Progress in Materials Science

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1 Progress in Materials Science 58 (2013) Contents lists available at SciVerse ScienceDirect Progress in Materials Science journal homepage: Mechanochemical synthesis of hydrogen storage materials J. Huot a, D.B. Ravnsbæk b, J. Zhang c, F. Cuevas c,, M. Latroche c, T.R. Jensen b a Université du Québec à Trois-Rivières, 3351 des Forges, Trois-Rivières, Québec, Canada G9A 5H7 b Center for Materials Crystallography (CMC), Interdisciplinary Nanoscience Center (inano), Department of Chemistry, Aarhus University, Langelandsgade 140, DK-8000 Århus C, Denmark c ICMPE, CNRS, UMR 7182, 2-8 rue Henri Dunant, Thiais Cedex, France article info abstract Article history: Received 13 April 2012 Accepted 9 July 2012 Available online 27 July 2012 New synthesis methods are of utmost importance for most materials science research fields. The present review focuses on mechanochemical synthesis methods for solid hydrogen storage. We anticipate that the general methods and techniques are valuable with a range of other research fields, e.g. the rapidly expanding fields of energy materials science and green chemistry including solvent free synthesis. This review starts with a short historical reminder on mechanochemistry, followed by a general description of the experimental methods. The use of milling tools for tuning the microstructure of metals to modify their hydrogenation properties is discussed. A section is devoted to the direct synthesis of hydrogen storage materials by solid/gas reactions, i.e. by reactive ball milling of metallic constituents in hydrogen, diborane or ammonia atmosphere. Then, solid/solid mechano-chemical synthesis of hydrogen storage materials with a particular attention to alanates and borohydrides is surveyed. Finally, more specialised techniques such as solid/liquid based methods are mentioned along with the common characteristics of mechanochemistry as a way of synthesizing hydrogen storage materials. Ó 2012 Elsevier Ltd. All rights reserved. Contents 1. Introduction Tuning of metal microstructures by mechanical milling BCC alloys Corresponding author. Tel.: ; fax: address: cuevas@icmpe.cnrs.fr (F. Cuevas) /$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved.

2 J. Huot et al. / Progress in Materials Science 58 (2013) Ti-based Mg-based BCC Amorphization Synthesis of hydrides by mechanically-induced solid/gas reactions Binary hydrides Magnesium hydride Titanium hydride Vanadium hydride Ternary hydrides ZrNi hydride TiNi hydride TiFe hydride LaNi 5 hydride TiV hydride Mg-based complex hydrides Mg 2 Fe hydride Mg 2 Co hydride Mg 2 Ni hydride Alanates Lithium alanates Sodium alanates Potassium alanates Mixed alkali alanates Alkali-earth alanates Synthesis of borohydrides by mechanical milling in diborane gas Synthesis of metal amides by mechanical milling in ammonia gas Synthesis of hydrides by mechanically-induced solid/solid and solid/liquid reactions Mechanochemical synthesis of metal borohydrides Synthesis of novel alane and metal alanates Novel quaternary hydrides Metal borohydride amides Metal alanate amides Solid liquid mechanically assisted synthesis Final remarks and conclusions Acknowledgments References Introduction Synthesis of innovative materials for energy conversion and storage has received increasing focus during the past decades due to the world s increasing energy demands and simultaneous needs for environmentally friendly energy technologies. Hydrogen is recognized as a possible renewable energy carrier, but its large-scale utilization is mainly hampered by unsatisfactory properties of known hydrogen storage materials. Hence, preparation and characterization of novel materials are receiving significant attention as reviewed elsewhere [1 4]. Traditionally, hydrogen storage materials, such as metallic or complex hydrides, were prepared by solvent-based synthesis methods or by direct solid gas hydrogenation reactions. However, during the past decade mechanochemical synthesis has become one of the most utilized preparation methods for this class of materials, and is still expected to hold a significant unexplored potential for development of novel approaches, e.g. for green chemistry including solvent free synthesis methods. In this work, recent progress within the experimental methods for preparation of hydrogen storage materials is surveyed. In the mid-eighties, several research groups initiated the use of mechanical activation methods for the synthesis of hydrides [5,6]. Mechanical milling (MM) of mixtures of elements under inert gas atmosphere was used to synthesize intermetallic compounds, which were subsequently exposed to hydro-

3 32 J. Huot et al. / Progress in Materials Science 58 (2013) gen in an external device to form hydrides. This two-step approach allowed modifying the microstructure of the alloy by milling to study its influence on the hydrogenation thermodynamics and kinetics [6 9]. Mechanical milling was very successful at improving hydrogenation kinetics via the synthesis of nanocrystalline materials and simultaneous incorporation of selective additives during the milling process [10 12]. Recently, it has been shown that Severe Plastic Deformation (SPD) techniques could be used for the synthesis and processing of metal alloys and their hydrides [13 26]. A few specific cases will be discussed in the forthcoming sections. A possible advantage of SPD techniques over conventional milling is easier scaling up to industrial level. However, some specific nanocrystalline structures and nanocomposites may only be synthesized through mechanical milling. In the early nineties, solid gas reaction facilitated by mechanical milling in reactive gases (nitrogen, oxygen and hydrogen) was investigated. This approach was initially designated Reactive Mechanical Milling (RMM) and used for preparation of hydrides in hydrogen atmosphere [27]. The experiments were performed in small-sized milling vials under moderate hydrogen pressures (below 2 MPa), often leading to an incomplete reaction between hydrogen and metals. The extent of the metal-hydrogen reaction was determined by ex situ XRD analysis of samples milled during a given period of time [28,29]. Modern devices for RMM synthesis are now equipped with pressure and temperature sensors that allow monitoring, e.g., hydrogen absorption during milling [30 32]. Hydrogenation reactions can be followed in situ as a function of milling time at working pressures up to ca. 15 MPa. Today, mechanochemical synthesis of metal hydrides using ball milling has grown to become one of the most frequently used methods. Typically, planetary ball mills are used, however other types such as rotational, vibratory or attritor mills are also operated [33]. The different types of mills differ in their milling efficiency and capacity and in some cases additional arrangements for cooling, heating, gas loading etc. can be applied. Typically a few grams of material and balls are placed in the planetary ball mill to give a ball-to-powder weight ratio of 10:1 50:1. This approach offers the advantage that the milling vial can be loaded, sealed and unloaded under inert conditions in a glove box, and, if equipped with valve connections, subsequently filled with reactive gas [30,31]. Thereby, the p,t phase space for mechanochemistry has expanded significantly. In some cases, especially for ductile materials, a process control agent (PCA) could be added to inhibit particle agglomeration [33]. The PCAs can be solids, liquids, or gases. A wide range of PCAs has been used in practice at a level of about 1 5 wt% of the total powder charge. The most common PCAs are stearic acid, hexane, methanol, ethanol, graphite and salts [33]. Several parameters can be varied for the ball-milling synthesis: milling speed, total milling time, vial and ball composition, powder-to-ball weight ratio, vial diameter, ball diameter and density, milling temperature, milling atmosphere and pressure of the selected gas. The latter two parameters require a special high-pressure vial. Most planetary mills only allow controlling the speed of the support disk. The speed of the planets, on which the milling vials are mounted, is usually fixed relatively to the speed of the main disk. However, for special mills, such as the Fritsch Vario-Planetary Mill Pulverisette 4, both the speed of the support disk and the planets can be varied freely [34]. Thereby, the trajectory of the balls within the vial may be controlled at least when the number of balls is low. Ideally the milling can be continuously changed from high-energy mode dominated by high-energy ball vial impacts to a grinding mode where the balls mainly follow the circumference of the vial [35]. The latter is also facilitated by a high number of balls in the vial. High-energy impacts tend to produce high mechanical pressure in the grain boundaries and in some cases make the high-pressure polymorph of the product. The grinding mode tends to produce more heat by friction and may in some cases lead to thermal decomposition of the product upon prolonged milling. Heating of the sample may be suppressed by using short periods of milling intervened by breaks where intrinsic heat produced in the grain boundaries can be dissipated and the sample can thermally equilibrate. Therefore, not only the total milling time is important for obtaining the desired compound, but breaks within the period of milling is in some cases crucial, which possibly also reduce agglomeration of the powder on the vial walls and balls [36 42]. Furthermore, the reactant mixture, balls and vial can be placed in a fridge or freezer prior to milling to lower the temperature further and/or the milling can be intervened by cooling of the vial. Milling at cryogenic conditions, i.e. at liquid nitrogen temperature (77 K), known as cryo-milling, has proven effective for preparation of some unstable metal hydrides [42].

4 J. Huot et al. / Progress in Materials Science 58 (2013) Within the past two decades, mechanochemistry has expanded widely both within the experimental methods and techniques but also within the variety of materials that can be prepared, e.g. binary and ternary metallic hydrides [43 47], or complex hydrides such as Mg-based transition metal hydrides [48 52], alanates [30,53 57], borohydrides [58], amides [59,60], and multi-component systems [61 64] including the broadly studied Reactive Hydride Composites (RHC) [65 68]. These topics are the focus for further discussion in this review. 2. Tuning of metal microstructures by mechanical milling The use of mechano-chemical methods for the synthesis and modification of hydrogen storage materials has generated an enormous amount of reports. In the last 10 years about a thousand papers have been published on the use of ball milling and mechanical alloying for this specific application. Therefore, this review focus on general aspects by discussion selected details and this section focus on the use of mechanochemical methods to tune the microstructure of metal hydride systems in order to improve their hydrogenation properties BCC alloys A body-centered cubic (BCC) structure is a coarse packing structure and has more interstitial space than face-centered cubic (FCC) and hexagonal close-packed (HCP) structures [69]. Thus, BCC alloys are more attractive candidates to be explored as possible interstitial hydrogen storage materials. Usually, BCC alloys are synthesized by arc melting or induction melting. However, for some alloys the desired composition is difficult to obtain by using these techniques because the constituting elements may have quite different melting temperatures. With mechanical alloying there is in principle no limitation on the nature and number of the raw elements used. For hydrogen storage applications one could distinguish two broad classes of BCC alloys: Ti-based and Mg-based. Each of these classes is discussed below Ti-based BCC alloys of systems Ti V Mn and Ti V Cr have been intensively studied for hydrogen storage [70 75]. This class of alloys may also catalyse hydrogen release and uptake in magnesium [76]. Fig. 1. X-ray powder diffraction pattern of arc-melted TiV 0.9 Mn 1.1 as a function of milling time (Cu Ka radiation) [77].

5 34 J. Huot et al. / Progress in Materials Science 58 (2013) Moderate hydrogen capacities as high as 3.6 wt% have been reported for Ti 25 V 40 Cr 35 alloy, which also possess prolific kinetic and thermodynamic properties. The effect of Severe Plastic Deformation (SPD) on BCC Ti 22Al 27Nb alloy has been investigated by Zhang et al. [25,26]. They showed that the first hydrogenation (activation) was much faster for the deformed alloy compared to the as-quenched sample. The deformed alloy also had 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. Cold rolling is a process by which a sheet metal or powder is introduced between rollers and then compressed and squeezed. 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 0.9 Mn 1.1 alloy [77]. 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. Fig. 1 shows the effect of milling on as-cast TiV 0.9 Mn 1.1. The presence of NaCl Bragg peaks is explained by the use of a small amount of this salt as an antisticking PCA. 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 structures was also observed for the system Fe Cu and may be due to an enhanced solubility due to the high dislocation density [78]. When milling was performed on the raw elements (Ti, V, and Mn), an identical result was obtained, i.e. formation of a nanocrystalline alloy composed of BCC and FCC phases [77]. The BCC alloys need activation, e.g. by cycling hydrogen release and uptake between p(h 2 ) = 5 MPa and vacuum at elevated temperature of 523 K. In Fig. 2 the hydrogen absorption and desorption isotherm (296 K) for arc-melted TiV 0.9 Mn 1.1 before and after 80 h of milling is presented. The maximum capacity of the as-melted alloy is 1.9 wt% at 7 MPa which corresponds to an H/M ratio of 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 none of them absorbs 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 the Ti V Cr system [79]. Fig. 2. Pressure composition temperature (PCT) curve, at 313 K, of arc-melted TiV 0.9 Mn 1.1 alloy before and after 80 h of milling [77].

6 J. Huot et al. / Progress in Materials Science 58 (2013) Fig. 3. TEM micrographs of TiV 1.6 Mn 0.4 after arc melting (top), after ball milled for 5 h (middle), and after 150 cold rolls (bottom). Micrographs on the left are bright field images and micrographs on the right are dark field images [82]. Singh et al. studied the effect of milling an arc-melted Ti 0.32 Cr 0.43 V 0.25 alloy [80]. As they used tungsten carbide balls, some contamination was observed after long milling time. Ball milling did not affect the crystal 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 sub-grain boundaries. Contamination from milling tools and microstructural changes caused an important decrease in the hydrogen storage capacity [80].

7 36 J. Huot et al. / Progress in Materials Science 58 (2013) Fig. 4. X-ray powder diffraction patterns of as-cast, milled 5 h and cold rolled 150 times TiV 1.6 Mn 0.4 alloy (Cu Ka radiation) [82]. Amira et al. compared the effect of ball milling and cold rolling for Ti Cr system [81]. 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 to ball-milled alloys despite different crystal structures. The alloy TiV 1.6 Mn 0.4 has been recently investigated by Couillaud et al. [82]. The effect of extended cold rolling as well as high energy ball milling was a reduction of crystalline size and lattice parameter but no change in the crystal structure. Fig. 3 shows TEM micrographs of TiV 1.6 Mn 0.4 alloy in the ascast, milled, and cold-rolled states. The dark field image shows that all samples are nanocrystalline. The bright field image of the coldrolled 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. Fig. 4 shows the X-ray diffraction patterns of as-cast, milled 5 h and cold rolled 150 times TiV 1.6 Mn 0.4 alloy. From these X-ray powder diffraction patterns it was determined that the crystallite sizes of as-cast, milled, and cold-rolled samples are respectively 17, 11, and 13 nm [82]. 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), which is a common feature of cold-rolled samples. Neither the ball-milled sample nor the cold-rolled samples 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 Mg-based BCC Recently, Mg-based BCC alloys have been explored in order to achieve higher gravimetric hydrogen storage capacity. In particular, Akiba s group has made an extensive study of the synthesis of Mg Ti [83,84], Mg Co [85 87], and Mg Ni [88,89] 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 [90,91], H 2 sources for fuel cells [84,92], switchable mirrors for smart solar collectors [93,94], and optical hydrogen detectors [95]. In the Mg Ti phase diagram equilibrium solid solubility of each metal in each other is less than 2 at.% and no intermetallic compound is found. Therefore, non-conventional synthesis methods based on melting or sintering can be used. Metastable single-phase Mg Ti thin films have been successfully synthesized over a large compositional range by

8 J. Huot et al. / Progress in Materials Science 58 (2013) Fig. 5. The powder X-ray diffraction pattern of a magnesium titanium Mg 50 Ti 50 mixture milled for 150 h (Cu Ka radiation) [84]. means of electron-beam and magnetron co-sputter deposition techniques [90,93,94,96,97]. 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 [83,91,98 101]. The synthesis of Mg Ti BCC alloys by mechanical alloying has been extensively studied by Asano et al. [83,84,92,100,101]. 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}h 12 10i while for titanium twinning deformation is more important [100]. 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 decrease its lattice parameter [100]. They found that by adding Li to Mg, the deformation of Mg was easier and the Ti crystallite size was reduced. This led to a decrease of synthesis time for BCC phase formation. In a subsequent study, they first synthesized a BCC Mg 50 Ti 50 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. Fig. 5 confirms that a BCC phase was obtained, and from the peaks width a crystallite size of 3 nm was determined [84]. A full hydrogenation at 423 K under 8 MPa of hydrogen and for 122 h resulted in the formation of Mg 42 Ti 58 H 177 FCC hydride phase and some MgH 2. 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 [83]. 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 [83,84]. If MgH 2 is used instead of Mg as the starting material then, after ball milling a 50MgH Ti mixture, the resulting compound is FCC Mg 33 Ti 50 H 94 plus some MgH 2 [92]. The importance of mechanical effect during milling is discussed in ref [101]. 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 an average diameter of 1 mm. These spherical particles are then crushed into spherical particles with a diameter of around 10 lm by introduction of cracks along the boundaries between Mg and Ti layers. Finally, the Mg 50 Ti 50 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 [101]. Recently, Çakmak et al. showed that mechanical milling of Mg 10 vol% Ti yields large Mg agglomerate, lm, with embedded Ti fragments of about 1 lm uniformly distributed within the agglomerates [102]. These Mg agglomerates are made of coherently diffracting volumes (crystallites) of small size. Crystallite size, as determined with X-ray diffraction analysis, can be as small as 26 nm after 30 h of milling.

9 38 J. Huot et al. / Progress in Materials Science 58 (2013) In an investigation of high-energy milling of 50Mg 50Ti mixture, Maweja et al. observed twinning in Ti-rich crystallites at intermediate milling time [103]. 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 straininduced martensitic transformation of the metastable x-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 [103]. 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 pre-milled before adding Pd then the alloying of Pd with pre-milled Mg 50 Ti 50 occurs very rapidly (few minutes) and is complete after 5 h of milling [104]. They also found that the crystalline structure of the Mg 50 Ti 50 alloy (BCC and HCP Mg Ti phase mixture) does not change significantly with the addition of Pd Amorphization In amorphization under mechanical driving forces, two categories of alloys could be defined [105]. The first category consists of intermetallic compounds induced to undergo polymorphic crystal-toamorphous transformation by deformation. In this process the introduction of defects by milling increases the free energy of the equilibrium alloy such that it goes above the free energy of the amorphous state. Thus, the amorphous phase becomes the lowest free energy state and the alloy becomes amorphous. The second category of amorphous alloys contains those formed by intermixing of individual elements that have a negative heat of mixing. In this case, deformation plays the role of enhancing such energy-lowering reactions through deformation-enhanced interdiffusion [105]. From a systematic study of Mg Ni system, Rojas et al. proposed the sequence of phase transformations during milling leading to amorphization as [106]: c-mg þ c-ni! nc-mg þ c-ni! amorphous þ nc-ni þ nc-mg! amorphous þ nc-ni þ nc-mg 2 Ni! nc-mg 2 Ni where c and nc denotes crystalline and nanocrystalline state, respectively. The first step shows the fact that grain refinement in nickel is slower than in magnesium. It has been demonstrated that the grain size attainable by milling depends on the crystal structure of the material being milled [107]. Usually, BCC materials tend to reach the smallest sizes, HCP materials somewhat larger grain sizes, and FCC materials tend to produce the largest grain sizes. Since the crystal structure of Mg is HCP and that of Ni is FCC, such a difference in the grain size after milling is expected. For Mg Ni system, amorphous phase could be prepared by ball milling in less than 10 h [108]. According to Varin et al., the presence of hard MgNi 2 phase helps to reduce crystallite size of Mg 2 Ni phase and thus facilitates amorphization [109,110]. For some compositions and milling parameters, a crystallization amorphous crystallization phenomenon could appear. One example of this was given by El-Eskandarany et al. for Co 75 Ti 25 [111]. A solid-state reaction took place during milling elemental Co and Ti powders and an amorphous phase of Co 75 Ti 25 was formed after 3 h. They showed that this amorphous phase crystallized into an ordered FCC-Co 3 Ti phase upon heating to 880 K. Further milling to 24 h also leads to crystallization and the formed phase was a metastable BCC-Co 3 Ti nanocrystalline phase. They attributed this transformation taking place in the ball mill to the inability of the formed amorphous phase to withstand the impact and shear forces that are generated by the milling media. When the milling time was further increased to 100 h, the crystalline phase was subjected to several points and lattice defects that raised the free energy from the stable BCC-Co 3 Ti phase to an amorphous less stable phase. In this case, the crystalline amorphous transformation which took place was similar to the mechanical grinding method in which the amorphization occurs by relaxing the short-range order without any compositional changes. Further milling leads to the formation of crystalline and/or amorphous phases depending on the milling time. Contamination from milling tools and temperature effect were ruled out as origin of this phenomenon [111]. Fig. 6 shows a schematic illustration of this crystallization amorphization crystallization process.

10 J. Huot et al. / Progress in Materials Science 58 (2013) Fig. 6. Schematic illustration of amorphous crystalline amorphous cyclic phase transformations that took place during ballmilling elemental powders of Co 75 Ti 25, using a rotation speed of 4.2 s 1 [111]. 3. Synthesis of hydrides by mechanically-induced solid/gas reactions Mechanical milling of metal powders under reactive gas, i.e. Reactive Mechanical Milling (RMM), is becoming a mature and powerful technique for the synthesis of metallic and complex hydrides. The mechanical treatment induces a chemical reaction between the solids and the gas. The synthesis of several metallic and complex hydrides by RMM is surveyed here. RMM under hydrogen gas allows for the synthesis of binary and ternary metal hydrides, Mg-based complex hydrides and alanates. More recently, this technique has been extended to other reactive gases such as diborane and ammonia for the synthesis of borohydrides and metal amides, respectively. Some particular phenomena such as ultra-fast hydride synthesis, reactive-milling induced amorphization, and multi-step reactions are reported. The obtained hydrides are typically nanocrystalline materials leading to fast kinetic for hydrogen release and uptake reactions useful for hydrogen storage applications. Significant progress on the understanding of RMM process has been provided by the in situ monitoring of the hydrogenation reaction during milling. In 2000, Dunlap et al. connected a ball-milling device to a large hydrogen reservoir by means of a rubber tube [112]. They could follow the hydrogen uptake as a function of milling time for several early transition metals (Ti, Zr, Hf, V, Nb and Ta). Experiments were conducted near atmospheric pressure (p(h 2 ) 0.1 MPa). A similar method was used by Bellosta et al. to monitor hydrogen release during ball-milling of sodium tetra-alanate with TiCl 3 additive [113]. Hydrogen release occurs due to titanium reduction to the zero-valent state on milling. A further improvement was reached by using telemetric systems instead of mechanical connections to in situ register both hydrogen pressure and vial temperature during milling [30,31]. The sensors were mounted on the lid of a stainless steel vial which was able to withstand high pressure

11 40 J. Huot et al. / Progress in Materials Science 58 (2013) (10 MPa). Under these conditions, the one-step direct synthesis of Ti-doped NaAlH 4 using NaH, Al and TiCl 3 as starting powders could be monitored. RMM is generally accomplished in tight stainless steel vials equipped with a connection valve for vacuuming and hydrogen filling. The vial is then placed in a milling device to promote the mechanochemical reaction leading to the hydride formation. For practical reasons regarding p T gauges attachment to vials, most of current milling devices used for RMM are planetary ball mills. Thus, this preparation technique is widely named as reactive ball milling. Nonetheless, also shaker, attritor, and vibration mills have been used successfully [51,112,114,115]. Today, compound synthesis can be anticipated from the in situ monitoring of the hydrogenation reaction and subsequently verified by ex situ crystallographic and chemical analyses. Furthermore, if thermodynamic parameters such as vial volume and gas pressure and temperature are accurately known, the quantity of absorbed hydrogen as a function of time can be reliably obtained. Zhang et al. have recently shown that hydrogen uptake can be determined with an accuracy of 95% [116]. In situ monitoring of changes in gas pressure is certainly a powerful tool for the study of hydride formation kinetics and reaction mechanisms on reactive milling Binary hydrides Though thermodynamically favourable, the formation of AH x binary hydrides by solid gas reaction between hydrogen gas and a metal (A, a metal with strong affinity for hydrogen, here stands for either alkaline earths (Mg) or early transition metals such as Ti and V) is very often hindered by kinetic barriers related to the presence of native oxide layers at the metal surface. Then, severe treatments at high temperature (typically above 700 K) and high pressure (several MPa) are needed in conventional gas-phase hydrogenation for activation. In the course of these treatments, oxygen at the surface might react with the bulk material leading to additional impurities. Such surface limitation can be overcome by RMM of pure metals in hydrogen atmosphere that allows achieving faster synthesis reactions under more moderate conditions. Synthesis conditions by RMM under hydrogen gas of representative binary hydrides are summarized in Table 1. Table 1 Representative binary and ternary metal hydrides synthesised by RMM under hydrogen gas. The employed device, reactants, initial hydrogen pressure, p(h 2 ), total milling time (tmt), milling speed (ms), ball-to-powder weight ratio (BTPWR) and ball diameter (Bd) are given. Compound Device Reactants p(h 2 ) (MPa) tmt (h) ms (rpm) BTPWR Bd (mm) Ref. MgH 2 Planetary Mg [117] MgH 2 Mg + graphite :1 [118] MgH 2 Fritsch P6 a Mg :1 10 [31] MgH 2 Fritsch P4 a Mg :1 12 [64] TiH 1.9 Planetary Ti [117] TiH 2 Fritsch P4 a Ti :1 12 [64] VH x Fritsch P5 V :1 7 [47] ZrNiH 3 Fritsch P5 ZrNi :1 10 [27] ZrH 2 + Ni Fritsch P5 Zr + Ni :1 10 [27] ZrH 2 + NiZr y H x Fritsch P5 Zr + Ni :1 10 [27] b-zrnih Fritsch P7 ZrNi :1 7 [28] c-zrnih 3 Fritsch P7 ZrNi :1 7 [28] TiNiH 3 Rod-mill Ti + Ni :1 10 [119] TiH 2 + Ni Ti + Ni :1 10 [120] TiH 2 + Fe Spex 8000 TiFe :1 6, 12 [114] a-lani 5 H 0.15 Fritsch P7 LaNi :1 7 [121] amph-lani 5 y H x Fritsch P7 LaNi :1 7 [121] BCC TiVH 0.9 Fritsch P5 TiV or Ti + V :1 10 [43] TiVH 2.8 Fritsch P5 TiV or Ti + V :1 10 [43] FCC TiVH 4.7 Fritsch P5 TiV or Ti + V :1 10 [43] Ti 0.20 V 0.78 Fe 0.02 H 2 Fritsch P4 * Ti 0.20 V 0.78 Fe :1 12 [122] a Pressure and temperature measured in situ in the Evico-magnetics vial.

12 J. Huot et al. / Progress in Materials Science 58 (2013) Magnesium hydride Magnesium hydride is classically prepared by reaction with hydrogen gas and Mg powder at temperatures around 700 K and hydrogen pressures in the range 7 8 MPa for several hours. However, even under these conditions, the presence of Mg is often detected by XRD as the reaction is not completed [123]. Indeed, a shell of magnesium hydride is reported to form at the surface of micrometer-sized magnesium grains, blocking further hydrogenation of the remaining metal core [124,125]. Consequently, the hydrogenation rate of bulk magnesium is slow. First attempts to form magnesium hydride by RMM were conducted by Chen and Williams using p(h 2 ) = 0.34 MPa and a vertical planetary mill [117]. Complete formation of MgH 2 hydride is reported to occur after long milling time (25 h). Later, similar experiments under MPa of hydrogen pressure were conducted by different groups [46,126,127]. Neither of them could attain hydride formation above 50 wt%, which was attributed to kinetic effects. This was finally overcome by performing RMM experiments at high temperature (573 K) with the addition of graphite to obtain complete hydrogenation within 1 h [118]. Graphite could act as a PCA to reduce particle agglomeration by cold-welding [33]. Doppiu et al. have however shown that fast MgH 2 formation can also be achieved near room temperature using high-pressure reactive ball milling [31]. The hydrogenation reaction could be monitored by in situ measurements of both pressure, p, and temperature, T, inside the vial during milling by using on-board sensors and radio transmitted data. Syntheses were done with pure Mg powders ball milled with a ball-to-powder weight ratio of 10:1 at 500 rpm and hydrogen pressures of 1, 4 and 9 MPa. From the data collection, it was first observed that the temperature of the vial increases up to 318 K mainly due to mechanical action. Mg absorbs hydrogen in less than 8 h for pressures larger than 4 MPa. Reaction rate was significantly slower for lower pressure (1 MPa). Moreover, a nucleation time, strongly dependant of the pressure is also reported; almost undetectable at 9 MPa, it reaches more than 2 h at 1 MPa. Further investigations by XRD at different milling times show the formation of the metastable orthorhombic c-phase along with the tetragonal b-mgh 2 one. The c-phase can also be achieved by ball milling of magnesium hydride [128]. With increasing milling time, the crystallite size decreases to finally stabilizing at 10 nm. Same amounts of hydride phases (>95 wt% for c + b) and identical crystallite sizes are obtained after 18 h of milling whatever the initial pressure though the rate of formation and the size reduction were faster for higher pressures. This was interpreted on the basis of two different factors. Higher pressures promote a more rapid formation of the hydride that is in turn known to exhibit higher plastic deformation. Then, for higher amounts of MgH 2, the mechanical action is more effective than for ductile Mg. However, it is worth noting that at long milling times, all samples reach the same chemical and microstructural states. Very similar results are also reported by Doppiu et al. who performed reactive milling of an elemental Mg 87 Ni 10 Al 3 powder mixture under hydrogen atmosphere [129]. Milling induces the synthesis of nanocrystalline MgH 2 at the first stage followed by the formation Mg 2 NiH 4 when a high degree of conversion of Mg in the hydride form was reached. A minimum value for the crystallite size of 8 nm P (MPa) (A) t (min) T (K) H/Mg (B) t (min) Fig. 7. Evolution of the pressure and temperature (A) during RMM of magnesium in hydrogen gas and the calculated H/Mg ratio in the solid state (B) as a function of time [64].

13 42 J. Huot et al. / Progress in Materials Science 58 (2013) was obtained. Small differences in the hydride stability were observed at different milling times. In spite of the oxidation of the sample, fast absorption desorption kinetics were obtained. A typical example of the evolution of the pressure, the temperature and the H concentration as a function of time during RMM of Mg is shown in Fig. 7. The starting material was coarse magnesium powder to reduce the amount of MgO that may develop at the grain surface of powder material. The milling vial was loaded with p(h 2 ) = 8 MPa and operated at 400 rpm. After initial heating due to ball friction, the pressure drop related to hydride formation is observed and the reaction is completed after 2 h. The final H/M value reaches 1.9, a value 5% smaller than expected for MgH 2. XRD analysis shows that the final product is made of 76 wt% of b-mgh 2, 21 wt% of c-mgh 2, and 3 wt% of MgO. The mean crystallite size for the hydride phases is close to 6 nm, in good agreement with previous results Titanium hydride The formation of titanium hydride by RMM with composition TiH l.9 was first reported by Chen and Williams in 1995 using a hydrogen filled container p(h 2 ) = 0.34 MPa for 67 h [117]. The hydrogenation reaction was completed in 5.5 h and the TiH 1.9 compound was stable during prolonged milling, with only a reduction of particle size being observed. Very similar results have been published by different groups [44 46]. Dunlap s group has extended this method to other early transition metals such as Zr, Hf, Ta, Nb and V [112]. Short reaction time was explained in terms of clean surface generation and severe reduction of diffusion path. Titanium powder is initially passivated by the presence of surface oxides. Upon milling, fresh and highly reactive surfaces are created promoting the formation of near-surface hydride precipitates. This causes hydrogen embrittlement of the metal, enhances its pulverisation and results in a shorter diffusion path for hydrogen absorption. The formation of hydride TiH 2 during RMM is shown in Fig. 8 [64]. Contrary to magnesium, no nucleation time is observed and the hydride formation proceeds readily and is completed after 10 min. XRD analysis confirms the formation of TiH 2 though small contamination with iron is observed, most probably due to the stainless steel vial abrasion during milling. The mean particle size determined from the diffraction peak widths is around 7 nm Vanadium hydride Orimo et al. reported on the preparation of nanostructured VH x prepared by mechanical milling under H 2 atmosphere [47]. Formation of the b 2 phase was observed after 5 min of milling at room temperature whereas conventional gas-phase hydrogenation would need activation treatments under 600 K and 3 MPa. The grain size of the b 2 phase decreases from 80 nm at 5 min milling time to 10 nm after 60 min. Additional milling time (up to 300 min) does not lead to further decrease of the grain size nor the formation of an amorphous state as the b 2 phase remains in crystalline state. From the relationship between hydrogen concentration and unit cell, the hydrogen concentration P (MPa) (A) T (K) t (min) t (min) Fig. 8. Evolution of the pressure, the temperature (A) and the H concentration (B) as a function of time during RMM of titanium in hydrogen gas. H/Ti (B)

14 J. Huot et al. / Progress in Materials Science 58 (2013) was determined as a function of the grain size. It decreases from 0.82 H/M for 80 nm down to 0.72 H/ M for 10 nm indicating a modification of the b 2 c phase boundary in the V H system for nanometer grain sizes. Lower concentration, nearly independent of the grain size and higher diffusivity of hydrogen are also reported in the intergrain domain Ternary hydrides Ternary metal hydrides of general composition ABH x can be easily synthesized by RMM providing that the hydrides are stable under the pressure and temperature milling conditions. In this formulation, A stands for a metal with strong affinity for hydrogen (an early transition or rare-earth metal) and B stands for a metal with weak hydrogen affinity (a late transition metal). Well-known hydrogen storage intermetallic compounds such as ZrNi, TiFe, TiNi and LaNi 5 have been used in RMM experiments (see Table 1). As a general behaviour, two effects are observed on prolonged milling: formation of amorphous ABH x hydrides and compound disproportionation into AH x + B species. Compound amorphization is driven by the large negative heat of mixing between A and B elements while its disproportionation is favoured by the different affinity between both elements for hydrogen ZrNi hydride RMM experiments in the Zr Ni system have been first reported by Aoki et al. [27]. They performed RMM (p(h 2 ) = 2 MPa) of both arc-melted ZrNi alloys and equiatomic Zr and Ni powder mixtures. In the first case, ZrNiH 3 hydride is formed after 3 h of milling. On prolonged milling (over 100 h), the crystallite size of the hydride decreases without apparent amorphization. The lack of amorphous phases is attributed to the difficulty to introduce defects in ZrNiH 3 hydride because of its brittleness. For RMM of elemental powders, a mixture of ZrH 2 and elemental Ni is formed at short milling time (<3 h). Further reaction over 100 h leads to the coexistence of ZrH 2 and amorphous Zr-poor NiZr 1 y H x phase. Orimo et al. have studied the effect of hydrogen pressure during RMM of ZrNi compound within the range 0 1 MPa [28]. b-zrnih and c-zrnih 3 hydrides are formed within 5 min of milling. Phase abundance depends on hydrogen pressure. Formation of the most stable b-hydride is observed at 0.1 MPa, whereas that of the less stable c-hydride occurs at 1 MPa. At intermediate pressures, 0.3 MPa, both phases are detected. As observed by Aoki, prolonged milling over 80 h results in the formation of ZrH 2 and amorphous Zr-poor NiZr 1 y H x phase. Such decomposition reaction seems to be delayed with increasing hydrogen pressure TiNi hydride RMM experiments (p(h 2 ) = 0.1 MPa) on equiatomic Ti and Ni powder mixture have been conducted in a rod-milling device [119]. Within the first 3 h of milling metallic Ti transforms to TiH 2 and metallic FCC Ni remains unreacted. Further milling up to 200 h leads to the gradual formation of a nanocrystalline (10 nm) single-phase FCC compound. The compound is described as an FCC TiNiH 3 solid solution, though details on the hydrogen content determination are not provided. This result is rather striking since TiNi compound only absorbs 1.4 H/f.u. under normal conditions of pressure and temperature [130]. In fact, later experiments at higher hydrogen pressure (1.1 MPa) failed to get the solid solution TiNiH 3 phase [120]. Instead, formation of poorly crystallized TiH 2 and Ni phases on milling for 40 h is reported TiFe hydride Chiang et al. have performed RMM experiments (p(h 2 ) = 0.5 MPa) in TiFe compound [114]. In situ manometric measurements reveal that a total hydrogen uptake of 1.6 H/f.u. occurs within 7 h of milling forming TiFeH 1.6. Ex situ XRD analysis reveals that the ternary hydride decomposes to TiH 2 and Fe LaNi 5 hydride Fujii et al. have performed RMM experiments (p(h 2 ) = 1 MPa) on LaNi 5 alloy and observe formation of solid solution a-lani 5 H 0.15 within the first 5 min of milling [121]. Formation of hydride b-lani 5 H 6 phase is not detected, which may indicate that the absorption plateau pressure of this hydride is above

15 44 J. Huot et al. / Progress in Materials Science 58 (2013) MPa at the temperature attained on milling. At longer milling times, 5 min < t < 3 h, the a-phase coexist with an amorphous phase. This phase forms faster in the presence of hydrogen than in ballmilling experiments performed under inert gas. The a-phase decomposes into Ni and amorphous Ni-poor LaNi 5 y H x phase upon prolonged milling, 3 h < t < 10 h. Subsequent thermodynamic measurements show a significant reduction of total and reversible hydrogen storage capacity for long-time milled as compared to pristine LaNi 5 compound. This is attributed to a lower hydrogenation capacity of both poorly crystallized inter-grain and amorphous regions as compared to the microcrystalline state. This concurs with the facts that nanocrystalline systems exhibit lower capacity than microcrystalline ones and that hydrogen binding energies expands over a wide energy range in amorphous systems [131,132] TiV hydride One should notice that the Ti V system differs from previous ones as concerns the affinity of constituting elements towards hydrogen. Both elements are A-type and exhibit comparable affinity for hydrogen which, in principle, precludes alloy disproportionation by hydrogenation. Furthermore, this system exhibits small heat of mixing so that alloy amorphization is expected to be difficult. RMM (p(h 2 ) = 0.2,0.4 and 2 MPa) of either equiatomic Ti and V powder mixtures or BCC TiV alloy have been conducted by Aoki et al. [43]. At long milling time (100 h), phase constitution of milled products does not depend on the nature of the initial powder. Hydrogen pressure plays, however, a major role. At low (0.2 MPa) and high (1 MPa) pressures, BCC TiVH 0.9 solid solution and FCC TiVH 4.7 hydride are formed, respectively. The hydrogen content of FCC TiVH 4.7 hydride is probably overestimated since maximum hydrogen uptake of both Ti and V is 2 H/M. Nevertheless, both BCC and FCC hydrogenated phases are crystalline. In contrast, at intermediate pressure (0.4 MPa), amorphous TiVH 2.8 phase is obtained. The formation mechanism of this phase depends on the initial reactants. For Ti + V powder mixture, the amorphous phase is formed by reaction between TiH 2 and V, whereas for the arc-melted alloy it results from gradual amorphization of BCC TiVH 2.8 phase on milling. Strikingly, for both cases, the amorphization reaction occurs without changing the hydrogen content. RMM experiments on a BCC Ti V Fe alloy of composition Ti 0.20 V 0.78 Fe 0.02 have been carried out in a device equipped with p and T sensors at rotation speed of 400 rpm and p(h 2 ) = 8 MPa. The hydrogenation curve is shown in Fig. 9 [122]. The alloy absorbs 2 H/f.u. in only ten minutes indicating the formation of a stoichiometric (Ti,V,Fe)H 2 hydride. Further milling produces hydrogen desorption from the milled sample. Its hydrogen content decreases to 1.55 H/f.u. after 8 h of milling. XRD diffraction analysis (Fig. 10) shows that the RMM alloy consists of a mixture of FCC VH 2 -type hydride and amorphous H/f.u t (min) Fig. 9. Time-evolution of the H concentration in BCC Ti 0.20 V 0.78 Fe 0.02 alloy during RMM in hydrogen gas.

16 J. Huot et al. / Progress in Materials Science 58 (2013) I (counts) amorphous phase θ Fig. 10. Rietveld analysis of Ti 0.20 V 0.78 Fe 0.02 alloy after RMM for 480 min. Observed (dots), calculated (top line) and difference curves (bottom line) are shown. Vertical bars ( ) correspond to Bragg positions (Cu K a 1,2 ) for FCC VH 2 -type hydride. Large dots stand for the contribution of an amorphous phase to the calculated pattern. phase. The amorphous phase formation accounts for the spontaneous hydrogen desorption on milling since it stores less hydrogen Mg-based complex hydrides Mg-based Mg 2 TH x ternary hydrides (T = Fe, Co and Ni transition metals) are attractive hydrogen storage materials due to their high specific (5.5, 4.5 and 3.6 wt%) and volumetric (150, 125 and 97 g/l) hydrogen contents for Mg 2 FeH 6,Mg 2 CoH 5 and Mg 2 NiH 4, respectively [138]. The synthesis of these hydrides is problematic due to the great difference in vapour pressure and melting point between Mg and T and the lack of stable Mg 2 Fe and Mg 2 Co intermetallic compounds in their respective binary phase diagrams. From these facts, synthesis of Mg 2 TH x ternary hydrides was classically achieved by sintering methods from elemental powder mixtures. Temperatures and hydrogen pressures as high as 750 K and 9 MPa, respectively, and reaction time of several days are required [139]. Synthesis conditions of Mg 2 TH x ternary hydrides by RMM under hydrogen gas are summarized in Table Mg 2 Fe hydride The synthesis of Mg 2 FeH 6 by RMM (p(h 2 ) = 1 MPa for 20 h) of Mg and Fe powder was first attempted in 1997 [133]. Mg 2 FeH 6 hydride was not obtained but Mg powder got hydrogenated to form intimate MgH 2 Fe mixture. The failure to form the ternary hydride could be due to milling conditions not being sufficiently efficient (ball-to-powder weight ratio of 4:1). It was later discovered that the desired hydride can be obtained in two different ways. The first simply consists in a sintering treatment of the reactive milled product MgH 2 Fe for one day at 625 K under 5 MPa of hydrogen. The second, more complex, was reported in a subsequent paper [134]. Mechanical milling of MgH 2 and Fe powders in molar ratio 2:1 under argon atmosphere was performed for 60 h in a high-energetic shaker mill with ball-to-powder weight ratio of 10:1. The mechanical energy provided under these milling conditions was high enough to promote Mg 2 FeH 6 formation without subsequent sintering. Much probably, the following solid-state reaction takes place: 3MgH 2 ðsþþfeðsþ!mg 2 FeH 6 ðsþþmgðsþ ð1þ The formation of the ternary compound is driven by the fact that the Mg 2 FeH 6 phase is more stable than MgH 2 [140]. In situ SR-PXD patterns measured for a ball milled sample of MgH 2 Fe (2:1) reveal formation of Mg 2 FeH 6 at 673 K at p(h 2 ) = 10 MPa [141].

17 46 J. Huot et al. / Progress in Materials Science 58 (2013) Table 2 Mg-based complex hydrides synthesised by RMM under hydrogen gas. The employed device, reactants, initial hydrogen pressure, p(h 2 ), total milling time (tmt), milling speed (ms), ball-to-powder weight ratio (BTPWR), ball diameter (Bd), reaction yield and formed side products are given. Compound Device Reactants p(h 2 ) (MPa) tmt (h) ms (rpm) BTPWR Bd Yield (mm) (wt%) Side products Mg 2 FeH 6 Fritsch P5 2Mg + Fe : MgH 2,Fe [133] Mg 2 FeH 6 Spex MgH 2 + Fe : Mg, MgO, [134] Fe Mg 2 FeH 6 Uni-Ball-Mill II 2Mg + Fe :1 28 MgO, Fe [49] Mg 2 FeH 6 Szegvari attritor 2Mg + Fe : Fe [115] Mg 2 FeH 6 Retsch Mg + Fe s 1 16: b Fe [51] vibrating mill Mg 2 FeH 6 Fritsch P4 a 2Mg + Fe : MgO, Fe [116] Mg 2 CoH 5 Kurimoto 2MgH 2 + Co : [48] planetary mill Mg 2 CoH 5 Uni-Ball-Mill II 2Mg + Co :1 50 Co [135] Mg 2 CoH 5 Fritsch P4 a 2Mg + Co : MgO [116] Mg 2 NiH 4 Fritsch P5 2Mg + Ni : Mg, [29] MgH 2,Ni Mg 2 NiH 1.8 Fritsch P7 Mg 2 Ni : [136] Mg 2 NiH 4 Kurimoto Mg 2 Ni : b amph- [137] planetary mill MgNi Mg 2 NiH 4 Retsch Mg + Ni s 1 16: [51] vibrating mill Mg 2 NiH 4 Fritsch P4 a 2Mg + Ni : MgO [116] Mg 2 (FeH 6 ) 0.5 (CoH 5 ) 0.5 Fritsch P6 a 4Mg + Fe + Co : b FeCo [52] a b Pressure and temperature measured in situ in the Evico-magnetics vial. Estimated values. Ref. In 2002, direct though incomplete synthesis of Mg 2 FeH 6 by RMM of elemental powders was simultaneously reported [49,115]. Gennari et al. used a Uni-Ball-Mill II device under 0.5 MPa with hydrogen refilling every 5 h to maintain constant hydrogen pressure in the vial [49]. Mg 2 FeH 6 formation with a yield of 28 wt% was achieved after 60 h of milling. The synthesis was reported to occur in two steps. MgH 2 is formed during the first 40 h by mechanically activated solid gas reaction followed by the solid-state reaction between MgH 2 and Fe at longer milling times. Raman et al. [115] used a Szegvari attritor device under 1 MPa of hydrogen. The ternary hydride started forming after 14 h of milling and a maximum yield of 63 wt% was achieved at 20 h of milling as determined from XRD analysis. Reaction yield is probably overestimated since a high quantity of Fe (37 wt%) was identified as the unique secondary phase, which is not possible from mass-balance considerations. In fact, significant residuals in the Rietveld analysis likely related to MgO phase can be observed, which explains the presence of unreacted Fe (similar effects have been later observed [50]). Crystallite sizes of 12 and 18 nm are reported for Mg 2 FeH 6 and Fe phases, respectively. Formation on MgH 2 as intermediate phase was not detected. Prolonged milling to 30 h is reported to lead to amorphization of the ternary hydride. Milling under hydrogen of 2MgH 2 + Fe and 2Mg + Fe powder mixtures have also been compared [142]. It was found that a faster reaction and higher yield is achieved for elemental powders as compared to 2MgH 2 + Fe. The differences were attributed to the dissimilar mechanical properties and microstructures of the mixtures. The 2Mg + Fe mixture behaves as a ductile ductile pair that results in a higher contact surface between Mg and Fe, and a better intermixing and size reduction. On the contrary, the 2MgH 2 + Fe mixture performs as a ductile brittle combination, with less contact area between the reactants and hence lower yield and longer synthesis time. In 2008, the direct synthesis of Mg 2 FeH 6 by RMM of elemental powders was monitored in situ by manometric means by Baum et al. [51] RMM experiments were performed in a horizontal vibrating mill operated at 32 s 1 under a hydrogen pressure of 0.3 MPa. In spite of using mild milling conditions (one unique ball and ball-to-powder weight ratio of 16:1), the reaction was completed after only 8 h of milling time. The reaction yield is not specified, but judging from the total hydrogen uptake and XRD

18 J. Huot et al. / Progress in Materials Science 58 (2013) data it should be around 90 wt%. This record could be related to the fact that enough hydrogen pressure was kept in the system on milling. The system was refilled with hydrogen when the pressure decreased below 0.27 MPa. Moreover, from in situ hydrogen uptake curves the authors could infer, in agreement with Gennari et al. a two-step process with formation of MgH 2 during the first 2 h of milling followed by the formation of Mg 2 FeH 6 from MgH 2 and Fe at longer milling times [49]. These results have later been confirmed by Deledda and Hauback by in situ measurements during RMM in an Evicomagnetics vial operated at 5 MPa of hydrogen pressure [52]. The latter authors observed, however, that the first step exceeded the hydrogen capacity of MgH 2 and proposed additional hydrogen uptake at Mg/Fe interfaces. Such additional capacity is doubtful since the authors used the ideal gas law, which is not valid at the imposed pressures, to estimate hydrogen absorption. The reaction path during RMM synthesis (p(h 2 ) = 7.5 MPa for 12 h) of Mg 2 FeH 6 has been recently studied in an Evico-magnetics vial [116]. The evolution of the hydrogen uptake as a function of milling time is shown in Fig. 11. The result for a similar experiment using only Mg powder is shown in the same figure for comparison. Hydrogen absorption by 2Mg + Fe powder mixture occurs in two steps Mg+Fe Hydrogen uptake (H/f.u.) H/f.u. 2Mg 2 nd step st step t = 50 min Milling time (min) Fig. 11. Time-evolution of the H concentration in solid-state during RMM of Mg powder and 2Mg + Fe powder mixture. Fig. 12. XRD patterns and phase identification of RMM 2Mg + Fe powder mixture after the first (50 min) and the second (720 min) reaction step (Cu Ka radiation).

19 48 J. Huot et al. / Progress in Materials Science 58 (2013) The first step, with t < 50 min, corresponds to the formation of MgH 2 hydride as demonstrated by both the equivalent amount of absorbed hydrogen in the experiment conducted with elemental Mg and ex situ XRD measurements (Fig. 12). The second step corresponds to the reaction between MgH 2 and Fe to form Mg 2 FeH 6. The XRD pattern for 12 h milled product (Fig. 12) shows the presence of three phases Mg 2 FeH 6 (77 wt%), Fe (12 wt%) and MgO (11 wt%) with crystallite sizes of 8, 12 and 4 nm, respectively. In contrast to previous reports no amorphization is observed on prolonged milling [115,143]. MgO contamination is attributed to undesired surface oxidation of fine Mg powder (36 lm) in glove-box and accounts for unreacted Fe residual [144]. According to these results, the two-step reaction path for Mg 2 FeH 6 synthesis could be described as: 2MgðsÞþFeðsÞþ3H 2 ðgþ!2mgh 2 ðsþþfeðsþþh 2 ðgþ!mg 2 FeH 6 ðsþ It is worth noting that reaction kinetics for the first step, i.e. MgH 2 formation, is faster for the 2Mg + Fe mixture than for pure Mg powder. This striking result may be related either to catalytic effects for hydrogen dissociation at the Fe surface or to nucleation phenomena at Mg/Fe interfaces [145]. The reaction path given by Eq. (2) concurs with recent reports on hydrogen absorption by classical solid gas reaction in nanosized Mg + Fe mixtures [146,147]. Based on DFT calculations, it has been proposed that the reaction between iron and magnesium hydride may occur through the formation of a (MgFe)H 2 solid solution which becomes unstable with increasing Fe content with respect to Mg 2 FeH 6 [148]. As we saw in the previous sections, ball milling has been extensively used to synthesize Mg 2 FeH 6. In the case of Severe Plastic Deformation SPD techniques, investigation has only started recently and the literature is much less abundant. Lima et al. observed substantial improvement in the hydrogen sorption kinetics of a Mg Fe powder mixture processed by high pressure torsion HPT [16]. The authors noted that hydrogenation and dehydrogenation of the processed samples did not change the preferential orientation (002) of the Mg phase, i.e. the material retained the microstructure imposed by HPT. In a subsequent investigation, the same authors used a combination of ball milling and extrusion to synthesize Mg 2 FeH 6 [149]. Their results indicate that the iron in the 2Mg Fe mixture produced a beneficial pinning effect on the Mg grains by hindering grain coarsening even after annealing treatments. The desorption kinetics of samples processed by high pressure torsion (HPT) was faster than that of extruded samples, probably due to bulk diffusion limitations [149] Mg 2 Co hydride Similarly to Mg 2 FeH 6, first attempts to produce Mg 2 CoH 5 by RMM (p(h 2 ) = 1 MPa for 20 h) of Mg and Co powder were unsuccessful. Instead, a mixture of MgH 2 and Co phases was obtained [133]. Subsequent sintering treatment allowed synthesizing Mg 2 CoH 5 hydride though with a lower yield (26 wt%) than for Mg 2 FeH 6 (65 wt%). The synthesis of Mg 2 CoH 5 hydride by RMM (p(h 2 ) = 0.1 MPa for 10 h) of MgH 2 and Co powders in molar ration 2:1 under hydrogen pressure was first reported by Chen et al. [48]. Ex situ XRD measurements revealed that Mg 2 CoH 5 phase started forming at 1 h milling and became the major phase after 10 h milling. Later, the synthesis could be also achieved using Mg and Co powders as reactants under 0.5 MPa of hydrogen [135]. The powders were previously milled under Ar atmosphere in the same system for 200 h. It was supposed that intimate contact and homogeneity between Mg and Co phases is essential to reach hydride formation as occurring for sintering methods [150]. The Mg 2 CoH 5 phase starts forming after 40 h of milling and a yield of 50 wt% was achieved at 90 h. The formation of an intermediate MgH 2 phase occurs from 10 h of milling. The two-step reaction has been also observed by Baum et al. by in situ monitoring of hydrogen uptake during RMM experiments [51]. The reaction path during RMM synthesis (p(h 2 ) = 7.5 MPa for 12 h) of Mg 2 CoH 5 from Fe and Co powders in molar ratio 2:1 has been recently studied in detail [116]. The evolution of the hydrogen uptake as a function of milling time is shown in Fig. 13. Hydrogen absorption occurs in two steps which, after the analysis of XRD data (Fig. 14), correspond to the following reactions: 2MgðsÞþCoðsÞþ5=2H 2 ðgþ!2mgh 2 ðsþþcoðsþþ1=2h 2 ðgþ!mg 2 CoH 5 ðsþ MgH 2 hydride is formed as an intermediate phase for t < 50 min. The reaction is again faster as compared to Mg milled alone. The second step corresponds to the reaction between MgH 2 and Co to form ð2þ ð3þ

20 J. Huot et al. / Progress in Materials Science 58 (2013) Mg+Co Hydrogen uptake (H/f.u.) H/f.u. 2 nd step 1 st step t = 50 min 2Mg Milling time (min) Fig. 13. Time-evolution of the H concentration in solid-state during RMM of Mg powder and 2Mg + Co powder mixture. Fig. 14. XRD patterns and phase identification of RMM 2Mg + Co powder mixture after the first (50 min) and the second (720 min) reaction step (Cu Ka radiation). Mg 2 CoH 5. The hydrogen uptake corresponding to this reaction is lower but faster than for Mg 2 FeH 6 formation (Fig. 11). The XRD pattern for 12 h milled product (Fig. 14) shows the formation of nanocrystalline (8 nm) Mg 2 CoH 5 phase without significant amorphous contribution Mg 2 Ni hydride Since Mg 2 Ni compound exists as stable phase in the binary Mg Ni phase diagram, the synthesis of Mg 2 NiH 4 ternary hydride by RMM can be attempted either using 2Mg + Ni elemental mixture or Mg 2 Ni powders as initial reactants. This synthesis was first tried by RMM (p(h 2 ) = 0.5 MPa for 22 h) of 2Mg + Ni powders [29]. Some MgH 2 phase was formed from 2 h of milling but no ternary hydride could be detected. Either milling energy (ball-to-powder weight ratio was 4:1 and milling speed 325 rpm) or hydrogen supply was insufficient to promote hydride formation. Orimo et al. later investigated hydrogen absorption in Mg 2 Ni during reactive milling (p(h 2 ) = 1 MPa for 80 h) using stronger energetic conditions: rotation speed of 400 rpm and ball-to-powder weight