Mechanochemical synthesis of hydrogen storage materials

Transcription

Progress in Materials Science 58 (2013) 30–75
Contents lists available at SciVerse ScienceDirect
Progress in Materials Science
journal homepage: www.elsevier.com/locate/pmatsci
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
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, 94320 Thiais Cedex, France
b
a r t i c l e
i n f o
Article history:
Received 13 April 2012
Accepted 9 July 2012
Available online 27 July 2012
a b s t r a c t
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.
2.
Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31
Tuning of metal microstructures by mechanical milling . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33
2.1.
BCC alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33
⇑ Corresponding author. Tel.: +33 1 49 78 12 25; fax: +33 1 49 78 12 03.
E-mail address: [email protected] (F. Cuevas).
0079-6425/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved.
http://dx.doi.org/10.1016/j.pmatsci.2012.07.001
J. Huot et al. / Progress in Materials Science 58 (2013) 30–75
3.
4.
5.
2.2.
Ti-based . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
2.3.
Mg-based BCC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
2.4.
Amorphization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
Synthesis of hydrides by mechanically-induced solid/gas reactions. . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.1.
Binary hydrides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.1.1.
Magnesium hydride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.1.2.
Titanium hydride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.1.3.
Vanadium hydride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.2.
Ternary hydrides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.2.1.
ZrNi hydride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.2.2.
TiNi hydride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.2.3.
TiFe hydride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.2.4.
LaNi5 hydride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.2.5.
TiV hydride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.3.
Mg-based complex hydrides. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.3.1.
Mg2Fe hydride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.3.2.
Mg2Co hydride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.3.3.
Mg2Ni hydride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.4.
Alanates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.4.1.
Lithium alanates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.4.2.
Sodium alanates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.4.3.
Potassium alanates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.4.4.
Mixed alkali alanates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.4.5.
Alkali-earth alanates. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.5.
Synthesis of borohydrides by mechanical milling in diborane gas . . . . . . . . . . . . . . . . . . . . . . . .
3.6.
Synthesis of metal amides by mechanical milling in ammonia gas . . . . . . . . . . . . . . . . . . . . . . . .
Synthesis of hydrides by mechanically-induced solid/solid and solid/liquid reactions . . . . . . . . . . . . . .
4.1.
Mechanochemical synthesis of metal borohydrides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.2.
Synthesis of novel alane and metal alanates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.3.
Novel quaternary hydrides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.3.1.
Metal borohydride amides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.3.2.
Metal alanate amides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.4.
Solid–liquid mechanically assisted synthesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
Final remarks and conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
Acknowledgments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
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1. Introduction
Synthesis of innovative materials for energy conversion and storage has received increasing focus
during the past decades due to the world́s 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-
32
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].
33
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.
2.1. 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.
2.2. 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 TiV0.9Mn1.1 as a function of milling time (Cu Ka radiation) [77].
34
Moderate hydrogen capacities as high as 3.6 wt% have been reported for Ti25V40Cr35 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 TiV0.9Mn1.1 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 TiV0.9Mn1.1.
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(H2) = 5 MPa
and vacuum at elevated temperature of 523 K. In Fig. 2 the hydrogen absorption and desorption isotherm (296 K) for arc-melted TiV0.9Mn1.1 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 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 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 TiV0.9Mn1.1 alloy before and after 80 h of milling
[77].
35
Fig. 3. TEM micrographs of TiV1.6Mn0.4 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 Ti0.32Cr0.43V0.25 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].
36
Fig. 4. X-ray powder diffraction patterns of as-cast, milled 5 h and cold rolled 150 times TiV1.6Mn0.4 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 TiCrx (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 TiV1.6Mn0.4 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 TiV1.6Mn0.4 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 TiV1.6Mn0.4 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.
2.3. 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], H2 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
37
Fig. 5. The powder X-ray diffraction pattern of a magnesium titanium Mg50Ti50 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 {0 0 0 1}h1 2 1 0i 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 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. 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
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 [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 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 [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 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 [101].
Recently, Çakmak et al. showed that mechanical milling of Mg–10 vol% Ti yields large Mg agglomerate, 90–100 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.
38
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 Mg50Ti50 occurs very rapidly (few minutes) and is complete
after 5 h of milling [104]. 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.
2.4. 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-Mg2 Ni ! nc-Mg2 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 MgNi2 phase helps to reduce crystallite size of Mg2Ni
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 Co75Ti25 [111].
A solid-state reaction took place during milling elemental Co and Ti powders and an amorphous phase
of Co75Ti25 was formed after 3 h. They showed that this amorphous phase crystallized into an ordered
FCC-Co3Ti phase upon heating to 880 K. Further milling to 24 h also leads to crystallization and the
formed phase was a metastable BCC-Co3Ti 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-Co3Ti 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.
39
Fig. 6. Schematic illustration of amorphous–crystalline–amorphous cyclic phase transformations that took place during ballmilling elemental powders of Co75Ti25, using a rotation speed of 4.2 s1 [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(H2) 0.1 MPa). A similar method was used by
Bellosta et al. to monitor hydrogen release during ball-milling of sodium tetra-alanate with TiCl3 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
40
(10 MPa). Under these conditions, the one-step direct synthesis of Ti-doped NaAlH4 using NaH, Al
and TiCl3 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.
3.1. Binary hydrides
Though thermodynamically favourable, the formation of AHx 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(H2), total milling time (tmt), milling speed (ms), ball-to-powder weight ratio (BTPWR) and ball diameter (Bd)
are given.
a
Compound
Device
Reactants
p(H2) (MPa)
tmt (h)
MgH2
MgH2
MgH2
MgH2
TiH1.9
TiH2
VHx
ZrNiH3
ZrH2 + Ni
ZrH2 + NiZryHx
b-ZrNiH
c-ZrNiH3
TiNiH3
TiH2 + Ni
TiH2 + Fe
a-LaNi5H0.15
amph-LaNi5yHx
BCC TiVH0.9
TiVH2.8
FCC TiVH4.7
Ti0.20V0.78Fe0.02H2
Planetary
Mg
Mg + graphite
Mg
Mg
Ti
Ti
V
ZrNi
Zr + Ni
Zr + Ni
ZrNi
ZrNi
Ti + Ni
Ti + Ni
TiFe
LaNi5
LaNi5
TiV or Ti + V
TiV or Ti + V
TiV or Ti + V
Ti0.20V0.78Fe0.02
0.34
0.4
1–9
8
0.34
8
1
2
2
2
0.1
1
0.1
1.1
0.5
1
1
0.2
0.4
1
8
25
1
8
2
5.5
0.16
0.17
3
3
100
0.08
0.08
200
40
7
0.08
10
100
100
100
0.17
Fritsch P6a
Fritsch P4a
Planetary
Fritsch P4a
Fritsch P5
Fritsch P5
Fritsch P5
Fritsch P5
Fritsch P7
Fritsch P7
Rod-mill
Spex 8000
Fritsch P7
Fritsch P7
Fritsch P5
Fritsch P5
Fritsch P5
Fritsch P4*
Pressure and temperature measured in situ in the Evico-magnetics vial.
ms (rpm)
500
400
400
400
400
400
250
400
400
400
BTPWR
10:1
10:1
60:1
60:1
30:1
30:1
30:1
30:1
30:1
30:1
30:1
10:1
8:1
30:1
30:1
20:1
20:1
20:1
100:1
Bd (mm)
Ref.
12
[117]
[118]
[31]
[64]
[117]
[64]
[47]
[27]
[27]
[27]
[28]
[28]
[119]
[120]
[114]
[121]
[121]
[43]
[43]
[43]
[122]
10
12
12
12
7
10
10
10
7
7
10
10
6, 12
7
7
10
10
10
12
41
3.1.1. 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(H2) = 0.34 MPa and a vertical planetary mill [117]. Complete formation of MgH2 hydride is reported
to occur after long milling time (25 h). Later, similar experiments under 0.5–1 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 MgH2 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-MgH2 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 MgH2, 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 Mg87Ni10Al3 powder mixture under hydrogen atmosphere [129]. Milling induces the synthesis
of nanocrystalline MgH2 at the first stage followed by the formation Mg2NiH4 when a high degree of
conversion of Mg in the hydride form was reached. A minimum value for the crystallite size of 8 nm
320
9.0
2.0
8.5
316
312
7.0
308
6.5
6.0
304
5.5
(A)
5.0
0
60
120
t (min)
300
180
1.5
H/Mg
7.5
T (K)
P (MPa)
8.0
1.0
0.5
(B)
0.0
0
60
120
180
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].
42
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(H2) = 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 MgH2. XRD analysis shows that the final product is made of 76 wt% of b-MgH2, 21 wt% of c-MgH2, and 3 wt% of MgO.
The mean crystallite size for the hydride phases is close to 6 nm, in good agreement with previous
results.
3.1.2. Titanium hydride
The formation of titanium hydride by RMM with composition TiHl.9 was first reported by Chen and
Williams in 1995 using a hydrogen filled container p(H2) = 0.34 MPa for 67 h [117]. The hydrogenation
reaction was completed in 5.5 h and the TiH1.9 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 TiH2 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 TiH2 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.
3.1.3. Vanadium hydride
Orimo et al. reported on the preparation of nanostructured VHx prepared by mechanical milling under H2 atmosphere [47]. Formation of the b2 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 b2 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 b2 phase remains in crystalline state.
From the relationship between hydrogen concentration and unit cell, the hydrogen concentration
2.0
320
8.0
315
1.5
305
7.0
H/Ti
310
T (K)
P (MPa)
7.5
1.0
300
0.5
6.5
(A)
6.0
0
60
120
t (min)
180
295
290
240
(B)
0.0
0
60
120
180
240
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.
43
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 b2–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.
3.2. Ternary hydrides
Ternary metal hydrides of general composition ABHx 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 LaNi5 have been used in RMM experiments
(see Table 1). As a general behaviour, two effects are observed on prolonged milling: formation of
amorphous ABHx hydrides and compound disproportionation into AHx + 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.
3.2.1. ZrNi hydride
RMM experiments in the Zr–Ni system have been first reported by Aoki et al. [27]. They performed
RMM (p(H2) = 2 MPa) of both arc-melted ZrNi alloys and equiatomic Zr and Ni powder mixtures. In the
first case, ZrNiH3 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 ZrNiH3 hydride because of its brittleness. For
RMM of elemental powders, a mixture of ZrH2 and elemental Ni is formed at short milling time
(<3 h). Further reaction over 100 h leads to the coexistence of ZrH2 and amorphous Zr-poor NiZr1yHx
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-ZrNiH3 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
ZrH2 and amorphous Zr-poor NiZr1yHx phase. Such decomposition reaction seems to be delayed with
increasing hydrogen pressure.
3.2.2. TiNi hydride
RMM experiments (p(H2) = 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 TiH2 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 TiNiH3 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 TiNiH3 phase [120]. Instead, formation of poorly crystallized TiH2 and Ni phases on milling
for 40 h is reported.
3.2.3. TiFe hydride
Chiang et al. have performed RMM experiments (p(H2) = 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 TiFeH1.6. Ex situ XRD analysis reveals that the ternary hydride decomposes to TiH2 and Fe.
3.2.4. LaNi5 hydride
Fujii et al. have performed RMM experiments (p(H2) = 1 MPa) on LaNi5 alloy and observe formation
of solid solution a-LaNi5H0.15 within the first 5 min of milling [121]. Formation of hydride b-LaNi5H6
phase is not detected, which may indicate that the absorption plateau pressure of this hydride is above
44
1 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 LaNi5yHx 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 LaNi5 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].
3.2.5. 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(H2) = 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 TiVH0.9 solid solution and FCC TiVH4.7
hydride are formed, respectively. The hydrogen content of FCC TiVH4.7 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
TiVH2.8 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 TiH2 and V, whereas
for the arc-melted alloy it results from gradual amorphization of BCC TiVH2.8 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 Ti0.20V0.78Fe0.02 have been carried out in a
device equipped with p and T sensors at rotation speed of 400 rpm and p(H2) = 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)H2 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 VH2-type hydride and amorphous
2.0
H/f.u.
1.5
1.0
0.5
0.0
0
60
120
180
240
300
360
420
480
t (min)
Fig. 9. Time-evolution of the H concentration in BCC Ti0.20V0.78Fe0.02 alloy during RMM in hydrogen gas.
45
800
I (counts)
600
amorphous phase
400
200
0
30
40
50
60
70
80
2θ
Fig. 10. Rietveld analysis of Ti0.20V0.78Fe0.02 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 a1,2) for FCC VH2-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.
3.3. Mg-based complex hydrides
Mg-based Mg2THx 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 Mg2FeH6, Mg2CoH5 and Mg2NiH4, 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 Mg2Fe and Mg2Co intermetallic compounds in their respective
binary phase diagrams. From these facts, synthesis of Mg2THx 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 Mg2THx ternary hydrides by RMM under hydrogen gas are summarized in
Table 2.
3.3.1. Mg2Fe hydride
The synthesis of Mg2FeH6 by RMM (p(H2) = 1 MPa for 20 h) of Mg and Fe powder was first attempted in 1997 [133]. Mg2FeH6 hydride was not obtained but Mg powder got hydrogenated to form
intimate MgH2–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 MgH2–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 MgH2 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 Mg2FeH6 formation without subsequent sintering. Much
probably, the following solid-state reaction takes place:
3MgH2 ðsÞ þ FeðsÞ ! Mg2 FeH6 ðsÞ þ MgðsÞ
ð1Þ
The formation of the ternary compound is driven by the fact that the Mg2FeH6 phase is more stable
than MgH2 [140]. In situ SR-PXD patterns measured for a ball milled sample of MgH2–Fe (2:1) reveal
formation of Mg2FeH6 at 673 K at p(H2) = 10 MPa [141].
46
Table 2
Mg-based complex hydrides synthesised by RMM under hydrogen gas. The employed device, reactants, initial hydrogen pressure,
p(H2), 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(H2)
(MPa)
tmt
(h)
ms
BTPWR Bd
Yield
(rpm)
(mm) (wt%)
Side
Ref.
products
Mg2FeH6
Mg2FeH6
Fritsch P5
Spex 8000
2Mg + Fe
2MgH2 + Fe
1
0.1
20
60
325
4:1
10:1
10
10
0
0
Mg2FeH6
Mg2FeH6
Mg2FeH6
Uni-Ball-Mill II
Szegvari attritor
Retsch 2000
vibrating mill
Fritsch P4a
Kurimoto
planetary mill
Uni-Ball-Mill II
Fritsch P4a
Fritsch P5
2Mg + Fe
2Mg + Fe
2Mg + Fe
0.5
1
0.3
60
20
8
44:1
400
20:1
1
32 s
16:1
6
12
28
63
90b
MgH2, Fe
Mg, MgO,
Fe
MgO, Fe
Fe
Fe
2Mg + Fe
2MgH2 + Co
7.5
0.1
12
10
400
700
60:1
8:1
12
4
77
100
2Mg + Co
2Mg + Co
2Mg + Ni
0.5
7.5
0.5
90
12
22
400
325
44:1
60:1
4:1
12
10
50
81
0
Fritsch P7
Kurimoto
planetary mill
Retsch 2000
vibrating mill
Fritsch P4a
Fritsch P6a
Mg2Ni
Mg2Ni
1
1
80
10
400
885
30:1
5:1
7
7
100
70b
2Mg + Ni
0.3
16
32 s1 16:1
12
100
Co
[135]
MgO
[116]
Mg,
[29]
MgH2, Ni
[136]
amph[137]
MgNi
[51]
2Mg + Ni
7.5
4Mg + Fe + Co 5
12
20
400
400
12
10
79
95b
MgO
FeCo
Mg2FeH6
Mg2CoH5
Mg2CoH5
Mg2CoH5
Mg2NiH4
Mg2NiH1.8
Mg2NiH4
Mg2NiH4
Mg2NiH4
Mg2(FeH6)0.5(CoH5)0.5
a
b
60:1
40:1
MgO, Fe
[133]
[134]
[49]
[115]
[51]
[116]
[48]
[116]
[52]
Estimated values.
In 2002, direct though incomplete synthesis of Mg2FeH6 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]. Mg2FeH6 formation with a
yield of 28 wt% was achieved after 60 h of milling. The synthesis was reported to occur in two steps.
MgH2 is formed during the first 40 h by mechanically activated solid–gas reaction followed by the solid-state reaction between MgH2 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 Mg2FeH6 and Fe phases, respectively. Formation on MgH2 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 2MgH2 + 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 2MgH2 + 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 2MgH2 + 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 Mg2FeH6 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 s1 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
47
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 MgH2 during the first 2 h of milling followed by the formation of Mg2FeH6 from MgH2 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 MgH2 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(H2) = 7.5 MPa for 12 h) of Mg2FeH6 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.
6
2Mg+Fe
Hydrogen uptake (H/f.u.)
5
nd
4
2 step
3.4 H/f.u.
3
2
2Mg
1
st
1 step
t = 50 min
0
0
60
120
180
240
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).
48
The first step, with t < 50 min, corresponds to the formation of MgH2 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 MgH2 and Fe to
form Mg2FeH6. The XRD pattern for 12 h milled product (Fig. 12) shows the presence of three phases
Mg2FeH6 (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 Mg2FeH6 synthesis could be described as:
2MgðsÞ þ FeðsÞ þ 3H2 ðgÞ ! 2MgH2 ðsÞ þ FeðsÞ þ H2 ðgÞ ! Mg2 FeH6 ðsÞ
ð2Þ
It is worth noting that reaction kinetics for the first step, i.e. MgH2 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)H2 solid solution which becomes unstable with increasing Fe content with respect to
Mg2FeH6 [148].
As we saw in the previous sections, ball milling has been extensively used to synthesize Mg2FeH6.
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 (0 0 2) 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 Mg2FeH6 [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].
3.3.2. Mg2Co hydride
Similarly to Mg2FeH6, first attempts to produce Mg2CoH5 by RMM (p(H2) = 1 MPa for 20 h) of Mg
and Co powder were unsuccessful. Instead, a mixture of MgH2 and Co phases was obtained [133].
Subsequent sintering treatment allowed synthesizing Mg2CoH5 hydride though with a lower yield
(26 wt%) than for Mg2FeH6 (65 wt%).
The synthesis of Mg2CoH5 hydride by RMM (p(H2) = 0.1 MPa for 10 h) of MgH2 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 Mg2CoH5 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 Mg2CoH5 phase
starts forming after 40 h of milling and a yield of 50 wt% was achieved at 90 h. The formation of an
intermediate MgH2 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(H2) = 7.5 MPa for 12 h) of Mg2CoH5 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:
2MgðsÞ þ CoðsÞ þ 5=2H2 ðgÞ ! 2MgH2 ðsÞ þ CoðsÞ þ 1=2H2 ðgÞ ! Mg2 CoH5 ðsÞ
ð3Þ
MgH2 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 MgH2 and Co to form
49
5
2Mg+Co
4
nd
2 step
3.4 H/f.u.
3
2Mg
2
1
st
1 step
t = 50 min
0
0
60
120
180
240
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).
Mg2CoH5. The hydrogen uptake corresponding to this reaction is lower but faster than for Mg2FeH6
formation (Fig. 11). The XRD pattern for 12 h milled product (Fig. 14) shows the formation of nanocrystalline (8 nm) Mg2CoH5 phase without significant amorphous contribution.
3.3.3. Mg2Ni hydride
Since Mg2Ni compound exists as stable phase in the binary Mg–Ni phase diagram, the synthesis of
Mg2NiH4 ternary hydride by RMM can be attempted either using 2Mg + Ni elemental mixture or
Mg2Ni powders as initial reactants.
This synthesis was first tried by RMM (p(H2) = 0.5 MPa for 22 h) of 2Mg + Ni powders [29]. Some
MgH2 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 Mg2Ni during reactive milling (p(H2) = 1 MPa
for 80 h) using stronger energetic conditions: rotation speed of 400 rpm and ball-to-powder weight
50
ratio 30:1 [136]. Instead of ternary hydride formation, a gradual solid solution of hydrogen in the
Mg2Ni powder with a maximum hydrogen content of 1.8 H/f.u. after 80 h is reported. The position
of XRD diffraction lines does not depend however on the hydrogen content. The milled material has
a two-phase microstructure formed by nanocrystalline (15 nm) Mg2Ni regions with low H-content
(0.3 H/f.u.) and disordered intergrain regions that store high amounts of hydrogen. Such a particular
microstructure has not later been confirmed by other research groups.
Tessier et al. also performed RMM experiments using Mg2Ni powder as starting reactant. RMM
experiments (p(H2) = 1.0 MPa for 10 h) were performed in a Kurimoto planetary mill at a high rotation
speed of 885 rpm [137]. Contrary to Orimo et al., the formation of disordered intergrain regions with
high H-content is not observed. The synthesis of 70 wt% of Mg2NiH4 ternary hydride, as estimated
from total hydrogen content in the milled product, is detected. The hydride crystallizes as a mixture
of low- and high-temperature phases of Mg2NiH4 hydride [151,152]. The high temperature modification is usually formed above 510 K but it seems to be stabilized by mechanical milling [153].
Baum et al. have monitored the in situ hydrogen uptake during RMM of 2Mg + Ni elemental powder
mixture with identical milling parameters as mentioned above for the synthesis of Mg2FeH6 [51]. Successful formation of Mg2NiH4 ternary hydride is reported after 16 h of milling. The hydrogen uptake
curve exhibits a unique step. This result differs from experiments on Mg2FeH6 and Mg2CoH5 formation
and was tentatively attributed to a simpler process for Mg2NiH4 related to the existence of Mg2Ni
compound. However, this seems not to be the case. Fig. 15 displays the in situ hydrogenation curves
monitored during the synthesis of Mg2NiH4 hydride [116]. Indeed, one-hydrogenation step for
t < 50 min is observed. As proved by XRD diffraction studies of the sample milled for this time (top
pattern in Fig. 16), the reaction is however not completed since a high amount of the starting Ni powder remains unreacted. Further milling for 12 h leads to complete formation of Mg2NiH4 hydride as
shown by the XRD pattern displayed in bottom Fig. 16. This second step on the formation of the ternary hydride is a solid–solid state reaction, i.e. it occurs without any hydrogen absorption. The reaction path of Mg2NiH4 formation is then described according to reaction scheme:
2MgðsÞ þ NiðsÞ þ 2H2 ðgÞ ! 2MgH2 ðsÞ þ NiðsÞ ! Mg2 NiH4 ðsÞ
ð4Þ
According to this reaction and to the previous experiments reported by Baum et al. [51] and Zhang
et al. [116], the synthesis of all Mg-based Mg2THx ternary hydrides (T = Fe, Co and Ni) by RMM of elemental pure powders under hydrogen atmosphere occurs through a common reaction path:
x
ðx 4Þ
H2 ðgÞ ! Mg2 THx ðsÞ
2MgðsÞ þ TðsÞ þ H2 ðgÞ ! 2MgH2 ðsÞ þ TðsÞ þ
2
2
ð5Þ
with x = 6, 5 and 4 for T = Fe, Co and Ni, respectively.
4
2Mg+Ni
3.4 H/f.u.
3
2Mg
2
1
st
1 step
t = 50 min
0
0
60
120
180
240
Milling time (min)
Fig. 15. Time-evolution of the H concentration in solid-state during RMM of Mg powder and 2Mg + Ni powder mixture.
51
Fig. 16. XRD patterns and phase identification of RMM 2Mg + Ni powder mixture after the first (50 min) and after 720 min of
milling (Cu Ka radiation).
Baum et al. have also reported the possibility of synthesizing quaternary Mg2T0.5T’0.5Hx hydrides
with T, T’ = Fe, Co and Ni [51]. This result has been confirmed by Deledda and Hauback for T = Fe
and T’ = Co [52]. The later authors obtained a quaternary hydride of composition Mg2(FeH6)0.5(CoH5)0.5
in which both [FeH6]4 and [CoH5]4 and complex anions coexist in the same compound. The synthesized compound desorbed hydrogen in a one-step reaction at temperatures between 500 and 600 K
which is between those of Mg2FeH6 and Mg2CoH5. These results open up the possibility of synthesizing other Mg-based transition–metal complex hydrides, in which the hydrogen storage properties can
be tailored by varying the transition metals and the relative content of the complex anion.
3.4. Alanates
Since the pioneering work of Bogdanović and Schwickardi, who found that Ti-catalyzed NaAlH4 can
release hydrogen under moderate conditions, the potential interest of metal aluminum complex hydrides, also known as alanates, for reversible hydrogen storage remains vivid [154]. Alanates are
formed by anionic aluminum–hydrogen complexes (typically either [AlH4] or [AlH6]3) stabilized
by a cation (typically an alkali or alkaline earth metal) [2,155]. Their gravimetric (up to 10.8 wt%
though only 5.6 wt% reversible) and volumetric capacities (90 g/L) make these compounds attractive
for hydrogen storage.
Conventionally, alanates are prepared by a wet chemical route. For instance, to synthesize sodium
alanate, NaH and Al are diluted in THF and hydrogen pressure up to 20 MPa at 423 K is applied for few
days. Subsequently, the NaAlH4 in solution is filtered and dried. Since the solubility of NaAlH4 is very
low, a high amount of solvent is necessary. Dymova et al. have shown that the synthesis of NaAlH4 can
also be obtained from Na, Al and H2 at high temperature (553 K), where Na is in liquid state, and high
hydrogen pressure (17.5 MPa) [156].
In 1999, mechanochemical synthesis methods started being employed to synthesize Na3AlH6 and
Na2LiAlH6 complex hydrides using NaAlH4 and LiAlH4 as reagents [157,158]. Very recently, it has been
shown that NaAlH4 can be formed by milling NaH and Al under a hydrogen atmosphere of 8.3 MPa,
4 mol% of TiCl3 was added as a dopant [30]. The alanates synthesized by reactive mechanical milling
are surveyed hereafter (Table 3).
3.4.1. Lithium alanates
The synthesis of LiAlH4 and Li3AlH6 has been attempted by Kojima et al. under 1 MPa of hydrogen
using a planetary ball mill at 400 rpm [165]. After 24 h of milling with addition of TiCl3, the XRD pattern shows broad peaks of LiH and Al phases but no clear signature of alanate formation. However, the
Raman and 27Al MAS NMR spectra indicate that a small amount of LiAlH4 was formed. Actually, lithium alanate LiAlH4 is metastable and it is considered as a non-reversible hydrogen storage compound
52
Table 3
Alanates synthesised by RMM under hydrogen gas. The employed device, reactants, initial hydrogen pressure, p(H2), 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
NaAlH4
NaAlH4
NaAlH4
NaAlH4
Na3AlH6
Na3AlH6
Na3AlH6
KAlH4
Na2LiAlH6
K2NaAlH6
Mg(AlH4)2
Ca(AlH4)2
CaAlH5
Ba2AlH7
BaAlH5
a
Reactants
p(H2)
(MPa)
tmt
(h)
ms
(rpm)
BTPWR Bd
(mm)
Yield
(wt%)
Fritsch P7 NaH + Al
Fritsch P5 NaH + Al
8.3
0.6
4
120
500
350
50:1
50:1
10
Low
Fritsch P5
Planetary
Fritsch P6
Planetary
Fritsch
P4a
Fritsch P7
Fritsch
P4a
Fritsch
P4a
Fritsch P7
Fritsch P6
Fritsch P7
Fritsch P5
Fritsch P5
1.2
3
0.85
0.5
10
240
60
20
30
8
230
350
400
350
400
10:1
30:1
40:1
30:1
90:1
10
8
High
86
8
15
KH + Al
10
LiH + 2NaH + Al 10
10
4
400
400
90:1
2KH + NaH + Al 10
2
400
4
10
3
10
10
400
500
400
300
300
NaH + Al
NaH + Al
NaH + Al
NaH + Al
3NaH + Al
2AlH3 + MgH2
2AlH3 + CaH2
AlH3 + CaH2
BaH2 + Al
BaH2 + 2Al
0.1
0.1
0.1
0.8
0.8
Side products
Ref
[30]
[53]
58
92
NaCl, Al
Na3AlH6, NaH,
Al
NaCl, Al
Na3AlH6, Al
Al
Al
NaCl, NaH, Al
15
97
NaCl
[161]
[57]
90:1
15
98
NaH
[57]
175:1
10
13
10
81
AlH3, MgH2, Al [162]
[163]
CaH2, Al
[162]
BaH2, Al
[164]
Al
[164]
80:1
40:1
40:1
91
84
83
[53]
[159]
[160]
[159]
[57]
under moderate temperature and pressure conditions [166,167]. Indeed, the desorption of hydrogen
from LiAlH4 is reported to be an exothermic process [161]. Lithium alanate does not form directly from
LiH, Al and H2 by solid–gas reaction. It is classically produced by using liquid complexing agents followed by desolvatation [168,169].
3.4.2. Sodium alanates
The sodium tetra-alanate NaAlH4 is the best performing and most studied among all the alanates.
Bellosta von Colbe et al. have synthesized NaAlH4 by RMM (p(H2) = 8.3 MPa for 4 h) of NaH and Al
powders using 4 mol% TiCl3 as a dopant [30]. The obtained alanate exhibits a reversible capacity of
4 wt% and very fast kinetics without activation. In the same year, Wang et al. also tried the synthesis
of NaAlH4 by the same means but at a lower pressure of 0.85 MPa and using TiF3 as a dopant [160].
Within the first 5 h of milling, alanate formation was not observed. However, further milling up to
20 h led to partial synthesis of the more stable Na3AlH6 phase.
Eigen et al. have also tried the synthesis of NaAlH4 at moderate hydrogen pressures (from 0.6 to
1.2 MPa) [53]. NaAlH4 could be formed with and without the addition of TiCl4 as catalyst. However,
a long milling time (100 h without TiCl4, 20 h with TiCl4) is needed to obtain a small fraction of NaAlH4
under 0.6 MPa of hydrogen. Generated heat due to material plastic deformation and the friction between the milling balls and vial wall during milling could lead to a temperature at which the NaAlH4
is not stable at the imposed hydrogen pressure. A complete formation of NaAlH4 could be achieved by
reducing the milling energy (ball-to-powder weight ratio and rotational speed were decreased from
50:1 and 350 rpm to 10:1 and 230 rpm) and increasing the hydrogen pressure (1.2 MPa).
Xiao et al. have also actively worked on the synthesis of NaAlH4 by RMM under hydrogen pressure
of 0.5–3 MPa [159]. They observed that while Na3AlH6 forms under 0.5 MPa hydrogen pressure,
NaAlH4 only forms above 0.8 MPa [159]. These results are similar to those of Eigen et al. and prove that
the formation of NaAlH4 needs a minimum hydrogen pressure about 1 MPa, which is much higher
than the equilibrium pressure of NaAlH4 at room temperature (about 0.1 MPa) [53]. This is attributed
to temperature rising on milling. By increasing the hydrogen pressure to 3 MPa, the synthesis yield of
NaAlH4 reached 87 wt% after milling for 60 h.
53
T (K)
310
305
300
Na 3AlH 6
295
Po = 9.6 MPa
Na 3AlH 6 (TiCl 3 doped) Po = 10.2 MPa
1.12
P/Po
1.08
1.04
1.00
0.96
0
60
120
180
240
300
360
420
480
Milling time (min)
Fig. 17. Time-evolution of the hydrogen pressure (bottom) and temperature (top) during the synthesis of Na3AlH6 by RMM
with (empty symbols) and without (full symbols) TiCl3 dopant [57].
Sodium hexa-alanate is not available commercially. Some research groups have synthesized the
hexa-sodium alanate by ball milling a mixture of the tetra alanate and the sodium hydride through
the solid-sate reaction 2NaH(s) + NaAlH4(s) ? Na3AlH6(s) without any additives [157,158]. By using
RMM (p(H2) = 10 MPa for 4 h), high yield is obtained through the following reaction [57]:
3
2mol%TiCl3
3NaHðsÞ þ AlðsÞ þ H2 ðgÞ ! Na3 AlH6 ðsÞ
2
ð6Þ
This experiment was performed in an Evico-magnetics vial. Fig. 17 shows the temperature and hydrogen pressure inside the milling vial for the formation of Na3AlH6 with and without TiCl3. In both cases,
the temperature first increases for t < 30 min due to the heat generated on milling before reaching a
plateau around 310 K due to the balance between internal friction heating and ventilation cooling.
The hydrogen pressure increases simultaneously because of the temperature rise in the vial. At longer
milling time, t > 30 min, the pressure drops markedly for the doped sample whereas it decreases much
more slowly for the undoped one.
A further advantage on the synthesis of sodium alanates by RMM is the simultaneous incorporation
of dopant. The addition of TiF3[159,160], TiCl4 [170], TiCl3 [57], ScCl3 and CeCl3 [171] dopants during
milling not only speeds up the formation of NaAlH4 or Na3AlH6 compounds, but also promotes the dispersion of the dopant in the alanate. The materials synthesized by one-step reactive mechanical milling displayed superior features compared to doping of pre-synthesized NaAlH4 with TiCl3
[30,172,173].
3.4.3. Potassium alanates
Potassium alanate KAlH4 as well as K3AlH6 are stable alanates under normal conditions of pressure
and temperature. They decompose at moderate temperatures (above 550 K) [174]. Up till now no reports have been published of the direct synthesis of K3AlH6 by reactive ball milling. KAlH4 can be synthesized by RMM of KH + Al powder mixture after 10 h milling under 10 MPa of hydrogen pressure
[161].
3.4.4. Mixed alkali alanates
Three mixed alkaline alanates are reported to exist: Na2LiAlH6, K2LiAlH6 and K2NaAlH6 [175–177].
In 1999, Na2LiAlH6 was synthesized by ball milling of NaAlH4 and LiH powder mixtures [157]. Recently, Na2LiAlH6 and K2NaAlH6 mixed alanates have been synthesized by RMM (p(H2) = 10 MPa for
4 h) [57]. These compounds were prepared according to reaction schemes [56,57]:
3
2 mol%TiCl3
LiHðsÞ þ 2NaHðsÞ þ AlðsÞ þ H2 ðgÞ ! Na2 LiAlH6 ðsÞ
2
ð7Þ
54
3
2 mol%TiCl3
NaHðsÞ þ 2KHðsÞ þ AlðsÞ þ H2 ðgÞ ! K2 NaAlH6 ðsÞ
2
ð8Þ
In situ hydrogen uptake curves proved that the above reactions typically take place in 2 h. Ex situ XRD
analysis after milling demonstrate the completion of the reaction as can be seen, for instance, for Na2LiAlH6 synthesis in Fig. 18. In addition, a small amount of NaCl is detected as a result from the reaction
of TiCl3 with NaH [113]:
3
3NaHðsÞ þ TiCl3 ðsÞ ! 3NaClðsÞ þ TiðsÞ þ H2 ðgÞ
2
ð9Þ
I (a.u)
Jeloaica et al. have attempted to synthesize ternary alkaline hexa-alanates of composition Li4/3Na1/3K4/
3AlH6 and LiNaKAlH6 [56]. However these compounds could not be obtained by RMM. The obtained
20
30
40
50
60
70
80
2θ
I (a.u)
Fig. 18. Rietveld analysis of RMM Na2LiAlH6: observed (dots), calculated (solid line) and difference curves (bottom) are shown.
Vertical bars (|) correspond to Bragg positions for Na2LiAlH6 (top) and NaCl (bottom) phases (Cu Ka1,2).
20
30
40
50
60
70
80
2θ
Fig. 19. Rietveld analysis RMM LiNaKAlH6: observed (dots), calculated (solid line) and difference curves (bottom) are shown.
Vertical bars (|) correspond to Bragg positions for K2NaAlH6, KAlH4, LiNa2AlH6, NaH, LiH and NaCl phases from top to bottom (Cu
Ka1,2) [56].
55
materials contain a mixture of alkaline hydrides (LiH, NaH and KH), the tetra-alanate KAlH4 and binary
hexa-alanate (NaK2AlH6 and Na2LiAlH6) phases (Fig. 19). Actually, it has been shown by DFT calculations that the nominal Li4/3Na1/3K4/3AlH6 and LiNaKAlH6 phases are not stable with respect to the mixture of LiH, the tetra-alanate KAlH4 and binary hexa-alanate K2NaAlH6 or Na2LiAlH6.
3.4.5. Alkali-earth alanates
Conventionally, Mg(AlH4)2 and Ca(AlH4)2 have been synthesized by wet chemistry through the
metathesis reaction between MCl2 (M = Mg and Ca) and NaAlH4 [178,179]. From the same reagents
but by using mechanochemical milling, several research groups have also synthesized the magnesium
and calcium alanates [180–182]. But all these approaches produce significant amount of magnesium
or calcium chloride as by-products. These salts are ‘‘inert’’ to hydrogen and thus decrease the H-storage capacity of the final products.
Reactive mechanical milling of Mg(AlH4)2 alanate has been attempted in 2005 by Varin et al. using
a magneto-ball mill under 0.8 MPa hydrogen [183]. Three stoichiometric Mg + 2Al mixtures, (a) elemental Mg and Al powders, (b) elemental Al powder and commercial AZ91 alloy (Mg–Al–Zn alloy)
and (c) powder of as-cast Mg + 2Al alloy have been used. No successful synthesis of Mg(AlH4)2 has
been achieved. The hydride formed after 270 h of milling is b-MgH2. In 2009, several groups tried
the synthesis of Mg-containing quaternary alkali alanates by RMM under a much higher pressure of
hydrogen (10 MPa). No traces of Mg(AlH4)2 alanate were however detected [184,185]. The lack of success in the synthesis of Mg(AlH4)2 alanate by RMM could be related to the lower thermodynamic stability of Mg(AlH4)2 as compared to MgH2.
Interestingly, it has been demonstrated that Mg(AlH4)2 can be obtained by RMM (p(H2) = 10 MPa
for 4 h) of 2AlH3 and MgH2 powder mixtures [162]. The reaction yield attained 81 wt% of Mg(AlH4)2.
Experiments were conducted either in argon or hydrogen atmosphere. Strikingly, the milling
atmosphere had minor impact on the yield of Mg(AlH4)2. In contrast, the most important parameter
was the pause used to allow vial cooling (2 min of pause following every 10 min of milling). In fact,
it was found that during milling, alane decomposes due to the temperature rise and reduces the yield
of Mg(AlH4)2 formation. This suggests that the alanate formation occurs through an all solid-state
process:
MgH2 ðsÞ þ 2AlH3 ðsÞ ! 2MgðAlH4 Þ2 ðsÞ
ð10Þ
This reaction could be driven by lower stability of alane as compared to magnesium alanate. Although
this synthesis route is successful, it should be taken into account that the preparation of alane is still a
challenge [186–188].
As for the calcium alanate, progress on its synthesis by mechanical means is quite similar to the
magnesium one. The synthesis of Ca–M–Al–H (M = Na and Li) quaternary alanates by RMM of binary
alkali hydrides and Al at 10 MPa of hydrogen pressure was unsuccessful [184]. In contrast, calcium
alanate can be formed by mechanical milling using CaH2 and alane as reactants [162,163]. Strikingly,
either Ca(AlH4)2 or CaAlH5 compounds are reported to be formed depending on the molar ratio between reactants: 2AlH3 + CaH2 in the former case and AlH3 + CaH2 in the latter.
Other alkaline earth alanates that have been synthesized by RMM under hydrogen atmosphere are
Ba2AlH7 and BaAlH5 [164]. RMM experiments (p(H2) = 0.8 MPa for 10 h) have been performed using
BaH2 and Al powders as reactants. Three ratios of the reagent powder BaH2/Al 2:1, 1:1 and 1:2 have
been used. Surprisingly, formation of Ba2AlH7 was favoured for 1:1 ratio as compared to 2:1. 2BaH2/Al
powder mixture forms 42 wt% of Ba2AlH7 (together with 58 wt% of BaH2) whereas BaH2/Al powder
mixture forms 84 wt% of Ba2AlH7. Elemental Al is not observed in the former case though it is expected
from mass balance considerations. As for the synthesis of BaAlH5, a high yield (83.2 wt%) of this phase
is obtained using the ratio 1:2, i.e. with Al in excess. The reason of these striking results was not
discussed.
RMM has also been used as intermediary step in the synthesis of Sr2AlH7 [189]. Sr2Al compound
was milled under hydrogen (0.6 MPa) to get intimate 2SrH2 + Al mixture. Such a mixture could react
with hydrogen at 553 K and 7 MPa in an external device to form Sr2AlH7. A similar approach has been
used for the synthesis of mixed (Sr,Ca)AlH7 compounds [190].
56
To summarize, mechanical milling under hydrogen gas is a very effective method for the synthesis
of hydrides. Most hydrides formed by this method can be also obtained by solid–gas reaction:
metastable hydrides are rarely formed. However, reactive mechanical milling is much faster than conventional solid–gas methods even for milder conditions of pressure and temperature. Main reasons for
this are probably clean surface generation and shortening of diffusion paths as result of mechanicallyinduced particle pulverization. In systems with more than one metal-constituent, extended and clean
interface formation between reactive powders should also play a major role to speed solid/solid
reactions.
Hydrides produced by RMM are generally nanocrystalline. Prolonged RMM may however lead to
amorphization and compound disproportionation as it occurs in ternary metallic ABHx hydrides. This
effect has not been observed so far in complex ternary hydrides which may indicate that ionic-covalent compounds are less prone to amorphisation than metallic compounds.
3.5. Synthesis of borohydrides by mechanical milling in diborane gas
A novel technique for reactive ball-milling in diborane atmosphere was developed at EMPA, Switzerland, by Remhof, Friedrichs, Borgschulte, Züttel and co-workers. Solvent-free synthesis of lithium
borohydride, LiBH4 from lithium hydride, LiH in diborane, B2H6 gas at 393 K has recently been demonstrated [191]. Diborane gas is produced by heating LiZn2(BH4)5 prepared by ball milling of a
2LiBH4:ZnCl2 mixture. This is considered a convenient and relatively safe source of diborane as compared to pressurized bottles, since diborane is a poisonous and explosive gas [192]. The formation of
LiBD4 by an addition reaction of LiD and B2D6 was analysed by in situ neutron diffraction, which shows
that nucleation of LiBD4 already starts at temperatures of ca. 375 K, i.e. formation of orthorhombic
LiBH4. However, the reaction is incomplete and the yield is only 50% even at elevated temperatures
(460 K) and the product contains small amounts of Li2B12H12. A passivation layer of LiBH4 is suggested
to form on the surface of the LiH grains retarding the process [193]. Therefore, a custom-made mill
connected to a gas/vacuum supply via a flexible hose made of polyamide was used. The mill was
equipped with two stainless steel milling vials performing horizontal movements based on the seesaw principle. One vial was equipped with ceramic balls. This approach allows continuous removal
of the borohydride surface layer, see Fig. 20 [194].
This new mechanochemical method allows solvent-free synthesis of borohydrides at room temperature demonstrated by the synthesis of three of the most investigated borohydrides at present: LiBH4,
Mg(BH4)2 and Ca(BH4)2. Similarly, Y(BH4)3 was prepared in a reaction with YH3 and B2H6 with this
new RBM [195]. This new gas–solid mechanochemical synthesis method is based on the reaction of
metal hydrides with diborane to form the corresponding borohydrides. The synthesis method may
facilitate preparation of a wide range of different borohydrides in the future. Furthermore, with this
Fig. 20. Schematic presentation of a custom-made mill connected via a flexible hose to a gas/vacuum supply [194]. The gas
supply consists of a container that is situated in an oven and is filled with a borane desorbing material, i.e. LiZn2(BH4)5.
57
technique it was demonstrated that Li2B12H12 and an amorphous Li2B10H10 phase can be prepared by a
reaction with LiBH4 and diborane (B2H6) at T 475 K. This tend to suggest that formation of higher
boranes, such as Li2B12H12, may occur during thermal decomposition of metal borohydrides [196].
Recently, a metal borohydride was prepared by reactive ball milling of a metal boride in hydrogen
atmosphere. The starting materials, MgB2 and hydrogen gas, p(H2) = 10 MPa, was charged into a hardened steel high pressure vial (ball to powder ratio of 30:1) and milled at a rotation speed of 600 rpm
for up to 100 h. The product, unsolvated amorphous magnesium borohydride (yield 50%) appeared
to form directly but impurities of higher boranes Mg(BnHm)y (30%) were also identified, which may
have formed in side reactions [197]. Interestingly, a new polymorph c-Mg(BH4)2 with 30% open space
in the structure transformed from crystalline to semi-amorphous by mechano chemical treatment
[198].
The hydrogen absorption mechanism for ball milled samples of 2NaH–MgB2 was also investigated
in detail. At high pressures, p(H2) 5 MPa, NaBH4 formed after observation of an unknown compound,
NaMgH3 and a NaH–NaBH4 molten salt mixture. In contrast, NaBH4 apparently formed directly at lower hydrogen pressures, i.e. p(H2) = 0.5 MPa. This indicates that the reaction mechanism may be modified by mechano-chemical treatment of the reactants and by the applied hydrogen pressure [199].
This work reveals that reactive ball milling of metal hydrides in diborane, metal boride or mixtures
of metal hydrides and borides in elevated hydrogen pressures has a potential for providing new convenient solvent-free routes for preparation of metal borohydrides.
3.6. Synthesis of metal amides by mechanical milling in ammonia gas
During the past decade lithium amide, LiNH2 has received much interest as a hydrogen storage
material [200,201]. Recently, LiNH2 and a range of other metal amides i.e. NaNH2, Mg(NH2)2 and
Ca(NH2)2 have been prepared by mechanochemical methods from metal hydrides, MHx and ammonia
gas, NH3 according to reaction scheme Eq. (11) [202,203].
MHx ðsÞ þ xNH3 ðgÞ ! MðNH2 ÞxðsÞ þ xH2 ðgÞ
ð11Þ
In general, gas loadings of p(NH3) 0.4–0.5 MPa was used with refilled approximately every second
hour. A well-known problem of utilization of the metal amides as hydrogen storage materials is the
release of NH3 gas besides hydrogen during the thermal decomposition [59,204,205]. In order to avoid
this, a metal hydride can be added to absorb the NH3 according to reaction scheme Eq. (11). The efficiency of this approach has been studied by milling LiH in p(NH3) = 0.4 MPa corresponding to LiH–NH3
(1:1). Gas chromatographic analysis revealed that 70% of the initial NH3 reacted with LiH and transformed to H2 after 30 min of milling [206].
4. Synthesis of hydrides by mechanically-induced solid/solid and solid/liquid reactions
4.1. Mechanochemical synthesis of metal borohydrides
Metal borohydrides are well known for their properties as reducing agents and neutron absorbers.
However, due to their high gravimetric hydrogen density this class of compounds has recently received increasing interest as potential hydrogen storage materials [2–4]. Schlesinger et al. reported
in 1953 the first mechanochemical synthesis of a metal borohydride i.e. sodium borohydride from sodium hydride and boric oxide according to reaction scheme [207].
4NaH þ 2B2 O3 ! NaBH4 þ 3NaBO2
ð12Þ
Similar methods are still investigated for preparation of alkali borohydrides, e.g. NaBH4 or KBH4
from MgH2 and Na2B4O7 or KBO2, respectively [208,209].
Due to the high thermal stability of the alkali borohydrides much attention has been given to synthesis and characterization of novel metal borohydrides e.g. based on alkaline earth and transition
metals. In 1989 Mal’tseva et al. showed that ball milling of alkali borohydrides, MBH4 (M = Li, Na or
K) with zinc chloride, ZnCl2 resulted in a metathesis reaction, i.e. formation of MCl was observed by
58
Table 4
Metal borohydrides synthesized by mechanochemical methods, reactants used for the synthesis, optimal reactant ratio, formed
side products, total milling time (tmt) and milling speed (rounds per minute, rpm) used in the synthesis.
Compound
Reactants
Opt. ratio
Side products
tmt (min)
rpm
Ref.
NaBH4
KBH4
Sr(BH4)Cl
Sr(BH4)2
LiSc(BH4)4
NaSc(BH4)4
KSc(BH4)4
Y(BH4)3
NaY(BH4)2Cl2
Mn(BH4)2
MgH2 + Na2B4O7
MgH2 + KBO2
LiBH4 + SrCl2
LiBH4 + SrCl2
LiBH4 + ScCl3
NaBH4 + ScCl3
KBH4 + ScCl3
LiBH4 + YCl3
NaBH4 + YCl3
LiBH4 + MnCl2
NaBH4 + MnCl2
LiBH4 + ZnCl2
NaBH4 + ZnCl2
NaBH4 + ZnCl2
KBH4 + ZnCl2
KBH4 + ZnCl2
LiBH4 + CdCl2
NaBH4 + CdCl2
KBH4 + CdCl2
KBH4 + CdCl2
LiBH4 + AlCl3
LiBH4 + LiCl
LiBH4 + LiBr
LiBH4 + LiI
NaBH4 + NaCl
Ca(BH4)2 + CaI2
4:1
2:1
1:1
1:1
4:1
2:1
2:1
3:1
2:1
2:1
2:1
5:2
5:2
3:1
2:1
1:1
2:1
14:9
1:1
4:3
13:3
–
–
–
–
–
MgO, B2O3
MgO
LiCl, Sr(BH4)2
LiCl, Sr(BH4)Cl
LiCl
Na3ScCl6
K3ScCl6
LiCl
Na3YCl6, Na(BH4)1xClx
LiCl
NaCl
LiCl
Na2ZnCl4, NaCl
Na2ZnCl4, NaCl
K2Zn(BH4)xCl4x, K3Zn(BH4)xCl5x
–
LiCl
NaCl, Na6CdCl8
KCdCl3, K2Cd(BH4)4, Cd(BH4)2
KCdCl3
LiCl
–
–
–
–
–
60
120
120
120
180
120
120
120
120
350
350
120
120
120
350
120
30
30
20
20
300
120
120
120
120
120
2750
490
400
400
500
400
400
200
200
600
600
200
200
200
–
200
200
200
200
200
500
200
200
200
200
250
[207,209]
[208]
[39]
[39]
[211,212]
[213]
[214]
[215,216]
[38]
[40]
[40]
[217]
[217]
[217]
[218]
[219]
[37]
[37]
[37]
[37]
[220]
[221,222]
[223]
[224]
[225]
[226]
LiZn2(BH4)5
NaZn2(BH4)5
NaZn(BH4)3
KZn(BH4)3
KZn(BH4)Cl2
Cd(BH4)2
KCd(BH4)3
K2Cd(BH4)4
Li4Al3(BH4)13
Li(BH4)0.9Cl0.1
Li(BH4)0.47Br0.53
Li(BH4)0.3I0.7
Na(BH4)0.9Cl0.1
Ca(BH4)1.6I0.4
PXD [210]. Such mechanochemical methods have become one of the most common methods for preparation of novel metal borohydrides during the past decade and a significant increase in the number of
new compounds and studies of their structural, physical and chemical properties [3,4]. Table 4
provides an overview of reactants and conditions for synthesis of metal borohydrides along with side
products, which are obtained in some cases.
During the mechanochemical synthesis a variety of reactions can occur and in some cases several
competing reactions are observed simultaneously. Metathesis, or double substitution reaction, is a
well-known mechanism for chemical reactions during ball milling, here illustrated by the reaction between lithium borohydride LiBH4 and yttrium(III)chloride, YCl3, which results in formation of Y(BH4)3
and LiCl according to reaction scheme [215,216].
ð1 : 3Þ YCl3 þ 3LiBH4 ! YðBH4 Þ3 þ 3LiCl
ð13Þ
The polymorph a-Y(BH4)3, previously obtained from diethyl ether solutions of LiBH4 and YCl3 at RT
and also by MM, can be obtained with varying amounts of a new polymorph, denoted b-Y(BH4)3. aY(BH4)3 can be considered as the stable phase at ambient conditions, whereas b-Y(BH4)3, is a hightemperature polymorph formed by an a to b polymorphic phase transition. The high-temperature
b-polymorph can be quenched to ambient conditions [215,216,227,228].
It should be noted that many single-cation metal borohydrides have been reported to be prepared
by mechanochemical synthesis e.g. Zn(BH4)2 [229,230], Sc(BH4)3 [229], Ti(BH4)3 [229,231], V(BH4)2
and Cr(BH4)2 [232]. However, no structural data have been reported for these compounds and their
formations are based on observation of alkali chlorides and assumption of a simple metathesis reaction similar to reaction scheme Eq. (13). Further investigations of some of these systems have shown
that more complex reactions often take place [213,214,217,219].
The system ZnCl2–MBH4 (M = Li, Na or K) can be used to illustrate the complexity of the mechanochemical synthesis. Ball milling a mixture of ZnCl2–KBH4 (1:1) leads to an addition reaction and a
phase pure product, of KZn(BH4)Cl2 [219].
ZnCl2 þ KBH4 ! KZnðBH4 ÞCl2
59
ð14Þ
3
Interestingly, the unit cell volume of KZn(BH4)Cl2 (V/Z = 149.5 Å ) is nearly equal to the sum of formula volumes for the reactants ZnCl2 (V/Z = 74.3 Å3) and KBH4 (V/Z = 76.2 Å3). However, there are significant differences between the structure of reactants and the product. The latter contains the
heteroleptic complex ion [Zn(BH4)Cl2] where Zn coordinates to two chloride ions and two hydrogens
in g2–BH4, i.e. CN(Zn) = 4. This clearly demonstrates that ball milling induces a complex chemical reaction involving bond breaking and bond formation. In this case, the structure is fully ordered. The addition reaction may also provide a solid solution, e.g. Li(BH4)1xClx, which will be further discussed later
in this section.
The mechanochemical synthesis of other new borohydrides in the system ZnCl2–MBH4 proceeds
via more complex chemical reactions during ball milling [217]
2ZnCl2 þ 5LiBH4 ! LiZn2 ðBH4 Þ5 þ 4LiCl
ð15Þ
2ZnCl2 þ 5NaBH4 ! NaZn2 ðBH4 Þ5 þ 4NaCl
ð16Þ
ZnCl2 þ 3NaBH4 ! NaZnðBH4 Þ3 þ 2NaCl
ð17Þ
Reaction schemes (16) and (17) illustrates that small deviations in the composition of reactants may
lead to significantly different reaction products both in terms of the stoichiometry and the structural
topology, i.e. the structures of NaZn(BH4)3 and NaZn2(BH4)5 are significantly different, which may suggest that the synthesis mechanism for these compounds is also different. The compounds LiZn2(BH4)5
and NaZn2(BH4)5 are isostructural and built from two identical interpenetrated three-dimensional
frameworks consisting of isolated complex anions, [Zn2(BH4)5], whereas NaZn(BH4)3 consists of a single three-dimensional network, containing polymeric anions with the composition ½ZnðBH4 Þ3 n
n (see
Fig. 21) [42,217].
These new compounds were prepared by high-energy MM in short intervals (2 min) separated by
pauses (2 min). This procedure suppresses heating of the sample by friction heat and keeps the tem-
Fig. 21. Crystal structure of LiZn2(BH4)5 (left) built from two identical interpenetrated three-dimensional frameworks
consisting of isolated complex anions, [Zn2(BH4)5], and NaZn(BH4)3 (right) consisting of a single three-dimensional network,
containing polymeric anions with the composition ½ZnðBH4 Þ3 [42,217].
60
perature close to room temperature. This suggests that the synthesis is facilitated by high pressure
rather than elevated temperature. In fact, this class of M–Zn–BH4 compounds decomposes irreversibly
by heating at 373 K or prolonged MM with release of diborane [36]. Furthermore, these compounds
are metastable, i.e. they decompose when stored at RT within approximately 1 week. These observations suggest that ball milling may lead to a chemical equilibrium state rather than just a statistical
distribution of reactants.
Reaction scheme Eqs. 13,15,16, and 17 also illustrate a general drawback of the mechanochemical
approach involving metathesis reactions, namely that the products may be contaminated with ionic
compounds, most often binary metal halides but sometimes also ternary halides. Formation of
LiZn2(BH4)5 (reaction scheme Eq. (15)) proceeds completely, i.e. with formation of LiCl as the only side
product, whereas formation of NaZn2(BH4)5 and NaZn(BH4)3 (reaction schemes Eqs. (16) and (17))
only proceeds partly due to formation of a ternary metal chloride by a simultaneous and competing
reaction described in reaction scheme [42].
ð1 : 2Þ ZnX 2 þ 2MX ! M 2 ZnX 4
ð18Þ
Mechanochemical synthesis of M2ZnX4, M = Li or Na, X = Cl or Br from stoichiometric mixtures of
MX and ZnX2 has previously been reported by Solinas et al., who state that reaction times and activation energy decreases as Li2ZnCl4 > Na2ZnCl4 > Na2ZnBr4 [233]. This is in agreement with the results
for the zinc-based borohydride systems ZnX2–MBH4 (M = Li, Na, K and X = Cl, Br), where formation
of M2ZnX4 has increasing dominance over the heavier elements, i.e. K > Na Li and Br > Cl. However,
reaction Eq. (18) is only weakly coupled with the formation of the borohydrides. In other words, reaction Eqs. (16) and (17) are faster than reaction Eq. (18). This contrasts the mechanochemical synthesis
of NaSc(BH4)4 and KSc(BH4)4 from ScCl3 and NaBH4 or KBH4, respectively, which is suggested to proceed as described by reaction scheme [213,214].
ð1 : 4Þ ScCl3 þ 4NaBH4 ! NaScðBH4 Þ4 þ 3NaCl
ð19Þ
ð1 : 4Þ ScCl3 þ 4KBH4 ! KScðBH4 Þ4 þ 3KCl
ð20Þ
Surprisingly, powder X-ray diffraction data show no presence of NaCl or KCl in any of the ball-milled
samples of ScCl3–MBH4 (M = Na or K) in molar ratios 1:2, 1:3 or 1:4. This indicates that the mechanochemically induced reactions differ from reaction schemes Eqs. (19) and (20). Two sets of unidentified
Bragg peaks were observed in all the reaction products and one was assigned to MSc(BH4)4 and the
other to a new ternary sodium scandium chloride M3ScCl6. Furthermore, the samples with the starting
ratios ScCl3–MBH4 of 1:3 and 1:4 contain different amounts of MBH4 whereas no diffraction from ScCl3
was observed. The samples with the starting ratio of 1:2 show neither ScCl3 nor MBH4 peaks and appear to contain the largest fraction of the new compounds, NaSc(BH4)4 and KSc(BH4)4 for M = Na or K,
respectively [213,214]. This suggests that an addition reaction is responsible for the formation of the
ternary salts, Na3ScCl6 and K3ScCl6 according to reaction scheme
ð1 : 3Þ ScCl3 þ 3NaCl ! Na3 ScCl6
ð21Þ
ð1 : 3Þ ScCl3 þ 3KCl ! K3 ScCl6
ð22Þ
In conclusion, the optimal ratio of reactants ScCl3–MBH4 (M = Na or K) for synthesis of MSc(BH4)4
turns out to be 1:2. The above-mentioned observations can be explained assuming that the formations
of the ternary salts, M3ScCl6 (Eqs. (21) and (22)) are much faster than the formations of the borohydrides, MSc(BH4)4 (Eqs. (19) and (20)). Hence, the overall reactions for the samples ScCl3–MBH4 in
1:2 ratio are described in Eqs. (23) and (24) as a sum of the strongly coupled reactions for formation
of MSc(BH4)4 and M3ScCl6. These mechanochemical syntheses lead to maximum borohydride yields of
22 wt% and 18 wt% for NaSc(BH4)4 and KSc(BH4)4, respectively [213,214].
ð1 : 2Þ 2ScCl3 þ 4NaBH4 ! NaScðBH4 Þ4 þ Na3 ScCl6
ð23Þ
ð1 : 2Þ 2ScCl3 þ 4KBH4 ! KScðBH4 Þ4 þ K3 ScCl6
ð24Þ
61
In contrast, for the system ScCl3–LiBH4, only one chemical reaction occurs during ball milling,
which produces LiSc(BH4)4 and LiCl [211,212].
Ionic-covalent type hydrides, i.e. complex hydrides, may also form solid solutions, e.g. ca. 10 mol%
lithium chloride may dissolve in solid lithium borohydride by ball milling as described in reaction
scheme [221,222].
xLiCl þ 1 xLiBH4 ! LiðBH4 Þ1x Clx
ð25Þ
On the other hand, Rietveld refinement of powder X-ray diffraction data reveals that Cl readily
substitutes for BH
4 in the structure of hexagonal h-LiBH4 upon prolonged heating, e.g. at 498 K for
1 h resulting in formation of h-Li(BH4)1xClx, x 0.42 [221]. This observation indicates that formation of such solid solutions is more efficiently facilitated by elevated temperatures rather than by high
pressures generated by mechanochemical methods. These types of anion substitution reactions also
take place for a range of other metal borohydride-metal halide systems both as a result of mechanochemical treatment and subsequent heating, e.g. LiBH4–LiBr [223], LiBH4–LiI [224], NaBH4–NaCl [225],
Ca(BH4)2–CaF2, Ca(BH4)2–CaCl2 and Ca(BH4)2–CaI2 [226,234]. Furthermore, substitution of BH
4 by Cl
can also take place in the novel metal borohydrides formed during the mechanochemical synthesis,
e.g. Rietveld refinement suggests incorporation of 42 mol% of Cl on one of the two BH
4 sites in Al3Li4(BH4)13 [220].
Generally, the smaller anion tends to dissolve in the compound containing the larger anion, and the
structure of the latter tends to be preserved in the obtained solid solution in accordance with the
observation of a CaI2-type trigonal solid solution, tri-Ca((BH4)1xIx)2, i.e. dissolution of Ca(BH4)2 in
CaI2. This trend follows the relative size of the anions, I > BH
4 > Br > Cl derived by a comparison
of the unit cell volumes for different inorganic salts [235,236]. This trend in anion substitution reactions can be interpreted as an increase in the lattice energy due to the clearly observed decrease in the
unit cell volume, i.e. a decrease in the average distance between the ions in the structure.
Furthermore, anion substitution can occur due to structural similarities between the two compounds, which may partly explain the observation of two solid solutions of mechanically treated
LiBH4 —LiI, i.e. dissolution of LiBH4 in LiI and dissolution of LiI in LiBH4. The two solid solutions of
LiBH4 —LiI, b–LiI and h-LiBH4 all adopt isostructural hexagonal structures. Upon heating, the two solid
solutions of LiBH4 —LiI merge into one [224]. Similarly, prolonged heating of NaBH4 —NaCl produces
two solid solutions while NaBH4 and NaCl both share the same rock salt structure type [225]. This contrasts the behaviour of LiBH4–LiCl system where there are no indications of any dissolution of LiBH4 in
the alkali halide salts [221,222].
Recently, cation substitution in a borohydride by ball milling was observed for the first time by formation of a Mg1xMnx(BH4)2 solid solution. This substitution is most likely related to the close structural similarity of Mn(BH4)2 to a-Mg(BH4)2 and it is interesting to note that this compound retains the
framework structure [237].
Recently, a new series of double-cation double-anion borohydride chlorides based on rare-earths
elements, LiM’(BH4)3Cl, M’ = La, Gd, and Ce, were discovered using combined mechano-chemical synthesis and heat treatment using M’Cl3–LiBH4 (1:3) mixtures [238,239]. The novel cubic compounds,
LiM(BH4)3Cl contains isolated tetranuclear anionic clusters [Ce4Cl4(BH4)12]4 with a distorted cubane
Ce4Cl4 core charge-balanced by Li+ cations. The Li+ ions are disordered and occupy 2/3 of the 12d
Wyckoff sites and DFT calculations indicates that LiCe(BH4)3Cl is stabilized by larger entropy rather
than smaller energy. The new compound LiM’(BH4)3Cl simultaneously carries moderate amounts of
hydrogen, which is released at relatively low temperatures, and are fast Li-ion conductors.
4.2. Synthesis of novel alane and metal alanates
Among the most attractive materials for reversible hydrogen storage is the light-weight aluminium
based hydrides due to their relatively high gravimetric hydrogen capacities and, importantly, hydrogen release and uptake at moderate conditions [2,240]. In particular, Ti-doped NaAlH4 has received
significant attention in the past decade [154]. Hydridoaluminates containing a negatively charged
complex anion are denoted alanates in the following whereas polymorphs of aluminium hydride,
62
AlH3 are denoted alane. Alanates synthesised by RMM in hydrogen gas was discussed in Section 3.4 of
this review paper.
For the most stable alane polymorph, a-AlH3 the dehydrogenation enthalpy is determined by bomb
calorimetry to be 7.6(6) kJ/(mol H2) and the total dehydrogenation entropy is found to be 30.0(4) J/
(Kmol a-AlH3) [241]. From these values and the entropies of Al and H2, the hydrogenation pressure
of Al at room temperature according to reaction scheme Eq. (26) can be estimated to be above 10 GPa.
Al þ 3=2H2 ! AlH3
ð26Þ
Therefore, the traditional synthesis of Al-based hydrides is based on wet-chemistry methods where a
coordinating solvent forms a kinetically-stabilized hydride under mild conditions that can be isolated
from the solution [242,243]. Subsequently, the intermediate complex metal hydride solvent complex
can be decomposed to recover the pure hydride. Recently, it was discovered that a reversible hydrogenation reaction using Ti-catalyzed Al powder and triethylenediamine (TEDA) in tetrahydrofuran
forms the alane adduct (AlH3TEDA) at low pressure (p(H2) = 2 MPa, RT) [244].
In contrast, during the past few years mechanochemical techniques at a variety of different conditions have been used with increasing frequency for synthesis of alanates or for formation of homogeneous alanate-additive samples, e.g. in argon or hydrogen atmosphere or at liquid nitrogen
temperatures (cryo-milling) [245].
Interestingly, alane, AlH3 can be prepared directly by cryo-milling LiAlH4 and AlCl3 at 77 K resulting
in a metathesis reaction according to reaction scheme Eq. (27) [188,246,247].
3LiAlH4 þ AlCl3 ! 4AlH3 þ 3LiCl
ð27Þ
The reaction also proceeds at RT, however it results in low yields of alane and a high content of Al, due
to the high milling energy at RT which causes AlH3 to decompose to Al and H2. Furthermore, increased
brittleness at low temperatures leads to reduced particle sizes, hence shortened diffusion paths. Recently it was found that the yield of alane might be further increased by using NaAlH4 instead of
LiAlH4 or AlBr3 instead of AlCl3. The same study also revealed that the relative amount of a- and a’AlH3 formed in reaction scheme Eq. (27) can be somewhat controlled by adding 2.4 mol% FeF3 to
the reactant mixture [248].
The thermodynamically unstable AlH3, which has a three-dimensional covalent network structure,
decomposes slightly above room temperature but may be kinetically stabilized by thin oxide layers
[249]. This contrasts the ionic sodium hydride, NaH which decomposes at T 700 K [1]. Mechanochemistry reveals an interesting example, namely that reaction between a stable and an unstable
compound may provide a new material with intermediate stability. In this case a two-step addition
reaction occurs according to reaction scheme Eqs. (28) and (29) and sodium alanate is formed built
from discrete [AlH4] complex anions and sodium counter cations [65,155].
3NaH þ AlH3 ! Na3 AlH6
ð28Þ
Na3 AlH6 þ 2AlH3 ! 3NaAlH4
ð29Þ
This approach for tailoring thermodynamic properties is also used for ‘destabilising’ hydrides by
formation of reactive hydride composites, e.g. 2LiBH4–MgH2 and LiBH4–Al [66,250–252].
Sodium alanate, NaAlH4 has moderate formation enthalpies, however the rehydrogenation of
NaAlH4 is kinetically hampered and hydrogen release and uptake in this material need to be catalyzed
by titanium [154]. The materials resulting from mechanochemical doping NaAlH4 with small amounts
of catalytic TiCl3 (<2 wt%) have kinetic and cycling properties that are closer to those required for a
practical hydrogen storage medium [172,253]. Doping through mechanical milling of NaAlH4 with
dopant precursors is an effective mean of charging the hydride with catalyst but also activates the
material through reduction of the average particle size [173,253].
Furthermore, mechanochemical treatment can promote decomposition of metal alanates with
addition of catalysts [113,254,255]. Balema et al. and Easton et al. investigated the decomposition
of LiAlH4 into Li3AlH6 and Al with release of hydrogen at RT during short-time ball milling with addition of catalytic amounts of TiCl3 (3 and 2 mol%, respectively) [254,255]. Milling times as short as
5 min proved effective for this transformation and according to PXD data LiAlH4 was completely trans-
63
Fig. 22. Ball collision effects on reaction mixtures as a function of ball impact energy.
formed, i.e. only Bragg peaks from Li3AlH6, Al and LiCl were observed [255]. Monitoring of the gas
pressure within the milling vial as a function of milling time is shown in Fig. 22. The high catalytic
activity of TiCl3 and other additives may be attributed to the in situ formation of nano/microcrystalline, e.g. Al–Ti phases, during ball milling [166,256,257]. Similar results was obtained by mechanochemical treatment of NaAlH4 with TiCl3, ZrCl4, FeCl2 and FeCl3 [113].
Also it has been demonstrated that ball milling alone to some extent improves the dehydrogenating kinetics of undoped metal alanates, e.g. LiAlH4 and NaAlH4 [258,259]. A study of LiAlH4 showed
that milling between 1.5 and 14 h promoted decomposition at 333 K lower than for as-received
LiAlH4. Furthermore, it was also shown that LiAlH4 milled for less than 14 h decomposes via a twostep reaction pathway (reaction schemes Eqs. (30) and (31)).
3LiAlH4 ! Li3 AlH6 þ 2Al þ 3H2
ð30Þ
Li3 AlH6 ! 3LiH þ Al þ 1:5H2
ð31Þ
It has also been demonstrated that mechanochemical synthesis is a convenient method for preparation of bialkali metal hydrides from alkali metal hydrides and alkali metal alanates, e.g. Na2LiAlH6, is
prepared by ball milling sodium hydride, lithium hydride, and sodium alanate in a stoichiometric
composition which reacts by an addition reaction according to reaction scheme Eq. (32) [157].
NaH þ LiH þ NaAlH4 ! Na2 LiAlH6
ð32Þ
Bialkali alanates such as K2LiAlH6 and K2NaAlH6 have been prepared by similar reactions [176,177].
Magnesium and calcium alanate can be readily prepared from sodium or lithium alanate and magnesium or calcium chloride, respectively via a metathesis reaction according to reaction scheme Eq.
(33) [161,176,179,260–262].
M0 Cl2 þ 2MAlH4 ! M 0 ðAlH4 Þ2 þ 2MClðM0 ¼ Mg or Ca; M ¼ Li or NaÞ
ð33Þ
Interestingly, if the amount of lithium alanate in the reaction mixture is increased so that a composition of MgCl2:LiAlH4 1:3 is used, lithium magnesium alanate, LiMg(AlH4)3 forms according to
reaction scheme Eq. (34) [261].
MgCl2 þ 3LiAlH4 ! LiMgðAlH4 Þ3 þ 2LiCl
ð34Þ
Recently, also Sr(AlH4)2 and lithium beryllium hydrides LinBemHn+2m have been prepared mechanochemically [263,264]. Furthermore, rare earth metal chlorides have been ball milled with sodium
alanate and partial decomposition was observed due to a metathesis reaction and release of hydrogen
3
to form the more stable hexahydridoaluminate complex ion, AlH6 , according to reaction scheme
Eq. (35), which is suggested for R = La, Pr, Ce, Nd [265].
64
RCl3 þ 3NaAlH4 ! RðAlH4 Þ3 þ 3NaCl ! RAlH6 þ 2Al þ H2
ð35Þ
4.3. Novel quaternary hydrides
Recently several novel quaternary hydrides based on the amide anion, NH
2 and alanates, AlH4 or
borohydrides, BH
or
BH
have
been
prepared.
This
class
of
materials
is
of
high
interest for hydrogen
3
4
storage, since it combines both high hydrogen densities and storage properties of two types of storage
materials. A range of novel materials can be synthesized by mechanical milling utilizing several different types of reactions.
4.3.1. Metal borohydride amides
Formation of a quaternary hydride having the approximate composition Li3BH4(NH2)2, has been
achieved mechanochemically from LiBH4 LiNH2 in molar ratio 2:1 according to the addition reaction
shown in reaction scheme Eq. (36) [266,267].
2LiNH2 þ LiBH4 ! Li3 BH4 ðNH2 Þ2
ð36Þ
Pinkerton et al. also reported on the extent of the reaction as a function of the milling time. This
showed that a substantial amount of Li3BH4(NH2)2 is formed after 40 min of milling and that the reaction is completed after 300 min of milling. Milling for a total of 960 min caused no further changes in
the reaction product suggesting that the milling does not produce an amorphous phase.
Noritake et al. found that varying the LiNH2:LiBH4 molar ratio i.e. using 1:1, 2:1 and 3:1 yielded
products of different compositions according to reaction schemes Eqs. (36)–(38) [268].
LiNH2 þ LiBH4 ! Li2 ðBH4 ÞNH2
ð37Þ
3LiNH2 þ LiBH4 ! Li4 ðBH4 ÞðNH2 Þ3
ð38Þ
BH
4
NH
2
The
and
anions are preserved during the mechanochemical synthesis and in the structure of
Li4(BH4)(NH2)3 they are positioned on specific crystallographic sites, i.e. an ordered structure is formed
during ball milling as opposed to formation of a solid solution. The structures of Li2(BH4)NH2 and
Li3BH4(NH2)2 remain to be solved hence the precise stoichiometry of these compounds is not known.
However, the compositions of the ball-milled mixtures and the decomposition reactions suggest the
compositions stated in reaction schemes Eqs. (36) and (37).
Furthermore, a phase diagram study was carried out by Meisner et al. in which samples of LiNH2–
LiBH4 in a wide range of molar ratios of x = 0.33–0.80 were prepared by ball milling [269]. The study
shows that samples of x = 0.6–0.8 yield the cubic Li4BH4(NH2)3, while x < 0.5 yield another cubic compound, presumably Li3(BH4)(NH2)3. Furthermore, the study also implies existence of two other new
compounds, denoted c and d at x = 0.6 and 0.8, respectively. These compounds might be metastable
arising from the highly non-equilibrium mechanochemical process.
Several metal amidoboranes, M(NH2BH3)x have also been prepared by mechanochemical synthesis
from amidoborane, NH3BH3 and metal hydrides, e.g. LiNH2BH3, NaNH2BH3 and Ca(NH2BH3)2 according
to reaction scheme Eqs. (39) and (40) [270,271].
NH3 BH3 ðsÞ þ MHðsÞ ! MNH2 BH3 ðsÞ þ H2 ðgÞðM ¼ Li; NaÞ
ð39Þ
2NH3 BH3 ðsÞ þ CaH2 ðsÞ ! CaðNH2 BH3 Þ2 ðsÞ þ 2H2 ðgÞ
ð40Þ
Recently, LiNH2BH3 has been used as starting material for synthesis of Y(NH2BH3)3 by a metathesis
reaction according to reaction scheme Eq. (41) [272].
YCl3 þ 3LiNH2 BH3 ! 3LiCl þ YðNH2 BH3 Þ3
ð41Þ
In the same study similar syntheses were tested utilizing different combinations of MNH2BH3 (M = Li,
Na) and YX3 (X = F, Cl), however the reactions seem either to have a very low efficiency or the desired
reaction product forms as an amorphous phase.
65
4.3.2. Metal alanate amides
In general, ball milling mixtures of alkali amides, MNH2 and alkali alanates, MAlH4 (M = Li or Na)
result in the release of hydrogen caused by the occurring multistep reactions [273–276]. Furthermore,
3
the presence of amide might facilitate transformation of the tetrahedral AlH4 into octahedral AlH6 by
ball milling at ambient conditions. Dolotko et al. observed formation of Li3AlH6 already after 4 min of
milling of a LiAlH4–LiNH2 sample in molar ratio 1:1 [274]. Xiong et al. report release of approximate
one equivalent H2 as all LiAlH4 is transformed to Li3AlH6 according to reaction scheme Eq. (42) [273].
3LiAlH4 ! Li3 AlH6 þ 2Al þ 3H2
ð42Þ
In comparison negligible amounts of hydrogen were released from LiAlH4 ball milled for 36 h [273].
Prolonged milling results in formation of amorphous compounds and from combined MAS NMR
and PXD analysis the reactions shown in reaction scheme 43 and 44 are suggested [274].
2LiAlH4 þ LiNH2 ! Li3 AlH6 þ AlN þ 2H2
ð43Þ
Li3 AlH6 þ LiNH2 ! 4LiH þ AlN þ 2H2
ð44Þ
A similar sample of NaAlH4–NaNH2 in molar ratio 1:1 was found to react at a much slower rate and
requires 60 min of milling to reach its completion. The occurring reactions are similar to the reactions
observed for LiAlH4–LiNH2 also yielding Na3AlH6 as an intermediate [274].
Similar studies have been performed for samples of LiAlH4–NaNH2 (molar ratio 1:2, 1:1 and 2:1)
and NaAlH4–LiNH2 (molar ratio 1:1) [274–276]. These studies all showed release of two equivalents
H2 during milling. For samples containing LiAlH4–NaNH2 the first reaction step is a cation-exchange
reaction according to reaction scheme Eq. (45) starting already after 1 min of milling.
3LiAlH4 þ 4NaNH2 ! 3NaAlH4 þ Li3 NaðNH2 Þ4
ð45Þ
Upon further milling of LiAlH4–NaNH2 and NaAlH4–LiNH2 in molar ratio 1:1, LiNa2AlH6 is formed as an
intermediate decomposition product [274]. This formation is facilitated by the presence of the amide,
3
since neither of the pure alanates nor the mixture of the two could transform to AlH6 containing compounds by ball milling alone [277,278]. The amides and alanates present in the sample at this stage are
slowly consumed leading to formation of Al, LiH, NaH and an unknown compound, possibly AlN or
LiAl0.33NH [274,275]. After only 30 min of milling no further changes are observed.
Furthermore, Dolotko et al. reported formation of a solid solution, Na3xLixAlH6 in the NaAlH4–
LiNH2 sample according to reaction scheme Eq. (46).
NaAlH4 þ ð2 xÞNaH þ xLiH ! Na3x Lix AlH6
ð46Þ
The samples of LiAlH4–NaNH2 in molar ratio (1:2) and (2:1) exhibit different reaction pathways during
the prolonged milling due to the excess of NaNH2 or LiAlH4, respectively [276]. The overall reactions
suggested to occur during ball milling for the (1:2) and (2:1) samples are shown in reaction schemes
Eqs. (47) and (48), respectively.
LiAlH4 þ 2NaNH2 ! 2NaH þ LiAlN2 H2 þ 2H2
ð47Þ
2LiAlH4 þ NaNH2 ! NaAlH4 þ Li2 AlNH2 þ 2H2
ð48Þ
Apparently, up to now, a mixed borohydride–alanate compound has not been prepared. The system
LiBH4 —NaAlH4 was investigated mechanochemically and by hand-mixing, the latter treatment did
not lead to any reactions while a metathesis reaction was observed for the ball milled sample, see
Eq. (49) [279,280].
LiBH4 ðsÞ þ NaAlH4 ðsÞ ! LiAlH4 ðsÞ þ NaBH4 ðsÞ
ð49Þ
4.4. Solid–liquid mechanically assisted synthesis
In some cases addition of a solvent during mechanochemical synthesis is a necessity to promote
the desired reaction, e.g. Ca(AlH4)2 can be prepared by a so-called mechanically assisted synthesis
66
from NaAlH4 and CaCl2 mediated by tetrahydrofuran (THF) according to reaction scheme Eq. (50)
[281].
2NaAlH4 þ CaCl2 þ 4THF ! CaðAlH4 Þ2 4THF þ 2NaCl
ð50Þ
The synthesis reported by Fichtner et al. was performed in a glass ball mill reactor to create fresh particle surfaces. The mixture was heated under reflux and the grey solid precipitate was milled by the
glass balls. The high yield of the solvated product (84%) was attributed to the high dissolution capacity
for NaAlH4 of THF and that the adduct might form more easily compared to CaAlH4. After the synthesis
the solvent can be removed almost completely from freshly prepared Ca(AlH4)24THF under vacuum
yielding a fine white powder of Ca(AlH4)2.
5. Final remarks and conclusions
Recent progress in the field of hydrogen storage materials has received extensive support from
mechanochemistry methods. Mechanical milling has been widely used not only to tune metal microstructures for modifying their hydrogenation properties but also as an efficient tool for the synthesis
of hydrogen storage materials. Solid/solid, solid/liquid and solid/gas reactions can be activated by
mechanochemistry.
Mechanical milling acts as a combination of compression and shear on the powder between two
colliding balls and between ball and container wall. Upon impact, the powder particles trapped between them will first experience elastic deformations, which are reversible in the elastic region
(Fig. 22). If the load increases, the material enters into plastic region and irreversible deformations occur, which may be followed by breakage of the material. Therefore, collisions result in impulses of
compression and shear that generate plastic deformation and fracture of the particles. Most of the energy transferred to the powder is mainly used for creating new surfaces and concomitant particle
refining.
Very similar mechanical effects are produced by Severe Plastic Deformation (SPD) techniques such
as Equal Channel Angular Pressing (ECAP), Cold Rolling (CR), and High Pressure Torsion (HPT). Therefore, it could be expected that SPD will have similar impact on hydrogen storage behaviour than
mechanical milling.
In the particular case of solid/solid reactions, while comminution is an important result of milling,
agglomeration assisted by cold-welding processes and thereby the formation of active interfaces is
more relevant. Solid-state reactions likely happen at the interface between solid particles of the different reactants. The increase of the surface-area to bulk-volume ratio due to mechanically induced
decreasing particle size increases the interfacial contact area between reactive compounds and therefore also the rate of reaction. In addition, plastic deformations due to mechanical work create defects
providing fast atomic diffusion pathways and contributing to the increased reaction rate. The intense
mixing action provided by the random motion of milled powders and repeated particle fracture and
welding favours the formation of active interfaces and ensures the chemical homogeneity of the final
product at long milling time.
As concerns solid/gas reactions, particle comminution and related effects such as fresh surface generation and diffusion path reduction are determinant. Thus, Mechanical Milling under hydrogen gas is
characterized by fast formation of hydride compounds under moderate pressure and temperature. For
instance, several days are required for the synthesis of Mg2CoH5 and Mg2FeH6 hydrides by sintering
methods at temperatures as high as 750 K, whereas the reaction takes place in only 3 h by reactive ball
milling under the same hydrogen pressure (9 MPa) [116,139]. Embrittlement due to hydrogen absorption favours particle refinement and fast synthesis of binary metal hydrides. From this point of view,
milling energy (determined by milling process parameters such as rotation speed and ball-to-powder
mass ratio) and mechanical properties of reactants (toughness, fracture limit) must play a key role. For
the synthesis of ternary hydrides starting from elemental powders, the above-mentioned mechanisms
of cold-welding and interface diffusion of solid reactants should be also considered.
The feasibility of hydride formation in solid/gas reactions is governed by thermodynamics (i.e. hydride stability under external pressure and temperature conditions). Thus, the formation of highly sta-
67
ble hydrides such as TiH2, ZrH2 and MgH2 is straightforward. In contrast, the synthesis of hydrides that
are reversible near room temperature, such as LaNi5 and NaAlH4 requires high pressure for the hydride
to be stable at the temperatures reached during the milling process [121]. If the pressure is not high enough, only the a-solid solution LaNi5H0.15 or the more stable Na3AlH6 phase are observed, respectively.
In this context, the failure to obtain the LiAlH4 or ternary alkaline hexa-alanates by reactive ball milling
is not surprising since they are not stable at the usual operating pressure and temperature [121].
Therefore, in reactive ball-milling experiments, both macroscopic and local pressures as well as
temperatures have to be considered in detail. As concerns pressure, one can distinguish the mechanical pressure in the material trapped between two colliding steel balls (internal mechanical pressure)
and the gas vial pressure (external isostatic pressure). The gas pressure is not significantly affected by
milling due to the high compressibility of the gas phase. Most modern equipment allows for gas pressures up to 15 MPa. In contrast, the internal pressure of the material at mechanical impact may well
reach some GPa and is not isotropic. This would explain for instance the formation of the high pressure
c-MgH2 phase by mechanical milling, which otherwise occurs in anvil cells above 2 GPa [282]. Formation of c-MgH2 phase can be obtained by mechanical milling of thermodynamically stable b-MgH2
phase under argon atmosphere, which demonstrates that its formation is related to the internal
mechanical pressure.
As concerns the temperature, one can also distinguish between the temperature of the material
trapped between the milling tools (material temperature at impact) and the macroscopic material
temperature. For materials with high thermal conductivity (i.e. metals) or by milling under hydrogen
gas (which also offers high thermal conductivity), the macroscopic material temperature is not expected to differ much from the vial temperature. The temperature increase in the vial does not exceed
some tens of degrees and temperatures in the range 320–350 K have been monitored by temperature
gauges. This temperature increase is however non-negligible for hydrogen storage systems that are
reversible near normal conditions of pressure and temperature. Several methods can be followed to
circumvent this problem such as working at higher hydrogen pressures and minimizing temperature
increase by using short milling times (below 10 min). Another and more elegant alternative approach is working at low temperatures, i.e. cryo-milling, though the temperature of the system has
to be kept high enough to allow for hydrogen mobility in the bulk powder material.
Mechanical milling of powders can be performed under other reactive gases such as diborane and
ammonia or in a liquid medium such as THF for the synthesis of hydrogen storage materials. Once
again fresh surface generation induced by mechanical work is likely to be determinant to promote solid/gas and solid/liquid reactions. Contrary to reactive mechanical milling under hydrogen gas, these
preparation methods are only recently explored in the literature. Further progress is still needed to
understand the involved reaction mechanisms and the feasibility of compound formation by these novel routes. A wide research field remains open for the production of new hydrogen storage systems
with undiscovered hydrogenation properties.
Acknowledgments
We thank V. Balema for fruitful discussions regarding this review. The work was supported in part
by the Danish National Research Foundation (Center for Materials Crystallography), the Danish Strategic Research Council (Center for Energy Materials and the HyFillFast project), and by the Danish Research Council for Nature and Universe (Danscatt). We are grateful to the Carlsberg Foundation. J.H.
would like to thank the Natural Science and Engineering Council of Canada and also the Research
Council of Norway for additional funding that permitted a sabbatical leave at the Institute for Energy
Technology (IFE) in Norway. ML, FC and JZ would like to thank CNRS and the French agency ANR for
financial supports trough research programs ALHAMO and NANOHYDLI.
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Mechanochemical synthesis of hydrogen storage materials

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