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Study of the diffusion kinetics and mechanism ofelectrochemical hydriding of MgeNieMm alloys

D. Vojt�ech*, V. Knotek

Department of Metals and Corrosion Engineering, Institute of Chemical Technology, Prague, Technicka 5, 166 28 Prague 6, Czech Republic

a r t i c l e i n f o

Article history:

Received 14 October 2010

Received in revised form

8 February 2011

Accepted 11 February 2011

Available online 12 March 2011

Keywords:

Magnesium

Hydrogen storage

Nickel

Rare earths

Diffusion

* Corresponding author. Tel.: þ420 220444290E-mail address: Dalibor.Vojtech@vscht.cz

0360-3199/$ e see front matter Copyright ªdoi:10.1016/j.ijhydene.2011.02.063

a b s t r a c t

Electrochemical hydriding in a 6 M KOH solution at 20 and 80 �C for 480 min was applied on

a series of as-cast binary MgeNi and ternary MgeNieMm alloys (Mm ¼ mischmetal con-

taining 45% Ce, 38% La, 12% Nd and 4% Pr) containing 11e24 wt. % Ni and 0e6 wt. % Mm.

The kinetics and mechanism of the hydriding process, as well as hydrogen release

temperatures, were studied by glow discharge spectrometry hydrogen profiling, scanning

electron microscopy, energy dispersion analysis, X-ray diffraction and mass spectrometry.

A maximum hydrogen concentration of 1.1% was achieved in the eutectic MgNi24Mm5

alloy hydrided at 80 �C. In all cases, the main hydriding product was binary MgH2 hydride.

Mass spectrometry revealed its destabilization due to Ni and Mm because its decomposi-

tion temperature was lowered by about 100 �C. Both nickel and mischmetal showed

positive effects on hydriding and dehydriding kinetics. These effects are discussed in

relation to the hydriding mechanism, electronic structure and atomic size of additives and

structural variations of the alloys. Based on the H-concentration profiles, the diffusion

coefficients of hydrogen were estimated. For the eutectic MgNi24Mm5 alloy, the H diffusion

coefficient at 20 �C was 4 $ 10�10 cm2 s�1.

Copyright ª 2011, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights

reserved.

1. Introduction systems have been studied so far, but only some of them have

Hydrogen has attracted much interest as a pure fuel and

energy source for both mobile and stationary applications. Its

combustion produces only water, which is neither an atmo-

spheric pollutant nor a cause of global atmospheric warming.

Great effort has been devoted to find a safe and cost-effective

method for hydrogen storage. Among these methods,

hydrogen storage in the form of metallic hydrides is prom-

ising. Suitable metallic hydrides are stable at room tempera-

ture and can be stored for an extended period or transported

across a great distance. In a vehicle, power station or other

energy-consuming device, they are heated and decomposed

to release hydrogen at a defined rate. Variousmetallic hydride

; fax: þ420 220444400.(D. Vojt�ech).2011, Hydrogen Energy P

the potential for further commercial utilization. In particular,

magnesium hydride MgH2 has been the subject of extensive

research all over the world because it has one of the largest

hydrogen gravimetric densities, 7.6 wt. %, exceeding the

Department of Energy’s (DoE) requirement of 6 wt. %.

However, the well known drawback of this hydride is its high

thermodynamic stability associated with slow hydriding/

dehydriding kinetics and high hydrogen release temperature.

To destabilize the MgH2 phase, various attempts have been

made. These include the addition of transition metals (Ni, Cr,

Co, Fe, Ti, rare earths (RE) etc.) that form either binary or more

complex hydrides, such as Mg2NiH4, Mg2FeH6, Mg2CoH5,

Mg3MnH7, Mg3CrH6, Mg2RENiH7, REH3 and others [1e4]. The

ublications, LLC. Published by Elsevier Ltd. All rights reserved.

Table 1 e Chemical compositions (in wt. %) of hydridedmagnesium alloys (Mm-mischmetal).

Alloy designation Element

Ni Mm

MgNi11 10.9 e

MgNi11Mm4 11.1 4.0

MgNi11Mm6 10.8 6.1

MgNi15 14.8 e

MgNi15Mm4 15.4 3.7

MgNi24Mm5 23.9 5.0

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more complex hydrides usually decompose at lower temper-

atures than does the MgH2 phase. However, one disadvantage

of these hydrides is their lower hydrogen gravimetric density

in comparison with pure MgH2. For the improvement of the

hydriding/dehydriding behavior of MgH2, catalysts are

applied. These include RE, transition metal oxides, carbon,

halides etc. Some catalysts do not necessarily form special

hydrides. Although the exact mechanisms of their actions are

not presently known, it is believed that they may involve the

surface modification of magnesium, destruction of the MgO

surface layer, enhancement of the formation of atomic

hydrogen on the surface, acceleration of hydrogen diffusion

and so on [5]. Another approach to MgH2 destabilization is the

preparation of amorphous and/or nano-crystalline structure

of the powdered hydride through intensive ball milling, rapid

solidification and other techniques. When hydriding pure Mg,

a surface MgH2 layer forms, and after it achieves a certain

critical thickness, further inward diffusion of hydrogen is

suppressed because H-diffusivity in MgH2 is much slower

than in Mg [5]. The advantage of ball milling consists of

continuous mechanical destruction of the MgH2 layer present

on each powder particle that creates paths for hydrogen

penetration. Moreover, intensive mechanical forces during

milling induce high concentrations of lattice defects, grain

boundaries and interfaces that also represent good paths for

hydrogen diffusion. In an extreme case, sufficiently long and

intensive milling can form an amorphous powder structure

that contains a larger free volume for H interstitials as

compared to an ordered crystal. Rapid solidification tech-

niques, such as melt spinning or gas atomization, provide

beneficial structural features similar to those provided by ball

milling.

In this work, we focus our attention on the MgeNieMm

system (Mm ¼ mischmetal, i.e., an alloy of RE like Ce, La, Nd

and Pr). MgeNi alloys have been investigated in the context of

H storage many times, and nickel has been clearly identified

as an element supporting hydriding/dehydriding [6]. REs have

been known to act as both hydride formers and catalysts of

hydriding [7]. RE-containing hydrides include REH3, Mg2RE-

NiH7, REH2, Mg3REH9 and others. Due to the large atomic sizes

of RE, they enhance the glass-forming ability (GFA) of Mg

alloys and consequently help the formation of beneficial

amorphous structures during rapid solidification [8].

The preparation of Mg-based hydrides generally involves

intensive long-time milling of an original alloy in a hydrogen

atmosphere at high temperatures and pressures. An alterna-

tive to the elemental synthesis is electrochemical hydriding.

In this process, atomic hydrogen is formed during electrolysis

of a suitable water solution on the cathode surface. When the

cathode is made of a Mg alloy, hydrogen directly enters the

cathode and forms hydride phases in its structure. We have

demonstrated that electrochemical hydriding can achieve

hydrogen concentrations exceeding 1 wt.% [9,10] that

approach the compositions of some of the complex hydrides

prepared by the classical and expensive route at high

temperatures and pressures. Electrochemical hydriding does

not need gaseous hydrogen, high temperatures and pressures.

Although electrochemical hydriding has been investigated

and utilized for a long time, there are still unanswered ques-

tions about its mechanism. Our study is, therefore, devoted to

the electrochemical hydriding of MgeNieMm alloys with

particular attention to the diffusion kinetics of hydrogen and

hydriding mechanism.

2. Experimental

In our work, several MgeNieMm-based alloys (Mm ¼ mis-

chmetal containing 45% Ce, 38% La, 12% Nd and 4% Pr), see

Table 1, were hydrided by an electrochemical process (here-

after, all concentrations are in wt.%). As indicated before,

MgeNi alloys have been extensively studied in the past

decades as prospective hydrogen storage materials because

nickel is known to significantly destabilize magnesium hyd-

ride and decrease its decomposition temperatures. Another

additive in this experiment was Mm because REs are also

known to positively affect hydriding behavior of MgeNi alloys

in gaseous hydrogen.

The alloys were prepared by vacuum induction melting of

pure Mg, Ni and Mm (99.9% purity) under argon. The cylin-

drical ingots of 200mm in length and 20mm in diameter were

gravity-cast into a brass mold. Afterward, the ingots were cut

into 0.5-mm thick coupons for electrochemical hydriding.

Prior to hydriding, the surface of the coupons was mechan-

ically polished.

Electrochemical hydriding was carried out in a 6mol/l KOH

at 20 �C and 80 �C. The alloys were immersed in an electrolyte,

connected to a DC source and polarized as the cathode, while

a graphite rod of 10mm in diameter and 100mm in lengthwas

used as the anode. Current density during hydriding and

hydriding time were 100 A/m2 and 480 min, respectively. The

current density value was selected to ensure a sufficient

amount of atomic hydrogen on the cathode surface was

formed and to prevent excessive evolution of gaseous

hydrogen. In part of the hydriding experiments, 0.02e12 mg/l

of As3þ was added to the electrolyte. Arsenic is well known as

“a poison of hydrogen recombination”, i.e., it reduces the rate

of H þ H / H2 recombination on the cathode surface; there-

fore, it prolongs the lifetime of atomic hydrogen. We assumed

that the hydrogen concentration achieved by hydriding could

be increased by addition of As.

The structure of as-cast alloys was investigated by scan-

ning electron microscopy (SEM, Hitachi S 4700) and energy

dispersion spectrometry (EDS, Noran). Phase compositions

both before and after hydriding were determined by X-ray

diffraction (XRD, X Pert Pro). To measure the concentration

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profiles of hydrogen, glow discharge spectrometry (GDS,

Profiler 2) was employed. Due to the fact that hydriding was

performed in a strongly alkaline bath, the formation of

magnesium hydroxide and/or complex hydroxide surface

layers could be expected. However, the influence of such

layers on the hydrogen profile analysis should be minimized,

and only hydrogen present in the metallic phase should be

measured. For this reason, the oxygen concentration profile

was also analyzed to determine the exact position of the

hydroxide/metal interface. This position was taken as the

point at which oxygen intensity (concentration) reduces to

below the detection limit (see Fig. 1). The GDS analyzer was

calibrated with respect to Mg(OH)2 prepared by anodic

oxidation of pure Mg. Before calibration, Mg(OH)2 was verified

by XRD. The sputtering rate during depth profiling was

calculated from the surface topography measured after anal-

ysis by a surface profilometer.

The hydrogen released from the hydrided alloys was

monitored by mass spectrometry (MS, Setaram SETSYS

Evolution - 1750 temperature range of 50e400 �C, heating rate

of 5 �C/min, Ar atmosphere). The mass spectrometer was

adjusted to detect the masses of H2 and OH fragments and

distinguish between the evolution of water vapor and hyd-

rogen upon heating.

3. Results

3.1. Structures of alloys

Figs. 2 and 3 show SEM micrographs of MgNi11, MgNi11Mm6,

MgNi15Mm4 and MgNi24Mm5 alloys and selected X-ray

elemental maps. Both MgeNi and MgeRE systems are of the

eutectic type. On the Mg side of the MgeNi and MgeRE (Ce, La,

Nd, Pr) phase diagrams, there are eutectics composed of a-Mg

and a corresponding intermetallic phase: Mg2Ni and Mg12Mm.

The latter one refers to a solid solution of isostructuralMg12Ce,

Mg12La,Mg12Nd andMg12Pr phases (space group I4/mmm) [11].

Eutectic concentrations of Ni, Ce, La, Nd and Pr are 23.5, 20.6,

11.9, 18.2 and22.0%, respectively [11]. TheMgNi11alloy (Fig. 2a)

hasa typicalhypoeutectic structureconsistingofprimarya-Mg

dendrites (dark) and a-Mg þ Mg2Ni eutectic (light) in the

Fig. 1 e Analysis of hydrogen concentration profile by GDS.

The surface hydroxide layer indicated by an increased

oxygen concentration was excluded from analysis. Only

hydrogen dissolved in the metallic phase was measured.

interdendritic region. Similar structures are observed for the

MgNi11Mm6, MgNi15Mm4 (Fig. 2b, c and 3a). In the interden-

dritic regions of these alloys, there is a ternary eutectic

a-Mg þ Mg2Ni þ Mg12Mm, as was also observed in the XRD

patterns. Quite a different structure is observed for the

MgNi24Mm5 alloy shown in Figs. 2d and 3b. Due to the highNi-

concentration in this alloy approaching the eutectic point, the

structure is dominated by the ternary Mg þ Mg2Ni þ Mg12Mm

eutectic (Fig. 6). The detailed X-ray map in Fig. 3b shows that

the Mg2Ni phase forms relatively coarse, elongated and inter-

connected particles, while the Mm-rich one is represented by

fineand separatedplateson theupper-left sideof themap.The

occurrence or even predominance of dispersed eutectic

structures in the alloys was one factor determining their

selection in our experiment. In eutectics, there are high frac-

tions of interphase boundaries where hydrogen diffuses fast;

therefore, they can be expected to support hydriding of alloys.

3.2. Hydrogen concentration profiles

Concentration profiles of hydrogen in the metallic phase

beneath the surfacehydroxide layer arepresented in Fig. 4. The

profiles are shown separately to better observe the effects of

the experimental variables, i.e., Ni- and Mm-concentrations,

temperature and As concentration in the bath. It is observed

that all profiles show a typical diffusion shape, i.e., maximum

concentration on the surface and a progressive decrease

toward the alloy interior. At some depths, hydrogen concen-

trations fall below the GDS detection limit. In the following

text, these depths will be referred to as hydrogen penetration

depths. Fig. 4a compares the MgNi11 and MgNi15 alloys. It is

evident that nickel positively affects both the maximum

hydrogen concentration and hydrogen penetration depth. The

influence of mischmetal additions is demonstrated in Fig. 4b.

According to this figure, Mm appears to be very effective in

supporting hydriding because the original maximum H

concentration of 0.3% for the MgNi11 alloy increases to more

than twice that value to 0.8% by adding 4% Mm. Furthermore,

2%Mmdoesnot increase themaximumHconcentrationonthe

surface significantly but prolongs the hydrogen penetration

depth. The influence of both Mm and Ni are illustrated in

Fig. 4c, where MgNi11Mm6 and MgNi24Mm5 alloys are

compared. One can see that the maximum H concentrations

are almost identical. However, an almost twofold incre-

ase in the hydrogen penetration depth is observed for the

MgNi24Mm5 alloy in comparison with its Ni-depleted coun-

terpart. A comparison of Fig. 4aec reveals that the eutectic

MgNi24Mm5 alloy shows the fastest hydriding due to the

largest hydrogen maximum concentration (0.8%) and pene-

trationdepth (50 mm). Fig. 4d presents the effect of temperature

on the hydriding process of this alloy. As expected, the

temperature increase from 20 �C to 80 �C accelerates hydriding

because the maximum hydrogen concentration and penetra-

tion depth grow up to 1.1% and 60 mm, respectively. The influ-

ence of addition of As is illustrated in Fig. 4e. Surprisingly,

arsenic does not accelerate hydriding; instead, it seems that its

effect is slightly negative because it reduces both the

maximum concentration and penetration depth to some

extent. However, no systematic dependence of these param-

eters on As concentration is observed.

Fig. 2 e SEM micrographs of MgNi11 (a), MgNi11Mm6 (b), MgNi15Mm4 (c) and MgNi24Mm5 (d) alloys.

Fig. 3 e X-ray elemental maps of the MgNi11Mm6 (a) and MgNi24Mm5 (b) alloys (EDS).

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Fig. 4 e Hydrogen concentration profiles in the hydrided alloys: a) MgNi11 and MgNi15 at 20 �C, b) MgNi11, MgNi11Mm4 and

MgNi11Mm6 at 20 �C, c) MgNi11Mm6 and MgNi24Mm5 at 20 �C, d) MgNi24Mm5 at 20 and 80 �C, e) MgNi24Mm5 in the As-

containing baths at 20 �C (GDS).

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Hydrogen penetration depths estimated from the concen-

tration profiles in Fig. 4 are summarized in a three-dimen-

sional graph in Fig. 5. One can see that a minimum

penetration depth of 10 mm progressively increases with

increasing amounts of both Ni and Mm. As will be discussed

later, this behavior is attributable to both the catalytic effects

of additives and the formation of disperse eutectic. For this

reason, the maximum penetration depth is observed for the

eutectic MgNi24Mm5 alloy.

3.3. Phase composition after hydriding

To reveal the hydriding mechanisms, all alloys were analyzed

by XRD both before and after electrochemical hydriding.

Because the best hydriding behavior was observed for the

MgNi24Mm5 alloy with pure eutectic structure (Fig. 2d), phase

transformations during electrochemical hydriding are illus-

trated for this alloy. The same behavior was observed for the

other hydrided alloys (not shown). Fig. 6 shows XRD patterns

of the eutectic MgNi24Mm5 alloy before and after hydriding at

80 �C/480 min. As was already indicated before, the as-cast

alloy contains three phases in its structure: a-Mg, Mg2Ni and

Mg12Mm. At first sight, XRD pattern of the hydrided alloy

seems almost identical. However, a detailed investigation of

this pattern reveals the presence of two new phases after

hydriding, namely Mg(OH)2 andMgH2. Magnesium hydride, as

the main hydriding product, is characterized by the highest

peak at a diffraction angle of 28�. Surprisingly, no ternary,

Fig. 5 e Hydrogen penetration depths versus Ni- and Mm-

concentrations in the alloys (GDS).

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more complex or RE-based hydrides were detected by XRD,

suggesting their volume fractions were below about 5%. XRD

patterns of other hydrided alloys (not shown) confirm that

MgH2 is the only hydriding product of electrochemical

hydriding of both the binary MgeNi and ternary MgeNieMm

alloys investigated in this study. If we take the MgH2 peak

Fig. 6 e XRD patterns of the as-cast (a) and hydrided (b)

MgNi24Mm5 alloy. Hydriding was performed at 80 �C/480 min. The only hydriding product is MgH2.

height as an approximate measure of its volume fraction, the

highest volume fraction of hydride is found in the eutectic

MgNi24Mm5 alloy, as is also documented by H concentration

profiles in Fig. 4.

3.4. Hydride decomposition

Hydrogen evolution upon heating was monitored by mass

spectrometry. The mass spectrometer detected H2 and OH

particles. The former originated from both hydrogen and

water. We assume that water vapor forms by a decomposition

of hydroxides present on the alloy surface due to the strongly

alkaline solution used for hydriding. To distinguish between

the evolution of hydrogen and water, detection of both types

of particles was necessary. In Fig. 7, the results of MS are

illustrated for the eutectic MgNi24Mm5 alloy hydrided at

80 �C/480 min. Upon heating of this alloy, H2 starts to be

released at about 200 �C, and its evolution continues up to

about 400 �C. In contrast, the evolution of OH originating from

the hydroxide surface layer starts at higher temperatures,

about 300 �C. This suggests that hydrogen evolved at

200e300 �C originates from MgH2 hydride decomposition. At

300e400 �C, there are two hydrogen sources, MgH2 and

Mg(OH)2. The other hydrided alloys, both binary and ternary,

behave in a similar manner, i.e. magnesium hydride starts to

decompose at 200 �C. However, it is well known that pure

magnesium hydride starts to decompose at 300 �C [12]. It is

thus evident that this hydride is destabilized in the hydrided

MgeNi and MgeNieMm alloys significantly due to a reduction

of decomposition temperature by 100 �C.

4. Discussion

4.1. Mechanism of electrochemical hydriding

It is known that atomic hydrogen forms on the cathode

surface via an electrochemical reaction:

H2Oþ e�/HþOH� (1)

Fig. 7 e MS curves of H2 and OH fragments for the

MgNi24Mm5 alloy hydrided at 80 �C/480 min.

Table 2 e Crystallographic parameters of phases [14].

Phase Space group Unit cell parameters (A)

Mg P63/mmc a ¼ 3.24, c ¼ 5.26

MgH2 P42/mnm a ¼ 4.52, c ¼ 3.02

Mg2NiH0.3 P6222 a ¼ 5.25, c ¼ 13.43

Mg2NiH4 Fm3m a ¼ 6.47

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We have shown in our previous paper [10] that electro-

chemical hydriding does not work for the pure Mg, probably

due to the formation of a surface MgH2 layer preventing

further inward hydrogen diffusion. The diffusion coefficient of

H in MgH2 is much lower than that in Mg [5]. To support

hydriding, catalysts are required, and both Ni and RE are

known to show this effect. According to some authors [13],

these additives have unoccupied d-bands in their electronic

structure and, therefore, promote the formation of atomic

hydrogen, whose unpaired electrons may contribute to the d-

states. PureMg, in contrast, has fully occupied s-orbitals, so its

effect is negligible. Additionally, in the structure of both

binary MgeNi and ternary MgeNieMm alloys, there are

significant volume fractions of eutectic mixtures containing

dispersive Mg, Mg2Ni and Mg12Mm phases (Figs. 2 and 3).

Because nickel and rare earth elements catalyze hydriding, it

can be assumed that atomic hydrogen forms and preferen-

tially enters the eutectic structure. Inward hydrogen diffusion

in the eutectic structure is also accelerated by a large fraction

of boundaries between eutectic phases that represent good

paths for diffusion (Fig. 5). Due to the high nickel concentra-

tions in the hydrided alloys, in the eutectic there is a large

fraction of the Mg2Ni phase. Hydrogen then diffuses along

both the interphase boundaries and in this phase to form an

interstitial solid solution Mg2NiH0.3:

Mg2Niþ 0:3H/Mg2NiH0:3 (2)

Both Mg2Ni and Mg2NiH0.3 have hexagonal crystal lattices

(space group P6222). For this reason, these phases are not

distinguishable in the XRD patterns in Fig. 6. Finally hydrogen

diffusing in the eutectic regions preferentially reacts with Mg

to form MgH2:

Mg þ 2H/MgH2 (3)

By the mechanism suggested above, hydrogen is able to

penetrate relatively deeply into the material, which is neces-

sary to achieve high hydrogen gravimetric densities.

However, the question still remains: why MgH2 is formed

preferentially, i.e., why are other hydrides, for example, the

well known Mg2NiH4, not formed during electrochemical

hydriding of the MgeNi and MgeNieMm alloys? In these

alloys, one would expect additional hydriding reactions to

compete with that expressed by Eq. (3) [14]:

Mg2NiH0:3 þ 3:7H/Mg2NiH4 (4)

Mg12Mmþ 26H/12 MgH2 þMmH2 (5)

However, these reactions do not play an important role in

hydriding. Apparently, nucleation and growth kinetics of

MgH2 are faster than those of other hydrides. A similar finding

was recently reported by Denys et al [14]. who hydrided

MgeNieMm alloys and observed that the rate of MgH2

formation (Eq. (3)) is two orders of magnitude higher than that

of Mg2NiH4 formation (Eq. (4)). However, they did not provide

any explanation for this difference. Perhaps, the difference in

the nucleation rates of hydrides may be partially attributed to

the structural characteristics of the original phases and

resulting hydrides. The crystallographic parameters of Mg,

MgH2, Mg2NiH0.3, Mg2NiH4 phases are summarized in Table 2.

One can see that there is a similarity in the lattice parameters

between Mg and MgH2. In contrast, the lattice parameters of

both ternary hydrides differ significantly. Therefore, in the

case of MgeNieMm hydriding, a lower nucleation barrier for

the MgH2 hydride may be expected in comparison with that

for Mg2NiH4. By hydriding MgeAleTieFe-layered structures

and MgeMneNi alloys [15,16], it was shown that the best

hydriding behavior was also associated with the presence of

MgH2 as the main hydride phase. Additives mainly acted as

hydriding catalysts and diffusion accelerators.

4.2. Hydrogen diffusion

As indicated before, the hydrogen diffusion profiles in the

alloys shown in Fig. 4 result from two processes: 1) formation

of atomic hydrogen on the alloy surface (Eq. (1)); and 2) inward

diffusion of hydrogen. To support the former step, arsenicwas

added in several concentrations. However, its effect on the

resulting diffusion profiles is small (Fig. 4e), suggesting that

the amount of atomic hydrogen on the surface is sufficient

and that the rate-controlling step is H diffusion in the

structure.

It is known that once a MgH2 continuous layer is formed on

the surface of Mg dendrites, interstitial hydrogen diffusion in

this layer is slow due to strong MgeH interactions. For this

reason, hydriding would not proceed for pure Mg [10].

However, the addition of catalysts, in our case Ni and RE,

accelerate H diffusion in the structure (Fig. 5). It is believed

that the acceleration can be attributed to: 1) the reduction in

the MgeH bond strength; 2) increase in the free space in the

crystal lattice (mainly by large RE atoms); 3) formation of

disperse eutectic structures with a high fraction of interfaces

(Figs. 2 and 3). Therefore, the higher the catalyst contents, the

higher the fraction of eutectics and the faster the H diffusion.

This is seen in Fig. 4. When the eutectic fraction is small, as in

the case of the MgNi11 alloy (Figs. 2a and 4a), the hydrogen

concentration profile is low, and the penetration depth is

short (Fig. 5). The other extreme is the MgNi24Mm5 alloy in

which the disperse eutectic predominates. This alloy achieves

the maximum penetration depth of hydrogen and, thus, the

best hydriding behavior, see Figs. 2d and 4.

If we take into account that the rate-limiting step of the

hydriding process is H diffusion, the diffusion penetration

depths summarized in Fig. 5 can be used for an approximate

estimation of the H diffusion coefficients in the alloys. A

simplified version of Fick’s second law (X2 ¼ 2Ds) can be used

for this purpose in which X is the hydrogen penetration depth,

D the diffusion coefficient and s the diffusion time. The results

are given in Table 3. The diffusion coefficient estimated for the

pure eutectic structure (MgNi24Mm5 alloy) is in good agree-

mentwith that reported by Cui for the pureMg2NiH0.3 phase at

Table 3 e Estimated hydrogen diffusion coefficients D inthe alloys at 20 �C.

Alloy D (cm2 s�1)

MgNi11 2 $ 10�11

MgNi11Mm4 3 $ 10�11

MgNi11Mm6 1 $ 10�10

MgNi15 3 $ 10�11

MgNi15Mm4 1 $ 10�10

MgNi24Mm5 4 $ 10�10

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25 �C (3 $ 10�10 cm2 s�1) [17], suggesting that the hydrogen

diffusion in this phase plays a significant role, as indicated by

Eq. (2). He also reported that the H diffusion coefficient in the

Mg2NiH4 hydride is lower (9 $ 10�11 cm2 s�1) due to stronger

MgeH interactions. In our case, the diffusion coefficients are

also reduced with decreasing amounts of alloying elements

and with an increasing number of strong MgeH bonds.

4.3. Hydrogen evolution

As illustrated in Fig. 7, the MgH2 phase in the hydrided MgeNi

and MgeNieMm alloys is destabilized because its decompo-

sition temperature is reduced by about 100 �C in comparison

with that of pure MgH2. It can be assumed that in the struc-

tures of the hydrided alloys, MgH2 particles are in contact with

Ni- and Mm-rich phases. Apparently, both Ni and RE metals

facilitate the breaking of MgeH bonds in the hydride. As

indicated before, this may be due to the occupation of valence

electron bands in these metals. Nickel, as well as rare earth

metals, have unsaturated d-electron shells, i.e., there is

a tendency to stabilize the electron configuration by sharing

electrons with other atoms. This may be accomplished by an

interaction with one valence electron of atomic hydrogen

originating from magnesium hydride. Once a hydrogen atom

is released from the hydride, it diffuses through the structure,

which is facilitated by the presence of Ni and RE, as indicated

in the previous paragraph.

5. Conclusions

It is shown in this work that electrochemical hydriding of

MgeNieMm alloys is a viable process for the preparation of

MgH2 hydride. Ni- and Mm-containing binary and ternary

hydrides were not formed in this process, probably due to

kinetic reasons. Therefore, both Ni and Mm act as catalysts of

the hydriding process because they support the formation of

atomic hydrogen on the surface and accelerate hydrogen

diffusion. The formation of MgH2 as the only hydriding

product seems to be beneficial because the binarymagnesium

hydride achieves the maximum hydrogen gravimetric den-

sity. In addition, this phase becomes destabilized by nickel

and rare earths. Among the hydrided alloys, a maximum

hydrogen gravimetric density of more than 1% was achieved

for the eutectic MgNi24Mm5 alloy hydrided at 80 �C. At first

sight, this value appears low, but it is not very far from that for

a LaNi5-based hydride that is commercially used in Ni-MH

batteries (1.4% [18]). In addition, the structure of the as-cast

alloys used in our experiment was relatively coarse, in spite of

the presence of a disperse eutectic. We believe that structural

refinement by ball milling or rapid solidification would

certainly increase the maximum hydrogen concentrations

and MgH2 fractions achieved. It should be noted that signifi-

cantly higher gravimetric densities were obtained by many

authors for hydrides based on magnesium alloys. However,

they mostly used high-pressure and high-temperature hyd-

riding with gaseous hydrogen. The objective of our research is

to find amore simple and less expensive process to synthetize

hydrides. Electrochemical hydriding appears as a viable

process for this purpose. Our further research will be focused

on structural investigation of the hydrided alloys and on their

cyclic stability.

Acknowledgments

Research on electrochemical preparation of hydrides is finan-

cially supported by the Czech Science Foundation (project no.

104/09/0263), the Czech Academy of Sciences (project no.

KAN300100801) and by the Ministry of Education, Youth and

Sports of theCzechRepublic (projectsno.MSM6046137302 and

MSMT no. 21/2011).

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