Recent progress in hydrogen storage alloys for nickel/metal hydride ...

Review 

Recent progress in hydrogen storage alloys for nickel/metal 

hydride secondary batteries 

Xiangyu Zhao, Liqun Ma* 

College of Materials Science and Engineering, Nanjing University of Technology, 5 Xinmofan Road, Nanjing, 210009, P.R. China 

article info 

Article history: 

Received 26 June 2008 

Received in revised form 

23 February 2009 

Accepted 13 March 2009 

Available online 25 April 2009 

Keywords: 

Hydrogen storage alloy 

Mg-based alloy 

Composite 

Charge transfer 

Hydrogen diffusion 

1. Introduction 

international journal of hydrogen energy 34 (2009) 4788–4796 

abstract 

Hydrogen storage alloys, the main materials in the negative 

electrodes of nickel/metal hydride (Ni–MH) secondary 

batteries, have been extensively studied for many years 

because of their dominant role in the battery. Kuriyama et al. [1] 

have reviewed some efforts such as substitution of alloy 

components, heat treatment, and surface treatment to improve 

the performance of the hydrogen storage alloys, mainly LaNi5based 

alloys. Hong [2] has surveyed the development of 

commercial hydrogen storage alloys, also recommended is the 

work of Furukawa [3]. For a more comprehensive summary of 

hydrogen-absorbing alloys, see Akiba [4]. 

Only rare-earth LaNi 5-based and Zr–Ti–V-based laves 

phase hydrogen storage alloys have been used as negative 

electrode materials for the commercial production of Ni–MH 

Available at www.sciencedirect.com 

journal homepage: www.elsevier.com/locate/he 

This paper reviews the development of hydrogen storage alloys prepared by an effective 

method of mechanical alloying and milling. It emphasizes alloys based on Mg or that 

contain Mg due to their low cost, low weight and high hydrogen storage capacity. Hydrogen 

absorption/desorption and electrochemical measurements are briefly discussed. The 

electrochemical properties of the alloys that contain Mg are covered in detail, emphasizing 

the effects of changes in alloy composition. The system of Ti–Ni-based alloys is also 

introduced. At present, composite hydrogen storage alloys may be the most effective 

materials for practical application in new nickel/metal hydride secondary batteries. The 

steps of hydrogen absorption/desorption such as charge-transfer and hydrogen diffusion 

for evaluating the electrochemical properties of hydrogen storage alloys are discussed. The 

relationship between alloy composition and electrochemical properties is noted and 

evaluated. 

ª 2009 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights 

reserved. 

batteries [5]. However, these materials have low hydrogen 

storage capacities, i.e. low energy densities. Their high cost 

and heavy weight also limit their application. The design of 

hydrogen storage alloys depends on three aspects: alloy 

composition, bulk structure, and surface structure. Apparently, 

the preparation method is the primary approach for 

varying these three aspects. Sakai et al. [6] mentioned that 

commercial hydrogen storage alloys were produced by a rapid 

solidification process followed by a pulverization process into 

powders for preparing the electrode. Mechanical alloying 

(MA), a process of milling powder materials and a technology 

of preparing non-equilibrium materials, has several important 

advantages for preparing hydrogen storage alloys: 

It allows easy and controlled synthesis of equilibrium and 

metastable alloy phases such as crystalline, 

* Corresponding author. Tel.: þ86 25 8358 7243; fax: þ86 25 8324 0205. 

E-mail addresses: zhaoxycc@yahoo.com.cn (X. Zhao), maliqun@njut.edu.cn (L. Ma). 

0360-3199/$ – see front matter ª 2009 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights reserved. 

doi:10.1016/j.ijhydene.2009.03.023

nanocrystalline, amorphous, and quasicrystalline in spite of 

large differences in the melting points of the raw materials. 

It can effectively synthesize a composite hydrogen storage 

material which contains two or more hydrogen storage alloys. 

Its product is a powder which can be used directly for 

hydrogen storage materials without subsequent size 

reduction. 

It is a simple and inexpensive processing technology operating 

at room temperature. 

Suryanarayana [7] and El-Eskandarany [8] have published 

two classic texts on mechanical alloying and milling. During 

high energy milling the powder particles are repeatedly flattened, 

cold welded, fractured and rewelded. The typical 

morphology with a lamellar structure which is formed by the 

alternation of cold welding and fracture decreases the diffusion 

distance. Furthermore, the particles are heavily deformed, 

and a variety of crystal defects such as dislocations, vacancies, 

stacking faults, and an increased number of grain boundaries, 

which can enhance the diffusivity of solute elements into 

the matrix, result. Additionally, the slight rise in temperature 

during milling, perhaps including local high temperatures, 

further aids the diffusion behavior, and as a result, true alloying 

takes place among the constituent elements. 

The classes of hydrogen storage alloys can be principally or 

conventionally classified as AB5-type alloys, AB3-type alloys, 

A2B7-type alloys, AB2-type alloys, AB-type alloys, Mg-based 

alloys and V-based solid solution alloys. For preparing 

hydrogen storage alloys, the application of MA and/or 

mechanical milling (MM) began in the 1980s when the Mg2Ni 

hydrogen storage alloy was fabricated [9], and since then this 

technology for the study of hydrogen storage alloys has 

flourished worldwide. In this review, the electrochemical 

properties and the development of Mg-based, AB 3-type, A 2B 7type, 

AB-type alloys, and composite hydrogen storage alloys 

that are prepared by mechanical alloying and milling will be 

introduced and discussed. 

2. Steps of hydrogen absorption/desorption 

Ni–MH batteries contain an electrochemical system involving 

a hydrogen storage electrode and a Ni(OH) 2/NiOOH counter 

electrode. The hydrogen storage alloy plays a dominant role in 

the power and service life of a Ni–MH battery and determines 

the electrochemical properties of the battery, and is therefore 

still a research topic of great interest. The hydrogen evolution 

reaction (HER) for hydrogen storage alloys involves the 

following steps: 

M þ H2O þ e 4MHads þ OH (1) 

MHads4MHabs 

MHabs4MHhydride 

2MHads4M þ H2 

international journal of hydrogen energy 34 (2009) 4788–4796 4789 

MHads þ H2O þ e 4M þ H2 þ OH (5) 

(2) 

(3) 

(4) 

i.e., the steps of adsorption/desorption, surface penetration, 

hydrogen diffusion, and formation/decomposition of metal 

hydride. In the initial step of the HER, the adsorption of 

hydrogen atoms occurs on the surface of a hydrogen storage 

alloy by dissociation of water. These atoms can be absorbed to 

form metal hydride, but can also combine together to produce 

hydrogen gas according to the Tafel reaction Eq. (4) and the 

Heyrovsky reaction Eq. (5). These steps can be also effectively 

described in a charge/discharge curve for a hydrogen storage 

alloy, as shown in Fig. 1. In stage OA, the potential of the 

electrode rapidly changes with time or hydrogen content. This 

stage is the formation of a solid solution called the a-phase. 

After this stage, in the section AB, the potential remains 

constant as the hydrogen content increases. This constant 

plateau region indicates a progressive conversion of the aphase 

to a metal hydride called the b-phase. Eventually the 

curve begins to slope again, indicating that the conversion of 

a-phase to b-phase is finished. Finally, with further increase of 

the charging time, there is another plateau region which 

corresponds to the process of hydrogen evolution and indicates 

that the charge of the electrode is saturated. During 

discharge, the b-phase decomposes to the a-phase, reversing 

the charging process. It is obvious that the performance of the 

hydrogen storage alloy is determined in the stage where the 

two phases coexist, i.e., the kinetics of the charge transfer 

reaction at the electrode surface, the rate of hydrogen transfer 

between the absorbed state and the adsorbed state, and the 

diffusion of absorbed hydrogen between the bulk and the 

electrode surface [10]. 

Many efforts [11–17] have investigated the steps of 

hydrogen absorption/desorption, especially the chargetransfer 

reaction at the electrode/electrolyte interface and 

hydrogen diffusion within the bulk electrode. These are the 

two dominant factors, particularly for the high rate dischargeability 

(HRD) of hydrogen storage alloy electrodes 

[18–20]. The two steps can be expressed by the values of 

electrochemical kinetic parameters such as exchange current 

density, polarization resistance, and hydrogen diffusion 

coefficient. 

Potential (V) 

O 

A 

charge 

B 

C 

D 

Charge-discharge time (h) 

E 

discharge 

Fig. 1 – Schematic illustration of a charge–discharge curve 

for a hydrogen storage alloy. 

F 

G

4790 

The exchange current density of hydrogen storage alloy 

electrodes can be determined by linear polarization [18,21], 

Tafel polarization [21,22], and electrochemical impedance 

spectroscopy (EIS) [22,23]. The exchange current density I0 can 

be obtained from a linearized form of the Butler–Volmer 

equation in the low overpotential region (h < 10 mV). The 

linear equation can be written as: 

I0 ¼ IRT 

(6) 

Fh 

where I is the applied current density, R is the gas constant, T 

is the absolute temperature, F is the Faraday constant, and h is 

the overpotential of the electrochemical reaction for the 

hydrogen storage alloys. The Tafel region is controlled by the 

step of charge transfer at high overpotential (h > 118 mV). For 

porous hydrogen storage alloy electrodes, the Tafel region 

shows strong mass-transfer effects at high currents. Using the 

measured limiting current densities from the curve of Tafel 

polarization, the Tafel plots can be corrected for mass-transfer 

effects by plotting the logarithm of I/(1 I/IL) against the 

electrode potential, which can be expressed as: 

h ¼ 2:3RT 

bF log1 þ 

I0 

2:3RT 

bF log 

I 

1 I=IL 

where I L is the limiting current density, and b is the transfer 

coefficient for hydrogen desorption. For EIS, the semicircle in 

the high frequency region is related to the contact resistance 

between the current collector and the alloy pellet, and the 

semicircle in the low frequency region corresponds to the 

charge-transfer resistance Rct [23–25]. The exchange current 

density can be calculated using the following Equation [21,26]: 

I0 ¼ RT 

F 

1 

Rct 

Ratnakumar et al. [22] and Witham et al. [27] found that 

exchange current densities estimated from linear polarization 

were in agreement with those from EIS, but were quite 

different from the values obtained from Tafel polarization. 

Wang et al. [10] proposed that this discrepancy may be caused 

by a neglect of the effect of hydrogen transfer between the 

absorbed and adsorbed states. 

Many approaches have been proposed for evaluating the 

hydrogen diffusion coefficient, including nuclear magnetic 

resonance [28], quasi-elastic neutron scattering [29], and 

various electrochemical technologies of current pulse [30], 

cyclic voltammetry [31], EIS [32], and potential-step [33,34]. 

The potential-step method, a simple and convenient technology, 

is widely used to study the hydrogen diffusion coefficient. 

When a large potential-step is applied to hydrogen 

storage alloy electrodes, a drastic depletion of hydrogen 

occurs on the alloy surface. After a long response time, the 

current decreases slowly in a linear fashion. In this linear 

region, hydrogen diffusion controls the electrode process, and 

the hydrogen diffusion coefficient D can be calculated 

according to the following equation [20,33–35]: 

log i ¼ log 6FD 

da 2 ðC0 CsÞ 


p2 D 

2:303 a2t where i (A/g) is the diffusion current density, D (cm 2 /s) is the 

hydrogen diffusion coefficient, d (g/cm 3 ) is the density of the 

(7) 

(8) 

(9) 

hydrogen storage alloy, a (cm) is the alloy particle radius, C 0 

(mol/cm 3 ) is the initial hydrogen concentration in the bulk of 

the alloy, Cs (mol/cm 3 ) is the hydrogen concentration on the 

surface of the alloy particles and t (s) is the discharge time. 

This method needs a long time period, generally several 

thousand seconds, to measure the response of the anodic 

current. As a result, the calculated hydrogen diffusion coefficient 

will be an average value. Consequently, this method is 

unsuitable for determining the hydrogen diffusion coefficient 

at a certain depth of discharge (DOD). Feng et al. [17,36] 

described a new and relatively simple potentiostatic method 

for determining the hydrogen diffusion coefficient over 

a small time period, less than 500 s. The hydrogen diffusion 

coefficient can be found by measuring the ratio of intercept 

and slope from a linear plot of I(t) versus t 1/2 by using the 

following equation: 

IðtÞ ¼ FADðCs C0Þ 

1 1 

pffiffiffiffiffiffi 

pffiffit 

pD 

1 

a 

(10) 

where A is the surface area of the particles (m 2 /g). This 

method does not require knowledge of either the hydrogen 

concentration or the surface area of the alloy particles. The 

hydrogen diffusion coefficient at a certain DOD can be 

obtained by using this method. 

Apart from the steps of charge-transfer and hydrogen 

diffusion, the process of hydrogen transfer, which was 

neglected in previous studies, is also discussed by Wang et al. 

[10] who assume that the pressure plateau of a hydrogen 

storage alloy is flat and the fully charged hydride electrode is 

considered to be uniformly discharged at each DOD. 

3. Mg-based alloys 

Mg has a high reversible storage capacity of 7.6 wt.% (approx. 

2200 mAh/g) hydrogen, which is principle makes it a promising 

candidate as a storage medium in mobile applications. In 

recent years, more attention has been paid to developing 

applications of Mg-based alloys in Ni–MH secondary batteries 

because of their high storage capacity, low cost and low 

weight. MA is a simple and effective method for preparing Mgbased 

alloys with metastable or non-equilibrium phases, 

which do not appear in phase-diagrams, in spite of large 

differences in the melting points of the raw materials. The 

alloy composition can be simply controlled. Singh et al. [37] 

described a nanocrystalline Mg2Ni alloy with a grain size of 

about 4 nm formed after milling magnesium and nickel 

powders. Amorphous Mg2Ni alloy can also be prepared by MA 

[38]. Cui et al. [39] showed that the formation and growth of 

a magnesium oxide and/or hydroxide layer on Mg 2Ni 

increased the electronic resistance at the electrode/electrolyte 

interface and resulted in a large discharge overpotential, 

leading to low discharge capacity and sluggish kinetics of the 

alloy. In order to improve the electrochemical properties of 

Mg-based alloy, element substitution of Mg-based alloys has 

been extensively attempted. Although Zr [40], Ti[41], Co[42], 

Al [43],Ce[44], Y[45], Ca[46] and Fe [47] have been introduced 

into Mg2Ni alloy to improve their hydrogen absorption/ 

desorption properties, there was no apparent improvement.

Liu et al. [48] described a composition of Mg 0.7Ti 0.225Al 0.075Ni 

which showed a low decay of discharge capacity due to the 

formation of MgTi2O4 instead of Mg(OH)2. A similar positive 

result was reported by Tian et al. [49] who found that 

a composition of Mg0.8Ti0.1Pd0.1Ni possessed good cycle 

stability because of a protective layer of (NiO)x(PdO)y(TiO2)z. 

Apart from element substitution, Rongeat et al. [50] proposed 

that an increase of MgNi particle size could effectively 

improve the cycle stability and high rate dischargeability 

because larger particles had a lower sensitivity to oxidation. 

Then, a comprehensive work [51] that considered both 

element substitution and particle size showed that an amorphous 

Mg0.9Ti0.1NiAl0.05 alloy with particle size larger than 

150 mm displayed excellent electrochemical properties, as 

good as a commercial LaNi5-based alloy, due to a control of the 

charge input and a cooperative protection by Ti and Al. 

Obviously the cost and weight are much lower than a LaNi5based 

alloy. However, the high rate dischargeability is presently 

too low, which will limit the potential use of Mg-based 

alloys in real applications; this should be studied further. 

4. AB 3- and A 2B 7-type alloys 

AB3- and A2B7-type alloys have been studied for decades. Ivey 

et al. [52] reviewed binary alloys correlated to AB3 and A2B7 

alloys, and mentioned that the structures of AB3 and A2B7 can 

be related to those of AB5 and AB2. Based on the assumption 

that there are no long-range H–H interactions, these relations 

can be expressed as [53]: 

nðAB3Þ ¼ 1 2 

nðAB5Þþ 

3 3 nðAB2Þ (11) 

and 


nðA2B7Þ ¼nðAB5ÞþnðAB2Þ (12) 

where n(AkBm) represent the hydrogen concentration for each 

of the respective formula units. Zhang et al. [54] described the 

formation of amorphous LaNi 2H x and LaNi 5H y after hydrogen 

absorption by LaNi 3 at 298 K. The amorphous LaNi 2H x would 

decompose to LaH 2 and LaNi 5H z at 473 K. Similar results were 

reported by Srivastava [55] who also claimed that both LaNi 3 

and La2Ni7 had lower storage capacity than LaNi5 but better 

suppression of pulverization. In recent years, ternary or 

multiple AB3 and A2B7 alloys have been developed. Kadir et al. 

[56,57] reported a ternary system of R–Mg–Ni (R ¼ rare earth, 

Ca or Y) with a PuNi3-type rhombohedral structure, which was 

of great interest for many researchers. Liang et al. [58] presented 

a mechanically alloyed Ca2Ni7, the structure of which 

could be destroyed by substituting Mg for a small fraction of 

the Ca, resulting in the formation of an AB 3 phase with an 

accompanying nickel phase, but the as-milled alloys absorbed 

hydrogen with poor reversibility. Ml (La-rich mischmetal) can 

be effectively used for the A side [59]. The element Ni on the B 

side can be partially substituted by Co, Mn, Fe, Al and Cu to 

improve the kinetics and cycle stability [60,61]. For electrochemical 

properties, Zhang [62] fabricated a composition of 

La1.5Mg0.5Ni7 with a maximum discharge capacity of 

389.4 mAh/g, good activation characteristics and good cycle 

stability. Its HRD showed a high value of 92.3% at a current 

density of 900 mA/g. This year a new phase, MgNi 3, with the 

same cubic crystal lattice as AuCu3 has been reported [63]. 

More work needs to be carried out with AB3- and A2B7-type 

alloys. 

5. AB-type alloys 

Early work on AB-type hydrogen storage alloys was restricted 

to TiFe. There are two stable intermetallic compounds formed 

by the Ti–Fe system, TiFe and TiFe 2 [52]. The application of 

TiFe in Ni–MH batteries has been limited because of its poor 

kinetics of hydrogen absorption/desorption [64]. In recent 

years, more AB-type hydrogen storage alloys consisting of Ti, 

Zr or Hf on the A side and Fe, Ni, Al, Co, Mn, or Sn on the B side 

have been investigated [65–69]. The substitution of Ni for Fe 

could improve the activation performance and discharge 

capacity. TiFe 0.25Ni 0.75 showed a high discharge capacity of 

155 mAh/g on the third cycle [67]. Although the Ti–Ni system 

has been studied for many years, it is still in an active research 

topic [70–73]. Cuevas et al. [74] have shown that martensitic 

Ti 0.64Zr 0.36Ni exhibited much higher reversible capacity, about 

330 mAh/g. Nevertheless, the electrochemical properties of 

this system of alloys are unsatisfactory [73]. 

Ti–Ni-based alloys with a similar composition but a quasicrystalline 

structure have been described in the past few years 

[75,76]. There are many interstices which could be suitable 

sites for hydrogen atoms in a quasicrystalline Ti–Zr–Ni alloy, 

compared to a normal crystal structure [76]. Liu et al. [26] 

found that Ti 45 xZr 35 xNi 17þ2xCu 3 icosahedral quasicrystalline 

phase (I-phase) had a large discharge capacity of 269 mAh/g 

without initial activation, when x was up to 8. Increased nickel 

content improved the electrochemical kinetic properties and 

prevented oxidation of the alloy electrodes. The hydrogen 

storage alloys with quasicrystalline structure show good 

hydrogen storage properties. However, direct fabrication of 

this quasicrystalline phase by MA has not been reported. 

Takasaki et al. [77] explained that MA processes elemental 

powders mostly by dynamic force or mechanical collision, 

which leads to chemically inhomogeneous final products and 

makes it difficult to produce a single quasicrystalline phase. 

6. Composite hydrogen storage alloys 

Desirable electrochemical properties of hydrogen storage 

alloys include high storage capacity, easy activation, high 

resistance to corrosion, favorable kinetic performance, high 

HRD, and low cost. In practice, it is difficult to simultaneously 

obtain all these properties in a single alloy system, but using 

a composite hydrogen storage alloy is an effective way to 

achieve it. A composite hydrogen storage alloy contains two or 

more hydrogen storage alloys, or a hydrogen storage alloy and 

another intermetallic compound. Generally, the major 

component in a composite hydrogen storage alloy is an alloy 

with good hydrogen storage properties. The minor component 

is a surface activator to improve the activation properties and 

the kinetics of hydrogen sorption/desorption [78]. 

AB5- and AB2-type alloys are the two conventional 

hydrogen storage alloys. Each one has its own advantages,

4792 


such as easy activation of AB 5 and high storage capacity of 

AB2. Chen et al. [79] milled an AB2 Zr–Ti–V–Ni alloy with added 

LaNi5, and showed that the matrix alloy was successfully 

coated with fine nanocrystalline LaNi5 and the electrochemical 

activation of the AB2 alloy was effectively improved. 

A similar result was obtained by Han et al. [80] who found that 

the addition of LaNi5 improved not only the activation 

performance but also the high rate dischargeability of the 

matrix due to a segregated La–Ni phase which provided active 

sites and pathways for hydrogen diffusion. Another interesting 

result was that the addition of AB 5 alloy could drastically 

decrease the amount of V dissolved from the matrix by 

the KOH electrolyte [81]. On the contrary, an AB2 alloy added to 

a matrix of LaNi5 alloy [82] enhanced the discharge capacity of 

the matrix. 

Yang et al. [78] found that amorphous Mg–Ni alloy can be 

an excellent surface activator to eliminate the initial activation 

and substantially improve the kinetic properties of alloys 

such as Zr(Ni0.6Mn0.15Cr0.1V0.15)2 and ZrCrNi, which are usually 

considered unsuitable for practical application because of 

their poor kinetics and difficult activation. Choi et al. [83] 

demonstrated a similar result in a system of TiV 2.1Ni 0.3–MgNi 

composites, and explained that the amorphous MgNi alloy not 

only acted as a protective film against corrosion, but also 

formed a uniform layer with high electrocatalytic performance 

and high elasticity on the surface of the matrix. It is 

well known that Mg-based alloys have attractively high 

storage capacity, low cost and light weight. Unfortunately, 

they have poor hydrogen absorption/desorption characteristics, 

which seriously blocks their practical use in energy 

storage. Cui et al. [84] investigated a composite alloy of Mg 2Ni– 

40 wt%Ti2Ni. The addition of Ti2Ni particles inlaid on the 

surface of Mg2Ni particles improved both the charge-transfer 

reaction at the surface of alloy and hydrogen diffusion in the 

bulk alloy. Some intermetallic compounds such as CoB [85] 

and TiB [86] are also effective additions for preparing 

composite hydrogen storage alloys. In recent years, magnesium 

has been used as a substitutable element in the system 

of La–Mg–Ni alloys which were extensively studied as negative 

electrode candidates for Ni–MH secondary batteries due to 

their high discharge capacities [58–62]. However, the inferior 

cycle durability of these alloys limits their real application. 

Preparing composite hydrogen storage alloys is also a good 

method to improve their overall properties [87–89]. 

7. Electrochemical kinetics of hydrogen 

absorption/desorption 

The thermodynamic properties of a metal-hydrogen system 

are conveniently summarized by a pressure–composition– 

temperature (PCT) curve [90]. Thermodynamic parameters 

such as the enthalpy DH and the entropy DS of hydride 

formation can be determined from the Van’t Hoff equation 

[91]. 

Apart from thermodynamic properties, another important 

criterion for a metal-hydrogen system is the electrochemical 

kinetics of hydrogen absorption/desorption, i.e. the kinetics 

of the charge-transfer reaction at the electrode/electrolyte 

interface and hydrogen diffusion within the bulk electrode. 

Table 1 – Exchange current densities for hydrogen storage alloys at 50% DOD. 

Composition Phase structure Exchange current 

density, I0 (mA/g) 

Reference 

Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.25 C14 Laves phase, V-based 

solid solution phase 

142.8 [93] 

Ti0.8Zr0.2V2.7Mn0.5Cr0.5Ni1.25Fe0.3 178.7 [93] 

Ti0.8Zr0.2V2.7Mn0.5Ni1.25Fe0.8 46.96 [93] 

Ti45Zr35Ni17Cu3 I-phase 199.5 [26] 

Ti43Zr33Ni21Cu3 229.1 [26] 

Ti41Zr31Ni25Cu3 301.2 [26] 

Ti37Zr27Ni33Cu3 480.7 [26] 

La2MgNi9 PuNi3-type 88.5 [60] 

La2Mg(Ni0.09Al0.01)9 65.4 [60] 

La2Mg(Ni0.08Al0.02) 9 62.8 [60] 

La1.7Mg0.3Ni7.0 LaNi5, PuNi3-type 128.1 [62] 

La1.6Mg0.4Ni7.0 228.7 [62] 

La1.5Mg0.5Ni7.0 227.7 [62] 

La1.4Mg0.6Ni7.0 167.3 [62] 

La0.7Mg0.3(Ni0.85Co0.15) 3 PuNi3-type 186.5 [19] 

La0.7Mg0.3(Ni0.85Co0.15)4 222.3 [19] 

La0.7Mg0.3(Ni0.85Co0.15)5 190.1 [19] 

Mg0.86Ti0.1Pd0.04Ni Amorphous 256 [49] 

Mg0.84Ti0.1Pd0.06Ni 247 [49] 

Mg0.8Ti0.1Pd0.1Ni 184 [49] 

La0.7Mg0.3Ni3.5 PuNi3-type, LaNi5 79.6 [89] 

Ti0.17Zr0.08V0.35Cr0.1Ni0.3 V-based solid solution phase, C14 Laves phase 99.5 [89] 

La0.7Mg0.3Ni3.5 þ 5 wt.% Ti0.17Zr0.08V0.35Cr0.1Ni0.3 PuNi3-type, LaNi5, V-based solid solution 

phase, C14 Laves phase 

362.9 [89] 

La 0.7Mg 0.3Ni 3.5 þ 50 wt.% Ti 0.17Zr 0.08V 0.35Cr 0.1Ni 0.3 201.6 [89]


These can be expressed by the values of the exchange 

current density I0 and/or the polarization resistance Rp, as 

well as the hydrogen diffusion coefficient D and/or diffusion 

resistance RD. 

Tables 1 and 2 summarize the I0 and D values of hydrogen 

storage alloys in a half-cell containing a Ni(OH)2/NiOOH 

counter electrode and a Hg/HgO reference electrode in the 

KOH solution, respectively. Zheng et al. [18] have found that 

the exchange current density I 0 obtained on LaNi 4.27Sn 0.24 

alloy electrodes was a function of the bulk hydrogen concentration. 

Ramya et al. [94] noticed similar results for 

a TiMn 1.6Ni 0.4 alloy and proposed that a modified surface 

could improve the charge-transfer reaction. There is no doubt 

that the properties of hydrogen storage alloys are principally 

influenced by the electrochemical characteristics of the alloys 

in the region of the potential plateau. That is why the 

exchange current density was determined at 50% DOD in 

a number of studies. Enhancement of the exchange current 

density can be related to modification of the surface state. 

Many approaches to surface modification have been attempted 

to dissolve the oxide layer and promote the formation of 

an active surface with excellent electrocatalytic performance 

[19,20,26,49,78,89]. It is interesting that the exchange current 

densities of various research groups can be increased to the 

same level by either element substitution or composite alloys 

as shown in Table 1. The electrochemical properties of 

hydrogen storage alloys may be strongly influenced by 

transport properties in the bulk alloys. The potential-step 

method for calculating the value or the average value [17,36] of 

hydrogen diffusion coefficients has been extensively used as 

shown in Table 2. It can be seen that this method is effective 

for evaluating the change of hydrogen diffusion coefficients 

within each research group. For comparison between various 

research groups, as Feng et al. [17] remarked, values obtained 

at the same DOD should be used. 

Presently, there is no doubt that increasing the exchange 

current density and the hydrogen diffusion coefficient 

increases the charge and discharge capacities and efficiencies 

of hydrogen storage alloys. The exchange current density 

related to the charge-transfer step is determined both by 

crystallographic and electronic structure [95]. Different alloy 

compositions on the alloy surface influence the valence 

electron configurations, which essentially determined the 

charge-transfer reaction Eq. (1), i.e. the hydrogen dissociative 

reaction [96]. The surface s- and d-electrons play an important 

role in the hydrogen dissociation on a metal surface [96]. In 

the case of a substrate atom whose valence electron states are 

fully occupied, the H s-electron is repelled. Jaksic [97] has 

reported that a synergistic effect in hydrogen electrosorption– 

desorption should arise by alloying metals having unoccupied 

d-band states with those having internally paired d-electrons. 

For instance, the valence electron configurations of Mg, La, Ni 

atoms are 2p 6 3s 2 ,6s 2 5d 1 and 4s 2 3d 8 , respectively. In the case 

of Mg, the s- and p-orbitals are fully occupied and no d-orbital, 

Table 2 – Hydrogen diffusion coefficients D obtained by using Eq. (9) for hydrogen storage alloys. 

Composition Phase structure Hydrogen diffusion 

coefficient, D (cm 2 /s) 

Mg2Ni Mg2Ni 4.5 10 10 

Mg1.9V0.1Ni0.8Co0.2 1.3 10 9 

La1.7Mg0.3Ni7.0 LaNi5, PuNi3-type 4.5 10 10 

La1.6Mg0.4Ni7.0 6.4 10 10 

La1.5Mg0.5Ni7.0 8.4 10 10 

La1.4Mg0.6Ni7.0 7.7 10 10 

La2Mg(Ni0.95Al0.05)9 PuNi3-type 1.18 10 10 

La2Mg(Ni0.95Sn0.05)9 7.57 10 10 

La0.7Mg0.3(Ni0.85Co0.15) 3 1.61 10 10 

La0.7Mg0.3(Ni0.85Co0.15) 4 2.20 10 10 

La0.7Mg0.3(Ni0.85Co0.15)5 1.48 10 10 

Mg0.86Ti0.1Pd0.04Ni Amorphous 3.0 10 9 

Mg0.84Ti0.1Pd0.06Ni 4.1 10 9 

Mg0.8Ti0.1Pd0.1Ni 5.6 10 9 

TiNi Nanocrystalline 2.73 10 12 

TiNi0.8Mn0.2 9.2 10 12 

TiNi0.8Mn0.2 Amorphous/Nanocrystalline 1.2 10 10 

Ti45Zr35Ni17Cu3 I-phase 5.8 10 10 

Ti43Zr33Ni21Cu3 6.3 10 10 

Ti41Zr31Ni25Cu3 7.4 10 10 

Ti37Zr27Ni33Cu3 10.5 10 10 

Ti0.8Zr0.2V2.7Mn0.5Cr0.8Ni1.25 C14 Laves phase, V-based solid solution phase 4.6 10 11 

Ti0.8Zr0.2V2.7Mn0.5Cr0.5Ni1.25Fe0.3 6.01 10 11 

Ti0.8Zr0.2V2.7Mn0.5Ni1.25Fe0.8 2.7 10 11 

Ti0.9Zr0.2Mn1.5Cr0.3V0.3 C14 Laves phase 0.62 10 14 

Ti0.9Zr0.2Mn1.5Cr0.3V0.3 þ LaNi3.8Mn0.3Al0.4Co0.5 C14 Laves phase, LaNi5 3.19 10 14 

La0.7Mg0.3Ni3.5 þ 5 wt.% Ti0.17Zr0.08V0.35Cr0.1Ni0.3 PuNi3-type, LaNi5, V-based solid 

solution phase, C14 Laves phase 

1.67 10 11 

La0.7Mg0.3Ni3.5 þ 25 wt.% Ti0.17Zr0.08V0.35Cr0.1Ni0.3 1.28 10 11 

La0.7Mg0.3Ni3.5 þ 40 wt.% Ti0.17Zr0.08V0.35Cr0.1Ni0.3 1.51 10 11 

Reference 

[35] 

[35] 

[62] 

[62] 

[62] 

[62] 

[61] 

[61] 

[19] 

[19] 

[19] 

[49] 

[49] 

[49] 

[73] 

[73] 

[73] 

[26] 

[26] 

[26] 

[26] 

[93] 

[93] 

[93] 

[92] 

[92] 

[89] 

[89] 

[89]

4794 

resulting in the high energy barrier for hydrogen dissociation. 

In cases of Ni and La, although the s-orbital is fully occupied, 

the d-orbital is hardly occupied, resulting in the negligible 

energy barrier [95,98]. It is obvious from Table 1 that elemental 

substitution with the element with unoccupied d-band state is 

fairly useful to enhance the exchange current density of the 

alloy [26,49,60]. Moreover, a synergistic effect in a composite 

alloy substantially improves the electrocatalytic performance 

of the alloy substrate [88,89]. 

In absorption, the dissociated hydrogen atoms are incorporated 

into the solid lattice framework. The hydrogen 

diffusion coefficient of atomic hydrogen in the metallic 

lattices was shown to be dependent on the strength of the 

metal-hydrogen interaction and the hydrogen concentration 

[17,20,36]. It is known that the hydriding properties of MgH2 

are significantly improved by the addition of transition 

metals, such as Ni which acts as a catalyst, resulting in 

a weakening of the bonding between Mg and H atoms and 

consequent an increase in hydrogen diffusion. A similar 

result has been achieved by the partial submission of Mg by 

Pd [49]. Furthermore, a substitution of Mg by V increases the 

lattice volume and enhances the hydrogen diffusion [35]. 

Drenchev et al. [73] reported that amorphous phase in 

Ti–Mn–Ni alloy was beneficial to the hydrogen diffusion in 

the bulk alloy. Liu et al. [26] showed a quasicrystalline 

Ti-based alloy with high hydrogen infinity containing abundant 

tetrahedrons for hydrogen occupation and found that 

the hydrogen absorption/desorption properties of the alloy 

could effectively improved by increasing the Ni content in the 

alloy, and the hydrogen diffusion efficient could be fairly 

enhanced. Composite hydrogen storage alloys [89,92] are 

helpful to the increase in the hydrogen diffusion of the 

matrix alloy by a synergistic effect, which improves the 

hydrogen penetration or hydrogen transfer between the alloy 

surface and bulk alloy. 

8. Summary 


Hydrogen storage alloys have been extensively studied for 

many years. There is an apparent trend to concentrate on low 

cost, light weight and excellent charge–discharge properties. 

This paper presents a review of some interesting hydrogen 

storage alloys prepared by an effective and low cost method of 

mechanical alloying and milling. Alloys based on Mg and that 

contain Mg have received more attention in recent years in 

spite of stubborn difficulties due to their low cost, light weight 

and high hydrogen storage capacity. The system of Ti–Nibased 

alloys shows an attractive quasicrystalline structure 

which is quite favorable for hydrogen storage. At the present 

time, using composite hydrogen storage alloys may be most 

effective for practical battery applications. 

Charge-transfer and hydrogen diffusion are the two most 

important factors for evaluating the electrochemical properties 

of hydrogen storage alloys. They can be expressed by the 

exchange current density I0 and hydrogen diffusion coefficient 

D, respectively. Any comparison of these values in different 

alloys, or in the same alloy system, should be carried out at the 

same DOD. 

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