after about 12 cycles (Figure 9c), whatever the forging temperature. As for the sample forged at room temperature, the formation of solid solution Mg 2 NiH 0.3 induced a slight irreversibility at the end the first cycles (not shown here). This irreversibility was no more visible at the end of the third cycles (Figure 9b), which means that Ni was already fully recombined as Mg 2 NiH 0.3 at the end of the second cycles (the "zero" reference for the uptake was reinitialized before each cycle). Again, a two-stage behavior appeared progressively during the absorption, with a very fast kinetic at the beginning of the reaction (about 2 wt.% in a few seconds) followed by a much slower absorption and a hydrogen uptake limited to 3.5 wt.% in 1 h. The hydrogenation of the Mg 2 NiH 0.3 phase into Mg 2 NiH 4 corresponds to a maximum uptake of 1.3 wt.%. This means that the first absorption stage was not only due to the hydrogenation of this phase, and that at least a fraction of the pure magnesium phase remained activated after a large number of cycles. Surprisingly, desorption kinetics became faster when forging at higher temperatures (Figure 9c). This behavior could result from a better adherence between Mg and Mg 2 Ni particles when the Mg 2 Ni phase is formed directly during the forging process. This would enable a catalytic role of the Mg 2 NiH 0.3 or Mg 2 NiH 4 during the dehydrogenation of the MgH 2 phase. formed directly during the forging process. This would enable a catalytic role of the Mg2NiH0.3 or Mg2NiH4 during the dehydrogenation of the MgH2 phase.
EBSD and in situ neutron diffraction experiments are in progress in order to better analyze the induced texture, to understand the two-step behavior of the absorption reaction, and to seek to increase the part of magnesium that remains activated after further hydrogen cycles.

Conclusions
Magnesium is a promising material for hydrogen storage but its use-at least in its unprocessed form-is hampered by slow sorption kinetics. Here, we show that Mg sample forging leads to improved properties but sorption kinetics remain rather slow even after several hydrogen cycles (almost 2 h to desorb at 360 °C, 150 kPa). In contrast, introducing nickel into the metal sample drastically improves its sorption properties. EBSD and in situ neutron diffraction experiments are in progress in order to better analyze the induced texture, to understand the two-step behavior of the absorption reaction, and to seek to increase the part of magnesium that remains activated after further hydrogen cycles.

Conclusions
Magnesium is a promising material for hydrogen storage but its use-at least in its unprocessed form-is hampered by slow sorption kinetics. Here, we show that Mg sample forging leads to improved properties but sorption kinetics remain rather slow even after several hydrogen cycles (almost 2 h to desorb at 360 • C, 150 kPa). In contrast, introducing nickel into the metal sample drastically improves its sorption properties.
Low-temperature forging. Forging Mg-Ni samples in the brittle state (T < 225-250 • C) is found to be the most efficient method to activate samples (first absorption). Considering that no Mg 2 Ni forms upon forging at such temperature, this activation improvement is assumed to result from the forging-induced texture along the c axis together with the formation of internal strains, fractures, and/or structural defects that promote hydrogen diffusion through the bulk material. The unalloyed nickel combines into Mg 2 NiH 0.3 during the first desorption reaction. Considering the slow atomic mobility at the desorption temperature (T = 325 • C), the formation mechanism of Mg 2 NiH 0.3 remains to be understood.
High-temperature forging. Forging above T = 400 • C leads to the formation of Mg 2 Ni during the process with resulting microstructures strongly dependent on the forging temperature. Above the eutectic temperature, the growth of Mg 2 Ni crystallites occurs prior to the forging deformation, during the thermal stabilization. In contrast, as the forging temperature is decreased, the formation of Mg 2 Ni is likely to result from solid-state reaction as a result of local (i.e., spatially restricted) alloying/recrystallization during the mechanical deformation. However, regardless of the forging temperature, despite different resulting microstructures, similar sorption behaviors are observed after a few cycles. These results, which indicate that the forging temperature is mainly a critical parameter for the activation process, are promising since forging processes can be easily scaled up to produce very large sample amounts. At this stage, the hydrogen capacity of samples obtained using such techniques does not compete with HEBM. However, this simple and economical process is a mature technology that enables mass production without handling pyrophoric powders.
Author Contributions: Conceptualization P.d.R., D.F. and N.S.; processing by Fast Forging, P.d.R. and J.W.; investigation and analysis, P.d.R., J.W. and M.B.; writing-original draft preparation, P.d.R., D.F., N.S., J.W. and L.L.; writing-review and editing, P.d.R. and L.L. All authors have read and agreed to the published version of the manuscript.
Funding: This research was funded by the CDP Eco-SESA, from the University Grenoble Alpes IDEX.