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Effect of rare earth elements on electrochemical propertiesof La–Mg–Ni-based hydrogen storage alloys

Yuan Lia,b, Da Hanc, Shumin Hana,b,*, Xilin Zhub, Lin Hub, Zhong Zhanga, Yuwen Liub

aState Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao 066004, ChinabCollege of Environmental and Chemical Engineering, Yanshan University, Qinhuangdao 066004, ChinacDepartment of Chemistry, Xiamen University, Xiamen 361005, China

a r t i c l e i n f o

Article history:

Received 26 March 2008

Received in revised form

23 July 2008

Accepted 14 November 2008

Available online 23 December 2008

Keywords:

Hydrogen storage alloys

Rare earth elements

Electrochemical properties

Phase structure

Kinetics

* Corresponding author. College of Environm066004, Hebei Province, China. Tel./fax: þ86

E-mail address: hanshm@ysu.edu.cn (S. H0360-3199/$ – see front matter ª 2008 Interndoi:10.1016/j.ijhydene.2008.11.049

a b s t r a c t

La0.60R0.20Mg0.20(NiCoMnAl)3.5 (R¼ La, Ce, Pr, Nd) alloys were prepared by inductive melting.

Variations in phase structure and electrochemical properties due to partial replacement of

La by Ce, Pr and Nd, were investigated. The alloys consist mainly of LaNi5 phase, La2Ni7

phase and LaNi3 phase as explored by XRD and SEM. The maximum discharge capacity

decreases with Ce, Pr and Nd substitution for La. However, the cycling stability is improved

by substituting Pr and Nd at La sites, capacity retention rate at the 100th cycle increases by

13.4% for the Nd-substituted alloy. The electrochemical kinetics measurements show that

Ce and Pr substitution improves kinetics and thus ameliorates the high rate discharge-

ability (HRD) and low temperature dischargeability. The HRD at 1200 mA g�1 increases from

22.1% to 61.3% and the capacity at 233 K mounts up from 90 mAh g�1 to 220 mAh g�1 for the

Ce-substituted alloy.

ª 2008 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights

reserved.

1. Introduction these methods elemental substitution was found to be an

Metal hydride/Nickel (MH/Ni) batteries are widely used in

portable electronic devices, electric hand tools and even

electric vehicles owing to their high energy density, long cycle

life and environmental friendliness [1]. R–Mg–Ni-based

(R¼ rare earths, Y and Ca) alloys are considered as promising

negative electrode materials due to their higher capacity than

commercial AB5-type alloys in recent years [2,3]. However,

they are subject to poor cycling stability, failing to meet the

need of practical applications [4].

Several methods have been applied to improve hydrogen

storage properties of R–Mg–Ni-based alloys, such as elemental

substitution in alloy composition [5–7], rapid quenching [8],

composite alloying [9], and surface modification [10]. Among

ental and Chemical En335 807 4648.an).

ational Association for H

effective way to improve cycling stability of R–Mg–Ni-based

alloys. A lot of research has been focused on B-elements up to

now. Mn, Co, Al, Mo, Cr, Cu, Fe, etc were adopted in B-sites. In

LaNi3�xMnx alloys, Mn entered the Ni sites and changed the

structure from PuNi3-type to CeNi3-type, decreasing stability

of the hydride against amorphisation on hydrogenation [5].

This may be helpful to improve cycling stability. Co substitu-

tion and combining annealing treatment was believed to

improve cycling stability. The capacity retention rate at the

100th cycle increased from 65.5% for La0.7Mg0.3Ni2.45Co0.75

Mn0.1Al0.2 electrode to 80.5% for La0.7Mg0.3Ni2.15Co1.05Mn0.1Al0.2

electrode [6]. Moreover, the combination of Al and Mo

ameliorated high rate dischargeability, low temperature dis-

chargeability and cycle life [7]. Reports on A-site substitution

gineering, Yanshan University, 438# Hebei Street, Qinhuangdao

ydrogen Energy. Published by Elsevier Ltd. All rights reserved.

i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 1 3 9 9 – 1 4 0 41400

were not enough although investigation on B-site substitution

was systematic. Pan et al. [12] reported Pr substitution for La in

R–Mg–Ni-based alloys La0.7�xPrxMg0.3Ni2.45Co0.75Mn0.1Al0.2 and

found improvement in cycling stability. Zhang et al. [11]

prepared (La, Ce, Pr, Nd)2MgNi9 alloys through an orthogonal

design, and the results showed that Mishmetal composition

has great influence on the La2MgNi9-type alloys. Under the

work typified above, overall electrochemical properties of R–

Mg–Ni-based alloys have been greatly improved; however,

cycling stability is still unsatisfactory. In our previous work,

we have found that low-Mg and B-site multi-component

design was beneficial for improving cycling stability, and we

have also obtained alloys with higher capacity retention rate

by substituting La-rich mishmetal for elemental La [13,14]. But

the effect of each component in the mishmetal (La, Ce, Pr and

Nd) was not clear.

This paper will focus on the effect of rare earth elements on

electrochemical properties of La–Mg–Ni-based hydrogen

storage alloys. The low-Mg and B-site multi-component

La0.60R0.20Mg0.20(NiCoMnAl)3.5 (R¼ La, Ce, Pr, Nd) alloys were

prepared. Phase structure and electrochemical properties of

the alloys will be studied as a part of our systematic work.

10 20 30 40 50 60 70 80

o**

oo+

****

o

o*

+o*o

*o+ *

+

o

o*

+

* LaNi5o La2Ni7+ LaNi3

R = Nd

R = Pr

R = Ce

R = La

Inte

nsit

y (a

.u.)

2 Theta (degree)

Fig. 1 – XRD patterns of La0.60R0.20Mg0.20(NiCoMnAl)3.5

alloys.

2. Experimental

La0.60R0.20Mg0.20(NiCoMnAl)3.5 (R¼ La, Ce, Pr, Nd) alloys were

prepared by inductive melting constituent metals (La, Ce, Pr,

Nd, Ni, Co, Mn, Al) with a purity of 99.5% and Mg–Ni inter-

metallic alloy under argon atmosphere. A Mg–Ni intermetallic

alloy was used instead of Mg and a slight excess of Mg was

used to compensate for evaporative Mg loss during the

melting for the purpose of reducing Mg loss. Then the five-

kilogram ingots were annealed at 1173 K for 8 h. Chemical

composition of the alloys was examined by ICP analysis. The

prepared ingots were mechanically crushed and ground into

fine powders in the air. Powders of 200–300 mesh were used

for electrochemical test and those below 300 mesh for XRD

analysis.

X-ray diffraction (XRD) patterns were obtained on a D/Max-

2500/PC X-ray diffractometer (Cu Ka radiation). Phase identi-

fication and lattice parameters calculation have been carried

out using Jade-5.0 software. Morphology of the alloys was

observed by S-3400 scanning electron microscopy.

A testing electrode was prepared by cold pressing a mixture

of 0.15 g alloy powders and 0.75 g carbonyl nickel powders into

a pellet of 10 mm in diameter under 20 MPa. Electrochemical

measurements were performed in a tri-electrode system con-

sisting of MH electrode (working electrode), Ni(OH)2/NiOOH

electrode (counter electrode) and Hg/HgO electrode (reference

electrode). All the tests were performed at 298 K except for low

temperature dischargeability measurement. The electrodes

were fully charged (the over-charged ratio was approximately

30–50%) at current density of 60 mA g�1, and then discharged at

the same current density during charge/discharge tests.

Discharge capacities were measured at different discharge

current densities when high rate dischargeabilty (HRD) was

measured. Low temperature dischargeability was measured

after laying the tri-electrode system at constant temperature for

8 h using a low-temperature-equipment (WGD701). The

electrodes were discharged to cut-off potential of�0.6 V (vs. Hg/

HgO electrode) in all the discharge process.

The pressure–composition isotherms were obtained

through an electrochemical method. The fully charged elec-

trodes were discharged for 0.5 h, followed by a 0.5 h rest until

cut-off potential was �0.6 V (vs. Hg/HgO electrode). The

equilibrium potential (Eeq) was recorded and converted to

equilibrium pressure (Peq) of hydrogen on basis of the Nernst

equation [15].

EeqðV vs: Hg=HgOÞ ¼ �0:925� 0:03 log�Peq

�The micropolarization curves of the electrodes were

measured on a ZF-9 potentiostat by scanning electrode

potential at the rate of 5 mV min�1 from �5 to 5 mV (vs. open-

circuit potential) at 50% depth of discharge (DOD). For the

potentialstatic step discharge, the fully charged electrode was

discharged at þ500 mV step for 3000 s.

3. Results and discussion

3.1. Phase structure

The XRD patterns of La0.60R0.20Mg0.20(NiCoMnAl)3.5 (R¼ La, Ce,

Pr, Nd) alloys are shown in Fig. 1. The alloys consist mainly of

LaNi5 phase, La2Ni7 phase and minor LaNi3 phase from Jade-

5.0 analysis. The lattice parameters and cell volume of the

above phases are listed in Table 1. It is noticed that Ce-

substituted alloy has the smallest cell volume although rare

earth atoms’ radii decrease in the order La>Ce> Pr>Nd. This

may be ascribed to particular 4f electron structure of Ce

atoms, which makes Cerium be compacted more tightly by

other atoms, such as Ni atoms [16].

The back scattered electron (BSE) image of La0.60Nd0.20-

Mg0.20(NiCoMnAl)3.5 and EDX spectra for phases are shown

in Fig. 2. Three main areas with different colors can be seen

in Fig. 2(a). The EDX analyses in Fig. 2(b)–(d) show that the

light grey (a), dark grey (b) and black area (g) correspond to

AB3 (La0.62Nd0.16Mg0.22Ni2.58Mn0.09Co0.19Al0.11) phase, A2B7

(La0.72Nd0.17Mg0.11Ni3.04Mn0.12Co0.23Al0.26) phase and AB5

Table 1 – Lattice parameters and cell volume of LaNi5phase and La2Ni7 phase.

Samples LaNi5 phase La2Ni7 phase

a (A) c (A) V (A3) a (A) c (A) V (A3)

R¼ La 5.032 4.015 88.05 5.057 24.71 547.2

R¼Ce 5.012 3.986 86.68 5.043 24.68 543.5

R¼ Pr 5.013 3.994 86.90 5.059 24.54 543.9

R¼Nd 5.017 4.009 87.37 5.071 24.55 546.8

i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 1 3 9 9 – 1 4 0 4 1401

(La0.78Nd0.18Mg0.03Ni3.99Mn0.19Co0.36Al0.33) phase, respectively.

The results are consistent with those obtained from XRD.

3.2. P–C isotherms

The electrochemical pressure–composition isotherms that

examine the discharge level of hydrogen in an anode at 298 K

are presented in Fig. 3. The plateaus of the hydrogen desorp-

tion isotherms increase by substituting Ce, Pr and Nd at La

sites. This is closely related with decrease in cell volume after

Fig. 2 – BSE image and EDX analysis of La0.60Nd0.20Mg0.20(NiCoMn

substitution. The Ce-substituted alloy exhibits much sloppier

plateau and higher equilibrium pressure in desorption PCT

owing to the smallest cell volume. The hydrogen storage

capacity (H/M ) obtained from desorption PCT is also shown in

Table 2. The capacity decreases after La is substituted by Ce, Pr

and Nd, which results mainly from decreasing cell volumes.

Among the four alloys, Ce-substituted alloy has the lowest

hydrogen storage capacity.

3.3. Electrochemical properties

The discharge capacities of La0.60R0.20Mg0.20(NiCoMnAl)3.5

(R¼ La, Ce, Pr, Nd) alloys at different cycle numbers are shown

in Fig. 4. It can be seen that the Ce- and Pr-substituted alloys

can be activated in 2 cycles; the Nd-substituted alloy and the

original alloy need 3 cycles to be activated. The maximum

discharge capacity (Cmax) listed in Table 2 decreases with Ce,

Pr and Nd substitution. The capacity of Ce-substituted alloy

decreases from 372 mAh g�1 to 340 mAh g�1, decreasing by

8.6%, however, the capacity of Nd-substituted alloy decreases

slightly, from 372 mAh g�1 to 370 mAh g�1.

Al)3.5 alloy (a) BSE image (b) a phase (c) b phase (d) g phase.

0.0 0.2 0.4 0.6 0.8 1.01E-4

1E-3

0.01

0.1

1

10

R = LaR = CeR = PrR = Nd

p eq(

MP

a)

H/M

Fig. 3 – Electrochemical desorption P–C isotherms of

La0.60R0.20Mg0.20(NiCoMnAl)3.5 alloy electrodes.

0 20 40 60 80 100100

150

200

250

300

350

400

0 2 4 6 8150

200

250

300

350

400

R = La

R = Ce

R = Pr

R = Nd

Dis

char

ge c

apac

ity

(mA

h g-1

)

Cycle number (n)

Fig. 4 – Discharge capacity of La0.60R0.20Mg0.20(NiCoMnAl)3.5

alloy electrodes vs. cycle number.

60

80

100 R = La

R = Ce

R = Pr

R = Nd

%)

i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 1 3 9 9 – 1 4 0 41402

Cycling stability of metal hydride electrode is an important

factor to cycle life of MH/Ni batteries. It can be seen from Fig. 4

that slope of cycle life profiles becomes smaller with Pr and Nd

substitution for La, indicating that the cycle life is improved.

In order to illustrate cycling stability, capacity retention rate at

the 100th cycle (S100) is calculated as the ratio of C100/Cmax, and

is listed in Table 2. Table 2 shows that Pr- and Nd-substitution

causes an increase in S100. Nd substitution improves cycle life

significantly with S100 increasing from 79.8 to 93.2%.

It has been proved that discharge capacity degradation is

caused by serious pulverization and oxidation of alloy elec-

trodes. Structure of LaNi3 and La2Ni7 phases are both

described as a stacking of [LaNi5] and [La2Ni4] slabs, and

volume expansion of [La2Ni4] slabs can reach up to 59–63% on

hydriding [17]. Moreover, in the anisotropic hydride, expan-

sion of La2Ni7 phase and LaNi3 phase in c-axis direction is

larger than that in a-axis direction [18], causing serious inner

stress and thus easy pulverization. Therefore, pulverization is

an important factor to capacity degradation of R–Mg–Ni-based

alloys. Meanwhile, pulverization accelerates oxidation and

diffusion of rare earth elements towards the surface. During

charge/discharge process, rare earth elements tend to diffuse

towards alloy surface, and the disproportion is determined by

atom concentration, melting point and cell volume. Among

the four elements Ce diffuses to surface more quickly because

of low melting point and small volume, however, Nd diffuses

slowly due to its much higher melting point [19]. That is Nd

substitution can suppress oxidation of the alloy elements and

Table 2 – Electrochemical properties ofLa0.60R0.20Mg0.20(NiCoMnAl)3.5 alloy electrodes.

Samples H/M Cmax

(mAh g�1)Na HRD1200

(%)C233

(mAh g�1)S100

(%)

R¼ La 0.972 372 3 22.1 90 79.8

R¼Ce 0.907 340 2 61.3 220 79.6

R¼ Pr 0.957 363 2 49.6 141 81.5

R¼Nd 0.968 370 3 44.3 60 93.2

help to keep original structure of the alloys. As a result, the

Nd-substituted alloy has better cycling stability.

The high rate dischargeability (HRD) of La0.60R0.20Mg0.20

(NiCoMnAl)3.5 (R¼ La, Ce, Pr, Nd) alloys is shown in Fig. 5. The

HRD can be calculated by the following expression:

HRDð%Þ ¼ Cd

C60� 100

where Cd is discharge capacity at the discharge current

density Id (mAh g�1) and C60 is the discharge capacity at

60 mA g�1 (mAh g�1).

Ce, Pr and Nd substitutions for La are beneficial to HRD of

the alloys. High rate dischargeability at discharge current

density of 1200 mA g�1 (HRD1200) is listed in Table 2. Ce

substitution brings prominent improvement in HRD, and

HRD1200 of Ce-substituted alloy reaches 61.3%, which is 177%

higher than that of the original alloy (22.1%).

The discharge capacities of La0.60R0.20Mg0.20(NiCoMnAl)3.5

(R¼ La, Ce, Pr, Nd) electrodes at low temperature are

0 200 400 600 800 1000 1200 1400 16000

20

40

HR

D (

Discharge current density (mA g-1)

Fig. 5 – HRD of La0.60R0.20Mg0.20(NiCoMnAl)3.5 alloy

electrodes.

230 240 250 260 270 280 290 3000

100

200

300

400

R = La

R = Ce

R = Pr

R = Nd

Dis

char

ge c

apac

ity

(mA

h g-1

)

Temperature (K)

Fig. 6 – Discharge capacity of La0.60R0.20Mg0.20(NiCoMnAl)3.5

alloy electrodes at low temperature.

Table 3 – Electrochemical kinetics parameters ofLa0.60R0.20Mg0.20(NiCoMnAl)3.5 alloy electrodes.

Samples Polarizationresistance,

Rp (mU)

Exchangecurrent density,

I0 (mA g�1)

Hydrogendiffusion

coefficient, D(�10�11 cm2 s�1)

R¼ La 152.1 168.8 3.68

R¼Ce 129.2 198.7 12.1

R¼ Pr 119.5 214.8 6.55

R¼Nd 138.0 186.0 3.23

i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 1 3 9 9 – 1 4 0 4 1403

presented in Fig. 6. The discharge capacity decreases with

decreasing temperature. The discharge capacities at 233 K

(C233) listed in Table 2 are 90 mAh g�1, 220 mAh g�1,

141 mAh g�1 and 60 mAh g�1, and they are 25.6%, 68.8%, 42.5%

and 16.5% of the capacities at 298 K in the order of La, Ce, Pr

and Nd, respectively.

Iwakura et al. [20] have pointed out that high rate dis-

chargeability and low temperature dischargeability are

controlled by kinetics, including charge transfer rate at

surface and hydrogen diffusion rate in bulk. The micro-

polarization curves of La0.60R0.20Mg0.20(NiCoMnAl)3.5 (R¼ La,

Ce, Pr, Nd) alloy electrodes are plotted in Fig. 7. From Fig. 7

polarization resistance (Rp) can be obtained and consequently

exchange current density (I0) can be calculated according to

equation [21]

I0 ¼RTFRp

where I0 is the exchange current density (mA g�1); R is the gas

constant (J mol�1 K�1); T is the absolute temperature (K); F is

-6 -5 -4 -3 -2 -1 0 1 2 3 4 5 6-50

-40

-30

-20

-10

0

10

20

30

40

50

R = La

R = Ce

R = Pr

R = NdCur

rent

den

sity

(m

A g

-1)

Overpotential (mV)

Fig. 7 – Micropolarization curves of La0.60R0.20Mg0.20-

(NiCoMnAl)3.5 alloy electrodes at 50% DOD.

the Faraday constant (C mol�1) and Rp is the polarization

resistance (mU).

The polarization resistance and the exchange current

density are both presented in Table 3. The Ce- and Pr-

substituted alloys have much larger exchange current density

I0, indicating their higher charge transfer rate at electrode

surface.

For the hydrogen diffusion coefficient test, potentialstatic

step method was used. It can be seen from Fig. 8 that each

semilogarithmic curve of anode current density vs. time can

be divided into two time regions. In the first time domain, the

current decreases rapidly due to consumption of hydrogen on

the surface; while in the second time domain, the current

decreases slowly with a linear tendency. In this case,

hydrogen is supplied from the bulk alloy proportionally to the

concentration gradient of hydrogen and the current is

controlled by the diffusion rate of hydrogen atom. From slope

of the liner region in Fig. 8, hydrogen diffusion coefficient D

can be calculated by the following equation [22].

log i ¼ lg

�6FDda2ðC0 � CsÞ

�� p2

2:303Da2

t

where i is the anode current density (mA g�1), D is the

hydrogen diffusion coefficient (cm2 s�1), d is the density of the

alloy (g cm�3), a is the average radius of the alloy particles, C0

is the initial hydrogen concentration in the alloy (mol cm�3),

Cs is the surface hydrogen concentration of the alloy

0 500 1000 1500 2000 2500 3000-2.8

-2.4

-2.0

-1.6

-1.2

-0.8

-0.4

0.0

log(

i(A

g-1))

Time (s)

R = La

R = Ce

R = Pr

R = Nd

Fig. 8 – Correspondence of anode current density vs. time

of La0.60R0.20Mg0.20(NiCoMnAl)3.5 alloy electrodes at

D500 mV.

i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 1 3 9 9 – 1 4 0 41404

(mol cm�3) and t is the discharge time (s). Taking a as 13 mm, D

can be calculated according to the equation, and the results

are presented in Table 3. It can be seen that the Ce-substituted

alloy has the largest hydrogen diffusion coefficient D.

4. Conclusions

The effect of rare earth elements on phase structure and

electrochemical properties of La–Mg–Ni-based hydrogen

storage alloys has been investigated in the paper present. The

main results can be summarized as follows.

(1) The XRD patterns and SEM–EDS analyses show that the

alloys consist mainly of LaNi5 phase, La2Ni7 phase and

minor LaNi3 phase. The cell volumes of LaNi5 and La2Ni7phase in the Ce-substituted alloy decrease the most.

(2) The electrochemical capacity decreases and equilibrium

pressure increases with Ce, Pr and Nd substitution for La.

However, Pr and Nd substitution improves the cycling

stability. The capacity retention rate at the 100th cycle of

Nd-substituted alloy rises by 13.4%.

(3) Ce and Pr substitution greatly improves hydrogen diffu-

sion rate in the bulk and charge transfer rate at the surface,

and thus ameliorates the high rate dischargeability and

low temperature dischargeability. The HRD at 1200 mA g�1

and capacity at 233 K reach 61.3% and 220 mAh g�1, 177%

and 154% higher than those of the original alloy for the

Ce-substituted alloy, respectively.

Acknowledgements

This work was financially supported by the National Natural

Science Foundation of China (20673093) and the Natural

Science Foundation of Hebei Province (B2007000303).

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