Article
pubs.acs.org/JPCC
Mechanistic and Kinetic Study of the Electrochemical Charge and
Discharge of La2MgNi9 by in Situ Powder Neutron Diffraction
Michel Latroche,*,† Fermìn Cuevas,† Wei-Kang Hu,‡ Denis Sheptyakov,§ Roman V. Denys,‡
and Volodymyr A. Yartys‡,∥
†
Institut de Chimie et des Matériaux Paris-Est, UPEC−CNRS, UMR 7182, 2 rue Henri Dunant, 94320 Thiais, France
Institute for Energy Technology, Department of Energy Systems, 2007 Kjeller, Norway
§
SINQ-HRPT, Paul Scherrer Institute, 5232 Villigen, Switzerland
∥
Norwegian University of Science and Technology, 7491, Trondheim, Norway
‡
S Supporting Information
*
ABSTRACT: The intermetallic La2MgNi9 has been investigated as negative electrode
material for NiMH battery by means of in situ neutron powder diffraction. This hydrideforming compound exhibits suitable plateau pressures ranging within the practical
electrochemical window and leads to significant reversible electrochemical capacities. Charge
and discharge of the composite electrode have been performed in beam following various
current rates and galvanostatic intermittent titration. From the diffraction data analysis, phase
amounts and cell volumes have been extracted, allowing the interpretation of the hydride
formation and decomposition. From the evolution of the diffraction line widths, differences
are observed between charge and discharge with the possible formation of an intermediate γ
phase on charge. The electrode readily responds to current rate variations and does not show
any kinetic limitation in the range C/10 and C/5 (C/n: full capacity C in n hours). This
material shows excellent properties regarding electrochemical storage of energy.
■
strains induced by the significant volume change between the
intermetallic and its hydride.
However, despite all these efforts to design materials with
enhanced properties, the use of LaNi5-type compounds remains
limited by their low capacities, around 320 mAh g−1. This is
mainly due to the heavy rare earths that lead to high molar
mass. To overcome this latter point, Kohno et al.12 developed
materials for which part of the heavy rare earths can be replaced
by light element, such as magnesium. Indeed, it exists in several
phases in the Ni-rich part of the binary La−Ni diagram.
Between 75% and 79.2% of nickel, line compounds LaNi3,
La2Ni7, and La5Ni19 can be described as a stacking along the c
axis of [La2Ni4] + n[LaNi5] subunits with n = 1, 2 and 3,
respectively.13,14 Mg does not substitute La in the LaNi5
subunits, but it replaces the rare earth within the La2Ni4
subunits. This led to a real breakthrough in terms of increased
gravimetric H densities as magnesium allows significant
decrease of the molar mass. On the basis of the described
concept, new materials with various rare earths and transition
metals have been developed offering significant electrochemical
storage capacities up to 400 mAh g−1.2−4,15
Among them, La3−xMgxNi9 are key compounds as they
correspond to a simple stacking [La2−xMgxNi4] + [LaNi5] (x =
0−2),14 allowing Mg-rich phases. LaMg2Ni9 (2 [MgNi2] +
INTRODUCTION
Metallic hydrides are very convenient materials for hydrogen
storage.1 They can store a large amount of hydrogen gas at
thermodynamic conditions close to ambient pressure and room
temperature, making them suitable for energy storage. Metallic
hydrides have been also developed for electrochemical storage.
Indeed, they can react with water to electrochemically exchange
one proton and one electron. This allows the design of efficient
negative electrodes for alkaline secondary batteries (NiMH)
which offer better properties that the actual NiCad batteries.2−4
For two decades, the most developed materials for these
negative electrodes were derived from the LaNi5 intermetallic
compound. By proper substitutions on either the La site (by
other rare earths) or the Ni sites (by other transition or pmetals) and by playing with the stoichiometry (according to the
narrow domain of existence of the LaNi5+x compound; 0 < x <
0.4), a complex but efficient class of materials has been
developed.5−10
According to the thermodynamic rules, during charge and
discharge, reactions at the negative electrode should occur only
between two solid phases, the intermetallic (α) and the hydride
(β). However, it was evidenced by several authors that for the
best materials in term of cycle life, the appearance of an out-ofequilibrium, metastable, transient γ phase is observed during
charge.6,8,11 This unexpected γ phase (with intermediate
volume and H-loading) was reported to form at the interface
between the α and β phases, helping to accommodate the heavy
© 2014 American Chemical Society
Received: April 1, 2014
Revised: May 19, 2014
Published: May 19, 2014
12162
dx.doi.org/10.1021/jp503226r | J. Phys. Chem. C 2014, 118, 12162−12169
The Journal of Physical Chemistry C
Article
porosity of the electrode. The potential was monitored against
a solid-state Cd/Cd(OH)2 reference electrode.5,9
NPD data were recorded at room temperature and ambient
pressure with the high-resolution powder diffractometer for
thermal neutrons (HRPT) in high intensity mode (flux of
about 1014 n cm−2 s−1) at SINQ, PSI in Switzerland.
Using the large position sensitive (PSD)3He detector,
measurements were performed in the scattering angle ranging
from 5 to 165° with angular step of 0.05°. Wavelength was set
to 1.494 Å. The patterns were recorded by batches of eight,
varying the detector angle by step of 0.15° between 3.95° and
5°. Typical time acquisition was 225 s per pattern. The eight
patterns were finally combined into one signal file leading to a
time resolution of about 1800 s for each diffraction measurement. Because of the high flux of the spallation source and the
use of a large position-sensitive detector, the high-rate
capability of such electrodes can be followed in situ.
The diffraction patterns were sequentially refined using the
program FULLPROF.18
Preliminary Work. The composite electrode was first
charged ex situ before the NPD experiment at 20 mA g−1 (115
mA or C/16) up to a total capacity of 2300 mAh. Three
activation cycles were performed with a cutoff potential of 200
mV versus Cd/Cd(OH)2. The discharge capacities measured at
each cycles are given in Table 1.
[LaNi5]) can be considered to be the parent compound of the
La3−xMgxNi9 family. Earlier synchrotron X-ray diffraction and
neutron powder diffraction (NPD) studies revealed that
substitution of La by Mg in LaNi3 proceeds in an ordered
way and leads to the formation of LaMg2Ni9 at maximum
solubility of magnesium. Gradual increase in Mg content is
accompanied by a linear decrease of the unit cell volume.
Hydrogen interaction with the La3−xMgxNi9 alloys was
investigated by in situ neutron powder diffraction and
pressure−composition−temperature (PCT) studies.13 Magnesium was found to influence structural features of the
hydrogenation process and determine various aspects of the
hydrogen interaction with the La3−xMgxNi9 intermetallics
causing: (i) more than 1000-fold increase in equilibrium
pressures of hydrogen absorption and desorption for the Mgrich LaMg 2Ni 9 as compared to that of the Mg-poor
La2.3Mg0.7Ni9, which implies a substantial modification of the
thermodynamics of the formation and decomposition of the
hydrides according to the van’t Hoff relation; (ii) an increase of
the reversible hydrogen storage capacities following increase of
Mg content in the La1−xMgxNi3 to ∼1.5 wt % H (380 mAh g−1
in equivalent electrochemical units) for La2MgNi9; (iii)
improvement of the resistance against hydrogen-induced
amorphization and disproportionation; (iv) change of the cell
expansion mechanism of the hydrogenation from anisotropic to
isotropic. Thus, optimization of the magnesium content
provides different possibilities for improving properties of the
studied alloys as hydrogen storage and battery electrode
materials.
The present work deals with the electrochemical formation
and decomposition of the hydride of the representative parent
La2MgNi9 alloy for battery electrode applications. Studies of the
interaction with gaseous hydrogen revealed that the hydride
forms at pressures slightly below one bar, which is suitable for
electrochemical formation and decomposition in open cells.
Indeed, the electrochemical discharge capacity of such alloys is
larger than that for LaNi5-based electrodes.16 Importantly,
La2MgNi9 alloy can be synthesized directly from the melt by
using the rapid solidification technique.17 The outcome of the
present study is the optimization of the chemical composition
of ternary La−Mg−Ni compounds with respect to their
electrochemical hydrogen storage performance.
Table 1. Discharge Capacities at Each Cycle during
Activation of the Composite Electrode
cycle number (C/16)
capacity
first cycle
second cycle
thrid cycle
949 mAh (177 mAh g−1)
1257 mAh (234 mAh g−1)
1375 mAh (256 mAh g−1)
From the pressure−composition−temperature (PCT) isotherm curve of La2MgNi9 measured at 20 °C (Figure 1; data
from ref 13), the electrochemical capacities can be estimated
according to the practical pressure windows. This defines about
310 mAh g−1 of reversible capacity (between 0.001 and 0.1
EXPERIMENTAL SECTION
The raw metallic material La2MgNi9 was prepared by powder
metallurgy at IFE, Norway.13 From this starting material, a
composite electrode was made from intermetallic powder
sieved below 63 μm and mixed with carbon black and PTFE in
the weight ratio 90:5:5. This mixture was spread out in sheets
approximately 1.5 mm thick and compressed at 5 tons on both
sides of a 5 × 3.14 cm2 nickel grid which plays the role of
current collector. The final thickness of the composite electrode
was about 0.9 mm. These electrode plates were then rolled up
on themselves to form a cylinder of about 50 mm height and 10
mm diameter. The final working electrode contains 5.366 g of
active material and was introduced in a specially designed silica
cell. The electrode is sandwiched between inner (⌀ 8 mm) and
outer (⌀ 12 mm) counter-electrode cylinders made of nickel
grid rolled on themselves, with silica sheaths as separators on
each side of the working electrode. The electrode is immersed
in NaOD 5.5 M electrolyte and pumped under primary vacuum
in order to fully impregnate the working electrode with the
liquid and to remove any gaseous species trapped in the
■
Figure 1. Pressure−composition−temperature (PCT) isotherm curve
for La2MgNi9 during a full absorption−desorption cycle at 20 °C. The
dotted lines stand for the electrochemical window (between 0.1 and
0.001 MPa) and define about 310 mAh g−1 of reversible capacity (data
from ref 13), whereas the total capacity at 0.1 MPa reaches 380 mAh
g−1.
12163
dx.doi.org/10.1021/jp503226r | J. Phys. Chem. C 2014, 118, 12162−12169
The Journal of Physical Chemistry C
Article
543.65 Å3), as the low quantity of this phase and the bulky
pattern do not allow accurate refinement of these parameters.
These values are higher than those of the La 2 MgNi 9
intermetallic (a = 5.0314(2); c = 24.302(1) Å; V = 532.78
Å3; 1.2 Å3/f.u. La2MgNi9)19 and indicate the formation for the
active grains of the H/D solid solution containing about 0.5−1
atom H/D/f.u. La2MgNi9. This is consistent with the PCT
diagram shown in Figure 1.
MPa) and 380 mAh g−1 of total capacity (between 0 and 0.1
MPa). Thus, the reversible capacity obtained for the third
activation cycle represents about 82% of the expected reversible
electrochemical capacity.
The electrode was then fully charged just before the in situ
experiment and was measured once in beam under open circuit
voltage (OCV). The diffraction data were interpreted on the
basis of the structural parameters given for La2MgNi9 (i.e., the
phase α)13 and La2MgNi9D13 (i.e., the phase β)19 (see
Supporting Information for details). Beside the strong nickel
lines coming from the current collector and the counterelectrodes, the diffraction lines belonging to the main hydride
phase (i.e., the phase β) are clearly observed (Figure 2). Some
EXPERIMENTAL RESULTS
The electrode was first discharged at a constant current of 180
mA (33 mA g−1 or D/10) with a cutoff potential of 0.5 V (vs
ECd/Cd(OH)2). The total discharged capacity was 1494 mAh (278
mAh g−1). During the process, starting from 0.044 V, an
increase of the potential of the working electrode is observed.
An unexpected kink appears after 3 h of discharge at 0.158 V
(Figure 3) but was not observed during the previous or
■
Figure 2. Refined neutron powder diffraction pattern (measured, dots;
calculated, solid line) for the electrode at initial charged state.
Crystallographic structures are taken from Denys and Yartys.13,19
Vertical bars correspond to diffraction line positions for each phase:
La2MgNi9 (α) and deuterated La2MgNi9D13 (β) phases. Nickel lines
(heavily textured) arise from the current collector of the working
electrode and from the counter electrodes. Background around 30°
comes from the silica cell and the NaOD/D2O electrolyte. One can
also note an extra peak at 2θ = 17.5° that is attributed to PTFE.
Figure 3. Evolution of the potential Ew versus the Cd/Cd(OD)2
reference electrode during the first in-beam charge−discharge (C/10)
of the electrode at a current rate of 33 mA g−1.
weak lines coming from the deuterium-free phase (i.e., the
phase α) can also be observed, and this phase was also
introduced in the refinement. The presence of this phase
indicates that a small amount (around 14 wt %) of the
intermetallic compound is electrochemically nonactive because
of loss of electronic or electrolytic contacts for a few grains.
This partly explains the 82% observed for the reversible
capacity at the third cycle.
The cell parameters of the β phase are 5.3692 and 26.1328 Å
for a and c, respectively (space group R3m
̅ ), and the cell
volume is 652.42 Å3. This is slightly lower than the volume
reported by Denys and Yartys13,19 (675.11 Å3) for the fully
charged hydride under gas pressure (13 D/f.u.), but this latter
value was obtained under 1 MPa of gas pressure, 1 order of
magnitude larger than the present ambient conditions. The
volume decrease of 22.7 Å3 or 2.5 Å3/f.u. La2MgNi9 can be
translated into the loss of hydrogen storage capacity of
approximately 0.8 at H/f.u. from La 2 MgNi 9 D 13 to
La2MgNi9D12. This volume shrinking is obviously associated
with a depopulation of the D sites occupied in La2MgNi9D13.19
However, for the simplicity of the evaluation of the data, only
changes of the unit cell parameters where accounted, whereas
the atomic structure was considered as unaltered.
The cell parameters of the α phase were kept fixed to 5.0712
and 24.4105 Å for a and c, respectively (space group R3̅m; V =
following measurements. It is then assumed that it is probably
an artifact (related to bubble formation close to the reference
capillary). This is supported by the analysis of the EMH−ENi
signal (not shown here), which shows a flat plateau during the
whole discharge. After 8.5 h of discharge, the potential rises
rapidly to reach 0.500 V and the current was switched off. The
potential then decreases rapidly to stabilize around 0.030 V
during the rest period (OCV).
Following 1 h of rest (OCV), the current was then set to
−180 mA (33 mA g−1 or C/10) for 12 h assuming full charge of
the electrode over this time scale. Potential decreases rapidly to
reach a plateau around −0.163 V. It remains nearly constant
though slightly decreasing and becoming bulky after 7 h (i.e., t
= 16 h) of charge (Figure 3). This is attributed to hydrogen
evolution leading to large production of deuterium gas bubbles
close to the reference electrode capillary and thus involving a
noisy signal.
The 3D view of the NPD pattern evolution as a function of
time during the first cycle (D/10 + C/10) is shown in Figure 4.
One can observe that the strong diffraction lines belonging to
the β phase decrease rapidly during the discharge (time 0−8.5
h) whereas those of the α phase increase. After 9.5 h, the
transformation is almost fully completed and the charge was
12164
dx.doi.org/10.1021/jp503226r | J. Phys. Chem. C 2014, 118, 12162−12169
The Journal of Physical Chemistry C
Article
Figure 5. Comparison between the α and β phase amounts and the
capacity Q during the in-beam charge−discharge cycle at D/10 + C/10
for the working electrode. The initial charge state Q at t = 0 is set at
380 mAh g−1, a value determined from the PCT curve (Figure 1) at
0.1 MPa.
Figure 4. 3D view of the NPD pattern evolution as a function of time
during the first cycle (D/10 + C/10) of the working electrode at a
current rate of 33 mA g−1.
• from 9.5 to 12 h, a region attributed to the α phase solid
solution domain for which phase amounts do not vary
significantly;
• from 12 to 16 h, a two-phase domain where the α and β
phases are in equilibrium and transform into each other
rapidly;
• beyond 16 h, a region where the β phase solid solution
domain starts but in competition with the hydrogen
evolution. Then a slow increase of the β phase amount is
observed despite the linear augmentation of Q.
From the analysis of the diffraction patterns, the evolution of
the cell volumes of both phases can be plotted as a function of
time (Figure 6). During the discharge, the volume of the β
phase decreases continuously, which indicates that the phase is
progressively depleted in deuterium. Assuming that ΔV/atom
D = 3.26 Å3 (as shown later in the paper), we can estimate that
the lower content of deuterium in the β-phase reached during
started. The α phase then transforms reversibly into the β
phase, but contrary to the discharge, strong overlap of the
diffraction peaks of both phases is observed. After about 20 h,
the charge is completed and a full recovery of the diffraction
peaks of the β phase takes place.
For the whole cycle, the diffraction patterns have been
refined sequentially assuming three phases as observed in the
first diffraction pattern (Figure 2): α deuterium-free and β
deuteride La2MgNi9 phases plus the nickel one from the
counter electrodes. Only scale factors, cell parameters, and the
U parameters of the Caglioti function were refined (8
parameters). In addition, as the level of the background can
follow slight variations depending on the amount of produced
deuterium gas (leading to various amounts of NaOD/D2O
electrolyte in the beam), the 25 points of the interpolated
background were also refined leading to a total of 33 refined
parameters.
From analysis of these data, the relative amount of each
phase (α and β) can be extracted. The amount of Ni was
considered to be a constant. The results are shown in Figure 5
and are compared to the electrochemical capacity Q passed
through the electrode (right-hand scale). However, one should
keep in mind that during the charge step, because of hydrogen
(deuterium) evolution, the capacity Q does not reflect the exact
state of charge of the working material. To overcome this
difficulty, the fully charged state was set at 380 mAh g−1, a value
determined from the PCT curve (Figure 1) at 0.1 MPa, i.e., at
the top of the electrochemical window.
The results of the galvanostatic cycle can be split into
different domains.
During the discharge:
• from 0 to 3 h, a first step attributed to the β phase solid
solution domain (i.e., the formation of the α phase is
hardly detected);
• from 3 to 8.5 h, a two-phase domain where the β and α
phases are in equilibrium and transform into each other;
• from 8.5 h to 9.5 h, a relaxation period for which phase
amounts remain nearly constant.
During the charge:
Figure 6. Evolution of the cell volumes for the α and β phases during
the in-beam charge−discharge cycle at C/10 for the working electrode.
For sake of comparison, the evolution of the potential Ew versus the
Cd/Cd(OD)2 reference electrode is also shown on the right-hand
scale.
12165
dx.doi.org/10.1021/jp503226r | J. Phys. Chem. C 2014, 118, 12162−12169
The Journal of Physical Chemistry C
Article
its discharge is close to La2MgNi9D11.2. This is much less the
case for the α phase for which the cell volume is nearly constant
though slightly decreasing.
During the charge, the determination of the cell volumes is
less accurate, especially during the α-to-β transformation (t =
12−16 h) when diffraction peaks from both phases strongly
overlap (see 3D patterns in Figure 4). Nevertheless, one can
observe that both volumes increase rapidly during the first 6 h
of charge and then stabilize progressively though the β phase
volume continues to increase until the end of the charge. This
indicates that the charge still takes place in the β solid solution
domain. At the end of the charge, the β phase volume recovers
its initial value (t = 0).
Finally, the evolution of the diffraction peak broadening, as
determined by the fitting of a common U parameter for both
phases, is given in Figure 7. For the discharge and the charge,
Figure 8. Evolution of the phase amounts for the α and the β phases
during the in-beam charge−discharge cycle at C/5 for the working
electrode. For sake of comparison, the evolution of the potential Ew
versus the Cd/Cd(OD)2 reference electrode is also shown on the
right-hand scale.
Figure 7. Evolution of the half-width of the diffraction peaks for the β
and α phases during the in-beam charge−discharge cycle at D/10+C/
10 for the working electrode. For sake of comparison, the evolution of
the capacity Q is also shown on the right-hand scale. The initial charge
state Q at t = 0 is set at 380 mAh g−1, a value determined from the
PCT curve (Figure 1) at 0.1 MPa.
Figure 9. 3D view of the NPD pattern evolution as a function of time
during the GITT discharge cycle (D/7) of the working electrode at 46
mA g−1.
an enlargement of the peaks is seen for the β and α phases
when they are in equilibrium and transform into each other
(i.e., crossing the plateau region). However, this effect is much
larger (1 order of magnitude) during the charge, leading to
peak overlapping as observed in Figure 4.
For the next 10 h, the electrode was cycled again with a
higher current density (75 mA g−1 or ∼C/5). The discharge
capacity obtained at this rate is 1332 mAh (248 mAh g−1).
Except for this slightly lower capacity, the electrode behaves
very similarly to the previous cycle at C/10 involving the same
step: α solid solution, α-to-β transformation, β solid solution,
and hydrogen evolution at the end of the charge. In the same
way, cell volume variations and half-width evolutions are very
comparable to those observed at C/10. Phase amounts and
potential evolutions are given in Figure 8.
Finally, the electrode was discharged by the galvanostatic
intermittent titration technique (GITT) using the following
protocol: discharge at 46 mA g−1 (D/7) for 1.5 h or Ew > 0.5 V
followed by a relaxation period of 1 h. Evolution of the
diffraction patterns is shown in Figure 9. The procedure lasts
for six steps that are summarized in Table 2 giving discharge
time and cumulated capacities.
Table 2. Cumulated Discharge Capacities Obtained from
GITT Experimenta
step
discharge time (h)
capacity (mAh)
capacity (mAh g−1)
1
2
3
4
5
6
1:30
1:30
1:30
1:02
0:10
0:08
370
740
1100
1350
1392
1427
69
138
207
252
259
266
a
The current was set at 46 mA g−1 (D/7) for 1.5 h or Ew > 0.5 V
followed by relaxation periods of one hour.
One can observe in the 3D view of Figure 9 that the
diffracted intensities of both phases follow fairly well the
different current steps of the GITT. In addition, one can note
that the background level is very dependent on the current
density. This can be understood if one considers that switching
on and off the current involves important changes in the
amount of deuterium gas produced in the cell. When the
current is off, no gas is produced and the quantity of liquid
12166
dx.doi.org/10.1021/jp503226r | J. Phys. Chem. C 2014, 118, 12162−12169
The Journal of Physical Chemistry C
Article
(D2O, NaOD) is larger in the beam, causing increased
background. Thus, each maximum in the background can be
attributed to a relaxation period.
The total cumulated discharge capacity obtained at this rate
is 1427 mAh (266 mAh g−1), a value very close to that
measured during the first in situ cycle at C/10. The evolution of
the relative amount of each phase (α and β) is shown in Figure
10. Again, the electrode behaves very closely to the previous
cycles involving the same transformations.
Figure 11. Evolution of the phase amounts for the α and the β phases
during the in-beam charge−discharge cycles at C/10, C/5, and GITT
for the working electrode. For sake of comparison, the evolution of the
potential Ew versus the Cd/Cd(OD)2 reference electrode is also shown
on the right-hand scale.
Figure 10. Evolution of the phase amounts for the α and the β phases
during the in-beam discharge obtained by GITT experiments. The
current was set to 46 mA g−1 (D/7) for 1.5 h or Ew > 0.5 V followed
by relaxation periods of 1 h. For sake of comparison, the evolution of
the capacity Q is also shown on the right-hand scale. The initial charge
state Q is set at 380 mAh g−1, a value determined from the PCT curve
(Figure 1) at 0.1 MPa.
In addition, small potential plateaus appear during the OCV
periods which correspond to zero current flow for which no
phase transformation takes place. However, very small changes
in the α and β amounts can be observed during the resting
periods (Figure 10). This is consistent with deuterium diffusion
from the β toward the depleted α phase on the particle surface,
transforming some β phase into (saturated) α phase during the
electrode relaxation.
The charge and discharge capacities for each in-beam cycle
described so far are given in Table 3.
Figure 12. Evolution of the cell volumes for the α and the β phases
during the in-beam charge−discharge cycles at C/10, C/5, and GITT
for the working electrode. For sake of comparison, the evolution of the
potential Ew versus the Cd/Cd(OD)2 reference electrode is also shown
on the right-hand scale.
electrochemical capacity does not correspond to that of the
electrode materials, mainly because of the deuterium gas
evolution at the end of each charge. To overcome this difficulty,
data can be analyzed by separately considering capacity of
charge Qc and capacity of discharge Qd. This has been done in
Figure 13 for the cell volume evolution of both phases.
Interestingly, a linear behavior is observed for the two phases
though the volumes measured during charge are a bit scattered.
The β phase volume increases smoothly as a function of Q
following the equation
Table 3. Discharge Capacity at Each Cycle during in Situ
Measurements of the Working Electrode
4th cycle (D/10)
5th cycle (D/5)
6th cycle (GITT; D/7)
capacity (mAh)
capacity (mAh g−1)
1494 mAh
1332 mAh
1427 mAh
(278 mAh g−1)
(248 mAh g−1)
(266 mAh g−1)
To get an overview of the electrode behavior, the evolutions
of phase amounts and cell volumes for the α and the β phases
during the in-beam charge−discharge cycles at C/10, C/5, and
GITT are given and compared to the potential Ew in Figures 11
and 12.
From the precedent data analysis, all measured parameters
(phase amounts and cell volumes) have been processed
regarding time scale evolution. At this stage, it is worth looking
also to the data as a function of the electrochemical charge Q.
Once again, one has to be cautious with the values of Q as the
Vβ = 0.101(4)Q d + 608.5
(1)
whereas that of the α phase is nearly constant in the whole
range of Qd.
The same figure can also be drawn for the phase amounts,
though it is significant only for the discharge. This evolution is
shown in Figure 14 for which two domains can be clearly
identified. From 100 to 310 mAh g−1, a clear two-phase
transition domain can be seen. Then, for Qd larger than 310
12167
dx.doi.org/10.1021/jp503226r | J. Phys. Chem. C 2014, 118, 12162−12169
The Journal of Physical Chemistry C
Article
Using neutron diffraction data analysis, we observed that the
structural properties of the charged electrode material are
comparable to that published by Denys et al.19 for the hydride
loaded under gas pressure. Starting from this charged state, the
electrochemical cycling behavior has been investigated in beam
at different rates to follow the mechanism of the reversible
charge−discharge process.
At first glance, one can consider that the electrode material
behaves like a classic one following an intermetallic-to-hydride
transformation through the sequences α, α-to-β, and β
formation, a typical behavior for a two-phase reaction.
This is indeed the case during the discharge according to the
variations of the phase amounts, the cell volumes, and the line
diffraction half-widths. All these parameters agree well with a
two-phase behavior. This is, however, not so obvious for the
charge process. Despite the fact that a two-phase transformation seems to take place in both processes, important
increase of the half width during charge involves strong
overlapping between the α and β diffraction peaks. This
unexpected behavior can be attributed to a large concentration
gradient in the H concentration or heavy constraints during the
charge process. However, this is not supported by the small
hysteresis between the absorption and desorption branches or
by the relatively flat plateaus of the PCT curves. A better
hypothesis might be the formation of an intermediate γ phase
as previously observed in the LaNi5-type system.6,8,11 Similarly,
the γ phase was observed only during the charge process. This
out-of-equilibrium phase was related to constraint relaxation at
the α-to-β interface and plays a key role in the cycle life by
decreasing the decrepitation process responsible for the high
corrosion rate in alkaline medium. As the PuNi3-type structure
is built from the stacking of LaNi5 and MgNi2 subunits, such a
mechanism may also take place in this system. A better
characterization of this intermediate phase could be obtained in
the future by performing in-beam GITT during charge with
longer relaxation times assuming that the γ phase amount will
persist at the phase interfaces because of kinetic barriers. To
our knowledge, the formation of this intermediate γ phase was
never reported before in these stacking structures.
From Figure 13, it is observed that the cell volume for the β
phase is strongly dependent on the state of charge and follows a
linear dependence. From eq 1, one can derive that this
corresponds to 3.26 Å3/D atom, a value in agreement with
previously reported values ranging between 2.5 and 3.5 Å3/H
atom.20,21 On the contrary, the volume of the α phase remains
almost constant at all states of charge. Evolution of the phase
amounts as a function of the state of discharge allows to clearly
discriminate between β solid solution, β to α transformation
and α solid solution domains (Figure 14). This latter domain is
very small in the range of the electrochemical state of charge
studied in the present work, which explains the little variation
of the cell volume of the α phase.
In other words, the reversible capacity is mainly obtained
from hydrogen absorption−desorption in the β solid solution
and at the α-to-β transformation but not in the α solid solution.
This statement agrees with the shape of the PCT curve (Figure
1) for which the hydrogen trapped in the α solid solution is
difficult to extract at low pressure, i.e., out of the practical
electrochemical window.
Finally, it is interesting to note that the electrode material
readily responds to the electrochemical solicitations. A typical
example is shown in Figure 10 where the diffraction peak
intensities follow very closely the current variation imposed by
Figure 13. Evolution of the cell volumes for the α (black squares) and
the β (red circles) as a function of the state of charge Qc (solid
symbols) and Qd (empty symbols). For Qd, the fully charged state was
set at 380 mAh g−1, a value determined from the PCT curve (Figure 1)
at 0.1 MPa.
Figure 14. Evolution of the phase amounts for the α (black squares)
and the β (red circles) phases as a function of the state of discharge Qd.
The fully charged state was set at 380 mAh g−1, a value determined
from the PCT curve (Figure 1) at 0.1 MPa.
mAh g−1, the β solid solution domain starts and extends up to
the upper charge state (380 mAh g−1). Interestingly, for the α
phase, an instant formation of the electrochemically stable solid
solution is observed. As it is not possible to electrochemically
decompose it, the domain of the α solid solution with a variable
content of H/D is effectively nonexistent. It is also worth
noting that a fully charged state (i.e., 100% β) is never reached,
whereas the fullly discharged state (i.e., 100% α) is almost
complete.
DISCUSSION
The intermetallic La2MgNi9 readily and reversibly absorbs
hydrogen by solid−gas route, and thanks to its suitable
equilibrium pressures, the same compound can be used as
negative electrode in alkaline medium. Indeed, after a few
activation cycles, an excellent correlation between the capacities
obtained by solid gas and electrochemical measurements is
observed.
■
12168
dx.doi.org/10.1021/jp503226r | J. Phys. Chem. C 2014, 118, 12162−12169
The Journal of Physical Chemistry C
Article
(4) Liu, Y.; Cao, Y.; Huang, L.; Gao, M.; Pan, H. Rare Earth−Mg−
Ni-based Hydrogen Storage Alloys as Negative Electrode Materials for
Ni/MH Batteries. J. Alloys Compd. 2011, 509, 675−686.
(5) Latroche, M.; Percheron-Guégan, A.; Chabre, Y.; Bouet, J.;
Pannetier, J.; Ressouche, E. Intrinsic Behavior Analysis of Substituted
LaNi5-type Electrodes by Means of in Situ Neutron Diffraction. J.
Alloys Compd. 1995, 231, 537−545.
(6) Latroche, M.; Percheron-Guégan, A.; Chabre, Y. Influence of
Cobalt Content in MmNi4.3−xMn0.3Al0.4Cox Alloy (x=0.36 and 0.69)
on its Electrochemical Behaviour Studied by In Situ Neutron
Diffraction. J. Alloys Compd. 1999, 293−295, 637−642.
(7) Cuevas, F.; Joubert, J.-M.; Latroche, M.; Percheron-Guégan, A.
Intermetallic Compounds as Negative Electrodes of Ni/MH Batteries.
Appl. Phys. A: Mater. Sci. Process. 2001, 72, 225−238.
(8) Latroche, M.; Chabre, Y.; Percheron-Guégan, A.; Isnard, O.;
Knosp, B. Influence of Stoichiometry and Composition on the
Structural and Electrochemical Properties of AB5+y-based Alloys Used
as Negative Electrode Materials in Ni-MH Batteries. J. Alloys Compd.
2002, 330−332, 787−791.
(9) Latroche, M.; Chabre, Y.; Decamps, B.; Percheron-Guégan, A.;
Noréus, D. In Situ Neutron Diffraction Study of the Kinetics of
Metallic Hydride Electrodes. J. Alloys Compd. 2002, 334 (1−2), 267−
276.
(10) Latroche, M.; Joubert, J.-M.; Guegan, A. P.; Isnard, O. In Situ
Neutron Diffraction Study of Deuterium Gas Absorption by AB5+y
Alloys Used as Negative Electrode Materials for Ni-MH Batteries.
Phys. B: Condensed Matter 2004, 350 (1−3, Suppl. 1), e427−e430.
(11) Vivet, S.; Latroche, M.; Chabre, Y.; Joubert, J.-M.; Knosp, B.;
Percheron-Guegan, A. Influence of Composition on Phase Occurrence
during Charge Process of AB5+x Ni-MH Negative Electrode Materials.
Phys. B: Condensed Matter 2005, 362 (1−4), 199−207.
(12) Kohno, T.; Yoshida, H.; Kawashima, F.; Inaba, T.; Sakai, I.;
Yamamoto, M.; Kanda, M. Hydrogen Storage Properties of New
Ternary System Alloys: La2MgNi9, La5Mg2Ni23, La3MgNi14. J. Alloys
Compd. 2000, 311 (2), L5−L7.
(13) Denys, R. V.; Yartys, V. A. Effect of Magnesium on the Crystal
Structure and Thermodynamics of the La3−xMgxNi9 Hydrides. J. Alloys
Compd. 2011, 509, S540−S548.
(14) Crivello, J.-C.; Zhang, J.; Latroche, M. Structural Stability of ABy
Phases in the (La,Mg)−Ni System Obtained by Density Functional
Theory Calculations. J. Phys. Chem. C 2011, 115, 25470−25478.
(15) Petit-Ferey, A.; Cuevas, F.; Latroche, M.; Knosp, B.; Bernard, P.
Elaboration and Characterization of Magnesium-Substituted La5Ni19
Hydride Forming Alloys as Active Materials for Negative Electrode in
Ni-MH Battery. Electrochim. Acta 2009, 54, 1710−1714.
(16) Hu, W. K.; Denys, R. V.; Nwakwuo, C. C.; Holm, T.; Maehlen,
J. P.; Solberg, J. K.; Yartys, V. A. Annealing Effect on Phase
Composition and Electrochemical Properties of the Co-Free
La2MgNi9 Anode for Ni-Metal Hydride Batteries. Electrochim. Acta
2013, 96, 27−33.
(17) Nwakwuo, C. C.; Holm, T.; Denys, R. V.; Hu, W. K.; Maehlen,
J. P.; Solberg, J. K.; Yartys, V. A. Effect of Magnesium Content and
Quenching Rate on the Phase Structure and Composition of Rapidly
Solidified La2MgNi9 Metal Hydride Battery Electrode Alloy. J. Alloys
Compd. 2013, 555, 201−208.
(18) Rodríguez-Carvajal, J. FULLPROF: A Program for Rietveld
Refinement and Pattern Matching Analysis. Phys. B (Amsterdam,
Neth.) 1993, 192, 55−69.
(19) Denys, R. V.; Yartys, V. A.; Webb, C. J. Hydrogen in
La2MgNi9D13: The Role of Magnesium. Inorg. Chem. 2012, 51 (7),
4231−4238.
(20) Dorogova, M.; Hirata, T.; Filipek, S. M. Hydrogen-Induced
Volume Changes in ZrCr2 and Pseudo-Binary Compounds of ZrCr2,
ZrMn2 and ZrV2. Phys. Status Solidi A 2003, 198, 38−42.
(21) Yartys, V. A.; Burnasheva, V. V.; Semenenko, K. N. Structural
Chemistry of Hydrides of Intermetallic Compounds. Russ. Chem. Rev.
1983, 52, 529−562.
the GITT technique. Indeed, good kinetics are obtained at all
rates for the electrode though the maximum rate was limited to
only C/5 in this experiment.
CONCLUSIONS
The electrochemical behavior of a composite electrode made of
La2MgNi9 has been thoroughly investigated using deuterated
samples by in situ neutron powder diffraction at different
charge−discharge rates. From the data analysis, combining
diffraction and PCT measurements, the mechanisms and the
kinetics of the electrode have been determined. A fairly good
agreement is observed between the solid−gas and electrochemical capacities. The electrode material works by following
a hydride to intermetallic transformation through a β solid
solution domain and a β-to-α transformation with low capacity
attributed to the α solid solution domain. However, during
charge, heavy line broadening is observed, which might be
related to the formation of an intermediate γ phase as
previously observed in the LaNi5 systems. The electrochemical
reaction easily follows the current variations at all rates,
indicating no kinetic limitation of hydrogen exchange in the
materials between C/10 and C/5.
However, for the studied bulky electrode, part of the
reversible capacity is lost because of (a) formation of
electrochemically stable α H solid solution which contains up
to 1 H atom/f.u. and (b) incompleteness of the conversion of
the metal hydride anode electrode alloy into the β hydride
phase during the charging process. Thus, care should be taken
to achieve efficient performance during scaling up of the metal
hydride electrodes.
■
■
ASSOCIATED CONTENT
S Supporting Information
*
Crystallographic data taken from ref 13 for La2MgNi9 (R3̅m;
N°166) and from ref 19 for La2MgNi9D13 (R3̅m; N°166). This
material is available free of charge via the Internet at http://
pubs.acs.org.
■
AUTHOR INFORMATION
Corresponding Author
*E-mail: latroche@icmpe.cnrs.fr. Tel.: +33 1 49 78 12 10.
Notes
The authors declare no competing financial interest.
ACKNOWLEDGMENTS
This work was financially supported by the project NOVEL
MAGnesium based nanomaterials for advanced rechargeable
batteries (NOVELMAG) in the frame of the ERA.Net RUS
FP7 Programme 225.
■
■
REFERENCES
(1) Latroche, M. Structural and Thermodynamic Properties of
Metallic Hydrides Used for Energy Storage. J. Phys. Chem. Solids 2004,
65, 517−522.
(2) Notten, P. H. L.; Latroche, M. Secondary Batteries: Nickel
Batteries Metal Hydride Alloys. In Encyclopedia of Electrochemical
Power Sources; Garche, J., Ed.; ; Elsevier: Amsterdam, 2009; pp 502−
521.
(3) Liu, Y.; Pan, H.; Gao, M.; Wang, Q. Advanced Hydrogen Storage
Alloys for Ni/MH Rechargeable Batteries. J. Mater. Chem. 2011, 21,
4743−4755.
12169
dx.doi.org/10.1021/jp503226r | J. Phys. Chem. C 2014, 118, 12162−12169