Compare to other SPD techniques, High Pressure Torsion (HPT) is a relatively simple and quick processing technique. It is also very efficient to produce small grain size [3]. The principle of HPT is schematically illustrated in Figure 1.

The sample, generally in the form of a thin disk, is located between the piston and the anvil within a cavity. A hydrostatic pressure is applied and plastic torsional straining is achieved by rotation of the piston. The diameter of the cylindrical cavities and the initial diameter of the HPT sample are identical. However, the sum of both depths of the cavities is somewhat smaller than the initial height of the HPT sample. This implies that during loading a small amount of the material will flow in the ring shaped region between the two anvils. The friction in this region confines the free flow of the material out of the HPT tool and leads to a back pressure and induces a defined hydrostatic pressure within the processing zone. More details of this technique could be found in reference [4].

Application of HPT to hydrogen storage materials is relatively new. One of such study was performed by Kusadome et al. on MgNi₂ alloy [5] which does not usually absorb hydrogen. The crystallite size of MgNi₂ phase changed from more than 100 nm before HPT processing ton about 20 nm after 10 HPT revolutions. Moreover, the HPT processing introduced a substantial strain into the material, and MgNi₂ became a weak hydrogen absorber (0.1 wt.%) with hydrogen accumulating in grain boundaries of the alloy [5].

A certain similarity between HPT and ball-milling can be made based on the recent work by Leiva et al., who detected the formation of a metastable phase γ-MgH₂ and a substantial reduction of crystallite sizes from 30 to 20 nm while using HPT to consolidate the metal hydride powder [6]. Previously, the formation of γ-MgH₂ upon mechanical processing was reported for and chiefly associated with high-energy ball-milling [7]. HTP has also been used for Mg-Fe and Mg-Ni compounds. In the case of Mg–Fe system, Lima et al. were able to synthesize the ternary complex Mg₂FeH₆ and the binary MgH₂ hydrides by hydrogenation treatment at 350 °C, at 3 MPa during 24 h [8]. However, the hydride formation was not complete because after hydrogenation XRD pattern showed the presence of Mg, Fe, MgH₂ and Mg₂FeH₆. Révész et al. subjected 7Mg + 3Ni powders to ball milling followed by High Pressure Torsion (HPT) [9]. The result was a microstructural refinement accompanied with the creation of straining and/or lattice defects. Absorption kinetic data indicated that the maximum H-absorption capacity was increased by 30–50% after HPT due to the creation of new possible hydrogen absorption sites at the grain boundaries and at lattice defects. On the other hand, the initial absorption rate was somewhat decreased due to HPT [9].

BCC alloys of systems Ti–V–Mn and Ti–V–Cr have been intensively studied for hydrogen storage [37,38,39,40,41,42]. They could also be used as catalyst for magnesium [43]. Hydrogen capacities as high as 3.6 wt.% have been reported for Ti₂₅V₄₀Cr₃₅ alloy.

The effect of severe plastic deformation (SPD) on BCC Ti–22Al–27Nb alloy has been investigated by Zhang et al. [19,20]. They showed that the first hydrogenation (activation) was much faster for the deformed alloy compared to the as-quenched sample. The deformed alloy had also a faster absorption/desorption kinetics. However, the beneficial effect of deformation was lost after a few hydrogenation cycles. In these studies, SPD was obtained by cold rolling or compression. In the case of cold rolling one rolling was performed at 10.5% and 80% thickness reduction. Some of the 80% rolled specimen were further rolled to 10% thickness reduction in a perpendicular direction with respect to the first rolling.

Huot et al. have made a systematic study of the effect of milling on TiV_0.9Mn_1.1 alloy [44]. This composition is interesting to study because the as-cast alloy is a mixture of BCC and C14 phases. Therefore, it is a good system to test the effect of milling on the crystalline change and the interaction between phases. Milling was performed on as-cast alloy as well as on mixtures of elemental powders. Figure 6 shows the effect of milling on as-cast TiV_0.9Mn_1.1.

The presence of NaCl Bragg peaks is explained by the use of a small amount of this salt as an anti-sticking agent. It is clear that, with milling time, the C14 phase vanishes and a FCC phase appears. From Rietveld refinement it was found that, for the sample milled 80 h, the crystal structure is a mixture of cubic (FCC) solid solution phase and a BCC solid solution. The coexistence of FCC and BCC has been seen in the system Fe–Cu and was explained by an enhanced solubility due to the high dislocation density [45] but this explanation may not be applicable in the present case. From Rietveld refinement it was determined that the crystallite size of the BCC phase goes from 44 nm before milling to 2 nm after 80 h of milling. When milling was performed on the raw elements (Ti, V, and Mn) an identical result was obtained, that is formation of a nanocrystalline alloy composed of BCC and FCC phases [44].

Before measurement of hydrogen storage properties BCC alloys have to be activated. In the present case, the activation was made by cycling between high hydrogen pressure (5 MPa) and vacuum at elevated temperature (up to 250 °C). In Figure 7 the hydrogen absorption and desorption isotherm (23 °C) for arc-melted TiV_0.9Mn_1.1 before and after 80 hours of milling is presented. The maximum capacity of the as melted alloy is 1.9 wt% at 7 MPa which corresponds to a H/M ratio of 0.97. After 80 h of milling, the alloy does not absorb hydrogen up to 7 MPa. Because the as-milled materials present both FCC and BCC phases this means that both of them do not absorb hydrogen. In the case of BCC phase the reason may be reduction of lattice parameters. Iron contamination (even at this low level) may also play a role as shown by Santos et al. in Ti–V–Cr system [46].

Singh et al. studied the effect of milling on arc melted Ti_0.32Cr_0.43V_0.25 alloy [47]. As they used tungsten carbide balls, some contamination was observed after long milling time. Ball milling did not affect the apparent bulk structure of the alloy. Increase of ball milling time resulted in the increase in lattice strain and the decrease in crystallite size, which in turn increased subgrain boundaries. These contamination and microstructural changes caused an important decrease in the hydrogen storage capacity [47].

Amira et al. compared the effect of ball milling and cold rolling for Ti–Cr system [23]. Unlike ball milling, cold rolling of TiCr_x (x = 2, 1.8, 1.5) did not lead to the formation of metastable BCC phase. However, cold rolling was found to be effective to form nanocrystalline C14 Laves phase. Hydrogen sorption experiments showed that cold rolled alloys have similar hydrogen sorption properties than their ball milled counterparts despite different crystal structures. This result proves that hydrogen absorption/desorption properties do not depend only of the microstructure.

The alloy TiV_1.6Mn_0.4 has been recently investigated by Couillaud et al. [22]. The effect of extended cold rolling as well as energetic ball milling was a reduction of crystalline size and lattice parameter but no change in the crystal structure. Figure 8 shows TEM micrographs of TiV_1.6Mn_0.4 alloy in the as-cast, milled, and cold rolled states.

The dark field image shows that all samples are nanocrystalline. The bright field image of the cold rolled sample clearly shows the pile-up of dislocations. The dark field image shows that, contrary to the as-cast and ball milled samples, the crystallites tend to be aligned along dislocations. Figure 9 shows the X-ray diffraction patterns of as-cast, milled 5 hours and cold rolled 150 times TiV_1.6Mn_0.4 alloy.

From these X-ray powder diffraction patterns it was determined that the crystallite size of as-cast, milled, and cold rolled samples are respectively 17, 11, and 13 nm [22]. Apart from peak broadening due to the reduction of crystallite size, the pattern of the ball milled sample has the same relative intensities and lattice parameter as the as-cast sample. For the cold rolled sample, the lattice parameter is also the same as the as-cast sample but there is a very strong texture along (200). Texturization is a common feature of cold rolled samples. Neither ball milled sample nor cold rolled sample absorb hydrogen even after 10 cycles of hydrogen pressurization (10 MPa) and vacuum at 423 K. The reason for this significant loss of hydrogen capacity is still not clear.

Recently, Mg-based BCC alloys have been developed in order to achieve higher hydrogen storage capacity on a weight basis than Ti based ones. In particular, Akiba’s group has made an extensive study of the synthesis of Mg–Ti [48,49], Mg–Co [50,51,52] and Mg–Ni [53,54] BCC alloys by means of ball milling. In this review we will limit our discussion to the Mg–Ti system.

Binary Mg–Ti alloys are being intensively investigated for various applications such as: negative electrodes for Ni–MH batteries, [55,56], H₂sources for fuel cells [49,57], switchable mirrors for smart solar collectors [58,59], and optical hydrogen detectors [60]. In the phase diagram of the Mg–Ti binary system, equilibrium solid solubility of each metal to another is less than 2 at.% and no intermetallic compound is found. Therefore, non-conventional synthesis methods have to be used. Metastable single-phase Mg–Ti thin films have been successfully synthesized over a large composition range by means of electron-beam and magnetron co-sputter deposition techniques [55,58,59,61,62]. However, these techniques could not be scaled-up to industrial level and other methods have to be investigated. Mechanical alloying has demonstrated its high efficiency for producing metastable Mg–Ti alloys starting from elemental Mg and Ti powders [48,56,63,64,65,66].

The synthesis of Mg–Ti BCC alloys by mechanical alloying has been extensively studied by Asano et al. [48,49,57,65,66]. Although both Mg and Ti have a hexagonal closed packed (HCP) structure, during milling of a Mg and Ti mixture they react differently. In the case of magnesium, the deformation is mainly by basal plane slip (0001)< −12–10> while for titanium twinning deformation is more important [65]. In one investigation, Asano et al. had the idea of adding lithium to magnesium in order to reduce the yield stress of magnesium and also to decreases its lattice parameter [65]. They found that by adding Li to Mg the deformation of Mg was easier and the Ti crystallite size was reduced. These led to a decrease of synthesis time of BCC phase. In a subsequent study, they first synthesized a BCC Mg₅₀Ti₅₀ alloy by ball milling a mixture of 50Mg + 50Ti in a Fritsch P5 planetary ball mill for 150 h at a rotation speed of 200 rpm. Figure 10 confirms that a BCC phase was obtained and from the peaks width a crystallite size of 3 nm was calculated [49].

A full hydrogenation at 423 K under 8 MPa of hydrogen and for 122 h resulted in the formation of Mg₄₂Ti₅₈H₁₇₇ FCC hydride phase and some MgH₂.

By controlling milling conditions and Mg:Ti ratio, Asano et al. have also shown that BCC, FCC or HCP phase could be obtained in the Mg–Ti system [48]. In the case of HCP phase it is formed by solution of Ti into Mg while the BCC phase is produced by solution of Mg into Ti and the FCC phase is stabilized by introduction of stacking faults in Mg and Ti which have a HCP structure [48,49]. If MgH₂ is used instead of Mg as the starting material then, after ball milling a 50MgH₂ + 50Ti mixture the resulting compound is FCC Mg₃₃Ti₅₀H₉₄ plus some MgH₂ [57]. The importance of mechanical effect during milling is discussed in reference [66]. It shows that during ball milling of Mg and Ti powders in molar ratio of 1:1, plate-like particles first stuck on the surface of the milling pot and balls. After these plate-like particles fell off from the surface of the milling pot and balls, spherical particles in which concentric layers of Mg and Ti are disposed, are formed. These particles have a mean diameter of 1 mm. These spherical particles are then crushed into spherical particles with diameter of around 10 μm by introduction of cracks along the boundaries between Mg and Ti layers. Finally, the Mg₅₀Ti₅₀ BCC phase with a lattice parameter of 0.342(1) nm and a grain size of 3 nm is formed. During milling, Ti acts as an abrasive for Mg which had stuck on the surface of the milling pot and balls [66].

Recently, Çakmak et al. showed that mechanical milling of Mg–10 vol.% Ti yields large Mg agglomerate, 90–100 μm, with embedded Ti fragments of about 1 μm uniformly distributed within the agglomerates [67]. These Mg agglomerates are made of coherently diffracting volumes of small size (crystallites). These crystallites, as determined with X-ray diffraction analysis, can be as small as 26 nm after 30 h of milling.

In an investigation of high energy milling of 50Mg–50Ti mixture, Maweja et al. observed twinning in Ti-rich crystallites at intermediate milling time [68]. They attributed the twinning to the deformation of Ti particles. But they also pointed out that in the Mg–Ti system it might also indicate a strain induced martensitic transformation of the metastable ω-FCC into BCC. The crystallite boundaries acted as preferential sites for the heterogeneous nucleation of the twins and for the formation of solid solution by release of the lattice strain energy [68].

For electrochemical applications, mechanical alloyed Mg–Ti materials must be activated by adding of few at.% of Pd. Rousselot et al. have shown that if a 50Mg–50Ti mixture is premilled before adding Pd then The alloying of Pd with pre-milled Mg₅₀Ti₅₀ occurs very rapidly (few minutes) and is complete after 5 h of milling [69]. They also found that the crystalline structure of the Mg₅₀Ti₅₀ alloy (bcc and hcp Mg–Ti phase mixture) does not change significantly with the addition of Pd.