Lithium Manganese Nickel Oxides Li(MnNi1,)2_O2
I. Synthesis and Characterization of Thin Films and Bulk Phases
B. J. Neudecker,* R. A. Zuhr, B. S. Kwak,* and J. B. Bates*
Oak Ridge National Laboratory, Solid State Division, Oak Ridge, Tennessee 3 7831-6030, USA
J. D. Robertson
Department of Chemistry, University of Kentucky, Lexington, Kentucky 40502, USA
ABSTRACT
The series Li(Mn Ni1_ )2_O2 for x s 1.33 and 0.38
0.50 shows a very close relationship to its parent series
1.26 are a linear function of the concentration ratio
Li/(Mn + Ni) which in turn is proportional to the averaged valence state of the transition metals. Li(Mn9Ni1_9)202 is
able to reversibly coprecipitate/reinsert Li20 and release/absorb 02. This self-regulation mechanism seems to always
adjust the number of cations to an undisturbed oxygen sublattice according to the rule "cations/anions = 1," which holds
true at least for temperatures up to 800°C and oxygen partial pressures above i0 atm. Samples prepared in air and under
02 did not show nucleation of Li20, not even for x> 1.0. The series Li(Mn6Ni1_6)202 where 0.38 y 0.50 crystallizes
in a rhombohedral unit cell (space group R3m) for x < 1.15 and transforms into a single monoclinic phase (space group
C2/c) for x> 1.25. The similarity between LiNi2_02 and Li,(Mn0Ni1_9)202 strongly suggests a rhombohedral — cubic
transition at x 0.6 for the latter series. Derived from the linear dependence of the X-ray density on the stoichiometric
parameter x, an equation was found with which the lithium concentration of Li(Mn9Ni,_0)202 thin film phases over the
entire range 0 x 1.33 can be determined accurately without extensive ion-beam analysis. XPS measurements on a
film with the bulk stoichiometry Li110Mn039Ni05102 gave evidence for Mn4 and Mn3, but no indication was found for
nickel valence states other than Ni2. In order to meet the above-given stoichiometry, the averaged nickel valence state
had to increase with film depth.
y
LiNi202. The refineà lattice parameters for at least 0.93 x
lnfroduction
In recent years, LiNiO2 has been used as a cathode mate-
rial in high-voltage lithium-ion and lithium batteries.1-4
Electrochemical experiments, however, showed that a slight
deviation from the ideal starting stoichiometry LiNiO2 resulted in a poor cell performance caused by Ni2 defects on
Li sites.5 Moreover, upon extensive lithium deintercalation
wherein x in LiNi02 approaches zero and the nickel
valence state comes close to its maximum of +4, nickel ions
migrate from their nickel layer sites into the vacancies of
the lithium layers. Such nickel migration creates severe diffusional limitations to lithium reinsertion thereby reducing
cell performance.6
These drawbacks stimulated our research on the partially substituted manganese derivatives of LiNiO2, namely,
Li(Mn.4Ni1_9)202. In particular, it was felt that the introduction of Mn" "buffer ions," which easily form under typical solid-state preparation conditions such as annealing at
700°C in air; would lift the electrochemically detrimental
effect Ni2-defects impose on chemically prepared LiNiO2
cathodes. That is, Mn" ions could help balance the two oxy-
gen anions without the need of Ni2 occupying Li sites.
Although LiNi2_02 for x < 1 can be chemically prepared,7
attempts to synthesize single-phase compounds with x/(2
x)> 1 were unsuccessful,7 because the lithium excess always
caused formation of Li2CO3 or Li20 as a second phase.
Bronger et al.8 and Migeon et al.,9 however, reported the
preparation of single-phase Li0 65Ni0350 (= Li1 30Ni0 7002) in
dry 02 at 450°C and Li2NiO3 (= Li1 33Ni06702) under 150 bar
02 670°C, respectively. Therefore we were encouraged to
attempt the synthesis of Lr(MflyNiiy)202 films with x
1.33 and to compare this series with the unsubstituted lithi-
um-nickel oxides. Furthermore, should single-phase
Li(Mn.9Ni1_0)2_02 be obtained, we were interested in deter-
mining the extent of oxidation of manganese and nickel in
subsequent electrochemical experiments upon lithium deintercalation. These compounds could retain some Li in the
lithium layers, even when all manganese ions and all nickel
ions have reached their maximal valence states of Mn" and
Ni", respectively. This idea is of particular interest with
respect to possibly suppressing the detrimental Ni-migration into the lithium layers of highly charged cathodes.
*
Electrochemical Society Active Member.
4148
The purpose of this paper is to describe the synthesis
and characterization of Li(Mn0Ni10)202 in thin-film and
bulk forms. Our electrochemical studies of the thin films
are reported in a companion.10
Experimental
All of the lithium manganese nickel oxide thin films were
deposited by planar rf magnetron sputtering from single
disk-shaped targets, typically 50 mm in diam and 5 mm
thick. These targets and the powder sample for the thermogravimetric analysis (TGA) served as samples for studies on
bulk phases.
Preparation and characterization of bulk phases.—
Stoichiometric amounts of LiOHH2O (Alfa Aesar, mm
56.5 wt % LiOH), Mn02 (Alf a Aesar; Puratronic, 99.999%
metals basis), and NiO (Alfa Aesar, Puratronic, 99.998%
metals basis), leading to the nominal cation stoichiometries Li124Mn037Ni039 and Li119Mn033Ni040, were thoroughly
blended and ground in an A1203 mortar. The loose powders
were placed in A1203 crucibles and heated at 600°C in air
for 3 h. After regrinding, the powders were cold-pressed
into disks (57.2 mm in diam) at 1 kbar and were subsequently fired in air at 900°C for 12 h. During cooling, the
targets were held for 12 h at 500°C in order to minimize
any possible lithium concentration gradients. This tem-
perature was low enough to ensure that no additional
lithium evaporated from the targets, but it was still high
enough for rapid Li diffusion, thus equilibrating surface
and bulk. The pellets shrank by 14% in diameter during
the final firing at 900°C.
Powder X-ray diffraction measurements (XRD) of the
targets were performed on a PC-controlled Scintag ThetaTheta XDS 2000 diffractometer using Cu Kc6 radiation.
Data were acquired in the 2 0 range 10-90° at a continuous
scan rate of 0.1°/mm, and the peak positions were found
after background removal by manual spline curve fitting
and performing Ka2 (X = 1.544390 A) stripping. Scrapedoff target powder samples were mixed and ground with
mica powder1' which served as an internal standard.
Lattice parameters were determined by weighted least-
squares refinement with the program LATCON.12 Because
the targets consisted of two-phase mixtures belonging to
the quaternary system Li-Mn-Ni-0, we did not attempt to
apply Rietveld analysis.
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J. Electrochem. Soc., Vol. 145, No. 12, December 1998 The Electrochemical Society, Inc.
The Li/Mn/Ni ratios of the powders scraped from the
front and the back side of the targets were determined by
inductively coupled plasma spectroscopy (ICP). Six independent measurements were made on each sample and the
results averaged. Assuming full occupancy of the oxygen
sublattice (see Results and Discussion), the stoichiometries
of the sputtered target surfaces were Li118Mn039Ni043O2 and
Li119Mn031Ni35002 for targets with the nominal cation starting compositions Li124Mn037Ni039 and Li119Mn933Ni043,
respectively. The unsputtered target back sides, representing
the bulk stoichiometry, were Li123Mn037Ni04002 and
Li119Mn033Ni04802, respectively, and, with respect to the
cation ratios, were identical with the starting compositions
within experimental error. The overall lithium oxide loss,
which is commonly observed in lithium first-row transition
metal oxides when annealed above 700°C under atmospheres other than pure 02, appeared to be minimal. We
assume that this can be attributed to a relatively large mass
(—30 g) pressed into a target pellet with a relatively small
geometrical surface area of about 55 cm2. The difference in
stoichiometry between the target bulk and sputtered surface
is a consequence of preferential sputtering. As can be seen
from the results, the trend is not the same in both targets.
The sample for TGA was prepared from Li0HH20,
Mn02, and NiO by pressing a powder with the nominal
cation composition Li145Mn045Ni045 into a pellet. The pellet
was fired in air for 3 h intervals at 300, 600, and 750°C.
During cooling to room temperature, the pellet dwelled for
12 h at 500°C. The pellet was ground to a fine powder
which then was loaded into the Pt crucible of the balance.
Alternate runs in air and flowing N2 were performed. The
sample was heated at 10°C/mm to 800°C and held at this
temperature for 1 h before cooling to room temperature.
At the conclusion of the TGA experiment, the powder was
analyzed by ICP and XRD in the same way as the targets.
Preparation and characterization of thin films—Thin
films were simultaneously deposited onto graphite disks,
single-crystal silicon wafers, quartz glass slides, and gold
foil substrates by rf magnetron sputtering of the lithium
manganese nickel oxide targets in 20 mTorr of a 50:1 Ar/02
gas mixture. The substrates and the thickness monitor
were fixed at a distance of 5 cm from the target, and the
net applied rf power of 50 W yielded a deposition rate of
typically 20 A/mm. The films were deposited through an
aluminum mask which defined a 12.5 X 8 mm area and
ranged in thickness from 0.3 to 3.0 p.m. The masses of the
600
4149
as-deposited films were estimated from the deposition
time and the steady-state deposition rate measured with
the thickness monitor before and after film preparation.
Subsequent surface profile measurements (Dektak 3030)
of the as-deposited films on smooth substrates (quartz
glass slides or single-crystal silicon wafers) yielded geo-
metrical densities of approximately 4.2 g/cm2. After deposition, the films were annealed for 3 h in different atmospheres (flowing 02, N2, or air) between 700 and 800°C in
order to investigate the influence of different oxygen par.-
tial pressures on the film stoichiometry. Typical heating
and cooling rates were 3°C/mm. The geometrical density
increased during the crystallization process, but the films
maintained their black color regardless of the annealing
temperature and atmosphere.
Thin film analysis and characterization—Powder X-ray
diffraction (XRD) measurements of the 2-3 p.m thick films
were performed as described above for the targets. In these
cases, the substrate materials (usually Au foil) served as
internal standards. The as-deposited films were X-ray
amorphous. Lattice parameters of the crystallized films
were determined by least-squares refinement using the program LATCON. Rietveld analysis was not applied because
all of the films exhibited preferred orientation as demonstrated in rocking curve scans13 for the (101) and (104)
reflection of a Li110Mn045Ni04502 and a Li1 14Mn043Ni04302
films, respectively (Fig. 1).
Rutherford backscattering spectrometry (RBS) was used
to determine the Mn/Ni/0 ratios of the films. The measured energy spectra (Fig. 2) represent a convolution of
both mass and depth information that was separated by an
appropriate fitting program. Due to its low atomic number
and mass, lithium could not be detected directly, but its
presence was inferred from the stopping powers necessary
to fit the Mn-Ni-0 data. The stoichiometry of as-deposited amorphous films was determined on graphite substrate
disks whereas the annealed crystalline films were measured on Si wafers. We noticed a light Si0 tarnishing layer
on the nondeposited area of the wafer when annealed in
air at 800°C. Although this layer was much thinner than
the films to be investigated, the presence of lithium ions at
the film/silicon interface certainly enhanced its formation.
The tarnishing layer was determined by X-ray diffraction
to consist of Li25i02 and an unidentified compound. Because of the presence of the Si0 and Li2SiO3 layers, the
measured (Mn + Ni)/0 ratio is smaller than the actual
900
800
500
Fig. 1. XRD omega step scans
("rocking curves") (a) for the
=
(101) peak fixed at
36.68° of a 2.6 p.m thick thorn-
a. 700
U,
C,,
a.
•'a 400
bohedral Li7 1Mn045Ni0.O2 [aft-
CI,
>1
Si
>1
I-
U,
er RBS, PIXE, PIGE, and XRD]
600
film on Au foil annealed at
750°C in air for 3 h and (b) for
the (104) peakfixe4at28,,,,,
U,
300
C 500
0
Si
C
44.500 of a 9,000 A thick thombohedral Li1 14Mn043Ni04302 [aft-
er MS. PIXE, PIGE, and XRD]
400
film on Pt foil annealed at 750°C
under 02 for 3 h. The appearance of a peak in both omega
scans affirms the preferred ori-
200
300
100
10 15 20 25 30 35
w (Degrees)
entation of these polycrystalline
films. All other films gave similar
scans due to film texiure.
200
15 20 25 30 35 40
w (Degrees)
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4150
J. Electrochem. Soc., Vol. 145, No. 12, December 1998 The Electrochemical Society, Inc.
2500
0:i
2000
C
1
Fig. 2. a) RBS spectrum of a
5,600 A thick crystalline
Li111Mn039Ni0O2 [after RBS,
PIXE, PIGE, and XRD] film on Si
substrate. The RBS curve-fitting
1500
U)
results in Li111Mn041Ni0O2
C
with a thickness of 5,540 A. The
0.06 oxygen excess is attributed
0
1000
C)
to oxygen taken up by the Si
substrate during the postanneal
treatment. The insets show the
upper edge of the overlappin
energy scatter ranges of Mn an
Ni and the Si substrate and oxygen, respectively.
500
C)
>-
0
800
400
1200
1600
2000
Energy (keV)
ratio of the Li-Mn-Ni-U film. However, we show below
that the composition of all Li-Mn-Ni-U films can be represented by L(MflyNi_y)2_rU2 and that the deviation of
the true (Mn + Ni)/U ratio from the measured ratio falls
within the standard error of RBS measurements which is
estimated to be less than 10 atom %.
Particle-induced X-ray emission (PIXE) analysis14 was
used to determine the Mn/Ni ratio in the films. The lithi-
um content could not be measured with this technique
because the energy of the Li Ku X-ray emission line at
0.052 keV is well below the detector threshold value. The
areal masses of Mn and Ni in the films were analyzed by
comparing the Mn and Ni X-ray yields to those obtained
from thin-film standards (MicroMatter, Inc.). An example
of a PIXE spectrum recorded from a crystalline sample is
shown in Fig. 3. PIXE measurements confirmed that the
Mn/Ni ratios were 0.76 0.05 and 1.00 0.07 for the films
deposited from targets with the ratios Mn/Ni = 0.62
and
0.91, respectively, at the sputter face. These ratios deter-
mined for the amorphous films did not change upon
annealing, at least up to 800°C, independent of atmos-
phere. This is most likely due to the negligible volatility of
the manganese and nickel oxides.
The lithium and the manganese contents of the films
were determined by particle-induced gamma-ray emission
(PIGE) analysis.15'16 The nuclear reactions 7Li(p,p'y)7Li
= 477 keV and 55Mn(p,n-y)55Fe =
E 932 keV at a proton
bombarding energy of 2.0 MeV were used. This low bombarding energy precluded the determination of the nickel
content of the film, but had to be chosen in order to reduce
the high fly-ray count rate from the silicon substrate at 4
MeV. An example of a PIGE spectrum which was taken
E
from a postannealed crystalline Li126Mn037Ni037O2 film on
Si is shown in Fig. 4. Single measurements of the Li/Mn
atom ratio in a film yielded results with uncertainties on
160
_ 140
120
c.l
U)
4-
8O
C
00
60
a)
.E
150
100
>
40
U)
C
50
a)
4-
20
C
0
0
0
2
4
6
8
10
0
Energy (keV)
Fig. 3. PIXE spectrum of a 9,000 A thick crystalline
Li1 26Mn037Ni03702 [after RBS, PIXE, PIGE, and XRD] film on Si sub.
strate annealed at 750°C under 02 for 3 h. The sputter surface of
the target had the composition Li118Mn039Ni0,4302 [after ICP and
evaluating the oxygen stoichiometry].
1000
2000
Energy (keV)
Fig. 4. PIGE spectrum of a 9,000 A thick crystalline
Li1 26Mn037Ni03702 [after RBS, PIXE, PIGE, and XRDI film on Si sub-
strate annealed at 750°C under °2 for 3 h. The relevant Li and Mn
signals are marked with the appropriate nuclear reaction.
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J. Electrochem. Soc., Vol. 145, No. 12, December 1998 The Electrochemical Society, Inc.
Fig. 5. SEM photograph of a
A thick as-deposited amorphous Li,23Mn0Ni0,,6O2 film on Si substrate. The average grain
size is about 1,000 A.
the order of 10%. As in the PIXE measurements, the 5%
uncertainty in the certified concentration of each thinfilm standard from Micromatter, Incorporated, was the
major source of experimental error.
X-ray photoelectron spectra of the films were obtained
with a Physical Electronics PHI 3057 Subsystem XPS
equipped with a spherical capacitor energy analyzer and
an Omni Focus III small area lens. The Mg anode was
operated at 26.7 mA and 15 kV (400 W) and emitted Mg Ka
radiation of 1253.6 eV. High-resolution spectra at 11.75 eV
pass energy were recorded in the C (is), 0 (is), Ni (2p),
Mn (2p), Mn (3p), and Li (is) regions. The specimens were
analyzed at an electron takeoff angle of 45°, measured with
respect to the surface plane. Measurements were made on
a 2.6 p.m thick film prepared at 700°C in air which had the
bulk stoichiometry of Li1 10Mn039Ni051O2. The XPS peaks
were deconvoluted into asymmetric Gaussian envelopes
with the standard integral background subtraction
method of Shirley.'1 The spectra were corrected for the dif-
ference between the observed position of the C (is) peak
4151
Fig. 6. SEM photograph of the film shown in Fig. 5, but now crystallized after an anneal at 700°C under 02 for 3 h. The stoichiomefry
changed to Li126Mn037Ni03702. The average grain diameter
increased to approximately 0.3 an.
(285.57 eV) and the reference position of adventitious
hydrocarbons (285 eV).
Scanning electron micrographs were obtained with a
JEOL JSM-840 SEM equipped with a Tracor Northern
5525 energy dispersive X-ray (EDX) analyzer. The as-deposited amorphous films had an average grain size of
approximately 0.1 p.m (Fig. 5). After the postanneal step at
700°C in air, the crystalline films exhibited an enlarged
grain size of about 0.3 p.m (Fig. 6). No elements other than
Mn, Ni, 0, and elements of the underlying substrate were
detected in the EDX measurements.
Results and Discussion
Structure studies.—Bulk phases from the targets—At the
sputter surfaces and at the back side surface, the gray-black
targets were two-phase mixtures of a rhombohedral phase
and a monoclinic phase which could be indexed to the space
groups R3m (LiNiO2-relate&8) and C2/c (Li2NiO3 =
3/2 Li1 33Ni0 6702-related9), respectively (Fig. 7 and Tables I-
4000
U)
Fig. 7. XRD powder pattern of
the target sputter face with the
overall
stoichiomefry
Li,,9Mn031Ni0O2. Arrows indicate the peaks arising from the
0.
2400
mica standard." The pattern
comprises a rhombohedral
—
1600
phase (!jNiO2-related, space
group R3m) and a monoclinic
800
phase (Li2NiO3-related, space
group C2/c) which both belong
to the series
MnNi,..,J2..,,O2.
The inset shows the decisive 20
range where only peaks origi-
nating from the monoclinic
phase are expected to appear.
0
10 20 30 40 50 60 70 80 90
2 9 (Degrees)
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4152
J. Electrochem. Soc., Vol. 145, No. 12, December 1998 The Electrochemical Society, Inc.
Table I. Experimental and refined X-ray data of the rhombohedral
phase (space group R3m) sputtered from the twa-phase target
surface with the overall stoichionefry Li1 19Mn031Ni0O2 displayed
in Fig. 7 (Cu Ka1 = 1.540562 A). Experimental 20 values after
correction on the internal standard mica. Estimated stoichiomefry
(ci. Fig. 14 and 16) of the rhombohedral phase:
Table II. Experimental and refined X-ray data of the monoclinic
phase (space group C2/c) sputtered from the two-phase target
surface with the overall stoichionetry Li1,19Mn031Ni0O2 displayed
in Fig. 7 (Cu Ka1 = 1.540562 A). Experimental 20 values after
correction on the internal standard mica. Estimated stoichiometry
(cf. Fig. 14 and 16) of the monoclinic phase:
3/2Li125(Mn9Ni1..,)07502
Li1 15(MnyNii Y)0.8502.
20
(Exp.)
hk
1
18.697
36.572
37.895
38.195
44.355
48.535
58.547
64.348
68.048
81.812
0
0
0
0
1
0
1
0
3
0
1
8
1
1
2
3
4
1
0
0
1
0
1
0
1
6
2
4
5
7
20
20
(Calc.)
Relative
intensity
d (A)
(Exp.)
(Caic.)
(Exp.)
18.702
36.575
37.927
38.236
44.351
48.527
58.570
64.313
68.036
81.843
100
16
5
6
4.7421
2.4550
2.3723
2.3543
2.0406
1.8742
1.5753
1.4466
1.3766
1.1763
4.7407
2.4548
2.3703
2.3519
2.0408
1.8745
1.5747
1.4473
1.3768
1.1760
20.797
20.958
21.668
24.258
36.852
44.596
48.775
58.708
58.968
60.454
64.488
68.628
69.272
77.492
80.953
81.812
82.593
82.815
84.394
45
5
6
12
3
2
ci (A)
Hexagonal unit cell parameters: a = 2.8777(7) A), b = 14.222(4)
A, volume = 102.0(0) A3, and Z = 3.
III). The close structural relationship between LiNiO3 and
Li3NiO3 discussed below, where in both cases the number of
cations equals the number of oxygen anions, strongly suggests that the overall target stoichiometries also adopt this
ion ratio. Therefore, we can reasonably justify the stoichiometric parameters of oxygen given in the experimental section. The targets showed approximately 70% of their theoretical density based on the amount and density of the two
phases present in each target (see below).
In order to understand the structure of lithium manganese nickel oxide, we first demonstrate the structural
resemblance between rhombohedral LiNiO3 (R3m) and
monoclinic Li3NiO3 (C2/c). Considering (i) that Li3MnO3
crystallixes in the monoclinic space group C2/c with unit
cell parameters close to those of Li3NiO3 and (ii) that the
ionic radii33 of low-spin Ni" (r = 0.56 A) and low-spin Mn"
(r = 0.58 A) as well as Ni" (r = 0.48 A) and It4 (r = 0.53 A)
are very similar to each other within an octahedral environment, we propose that manganese ions occupy nickel sites in
manganese-substituted lithium nickel oxides. This statement is
supported by a neutron diffraction study on rhombohedral
LiMn83Ni08O3 showing that Mn substituted for Ni only on the
predominantly Ni-filled layers of the R3m structure.21 Thus,
the structure of rhombohedral Li(Mn9Ni3 y)3x°3 can be
derived from the respective rhombohedral parent compound of the series LiNi3O3. Judging from the cation radii
given above, we believe that the structure of monoclinic
Li(Mn9Ni3j302 can be obtained by simply substituting
manganese for nickel in monoclinic LcNi2_X02.
h
k
0
2
1
—1
1
2
—1
1
1
1
2
—1
3
—3
3
—1
—2
—4
—3
1
2
0
1
0
3
3
0
1
1
1
3
1
2
2
5
7
6
4
Li1 8€(MnNi1 .Y)1.1203
20
(Calc.)
l
0
0
1
1
0
3
3
4
6
1
4
—3
7
7
1
0
1
2
7
Relative
intensity
20.808
20.980
21.646
24.250
36.864
44.599
48.784
58.680
58.954
60.405
64.487
68.671
69.240
77.468
80.947
81.809
82.586
82.805
. 84.385
6
4
3
4
42
100
10
21
7
7
36
8
6
7
5
5
5
6
5
(Exp.)
d (A)
(Calc.)
4.2677
4.2353
4.0981
3.6661
2.4370
2.0301
1.8655
1.5714
1.5650
1.5301
1.4438
1.3664
1.3553
1.2308
1.1866
1.1763
1.1672
1.1646
1.1468
4.2653
4.2309
4.1021
3.6673
2.4362
2.0300
1.8652
1.5720
1.5654
1.5313
1.4438
1.3657
1.3558
1.2311
1.1867
1.1764
1.1673
1.1647
1.1469
ci (A)
Unit cell,parameters: e = 4.9454(4) A), 6 = 8.531(1) A, c =
9.6648(9) A, I = 99.96(l)°. Volume = 401.6(1) A3, Z = 8
(X Li3 33(Mn8Ni3)3 3303).
The Li3NiO3 structure was determined following the monoclinic 3-Li3SnO3 structure (C2/c) 9,22,23 wherein the oxygen
anions form a cubic close-packed array with alternate layers occupied either by lithium ions only or by lithium and
nickel ions in the ratio 1:2. Based on the similarities of nickel and manganese, the monoclinic unit cell of a hypothetical
Li3Mn8Ni1.9O3 phase was constructed as shown in Fig. 8. If
this monoclinic cell is extended by 150% in c direction, the
hexagonal unit cell of LiNiO2 can be seen as illustrated in
Fig. 9 and 10 for the case wherein manganese has been par-
tially substituted for nickel. For better orientation, we
marked an arbitrary lithium ion surrounded by six transition metal neighbors on one of the Li-Mn-Ni layers. The
pseudo-hexagonal unit cell in Fig. 9 is almost identical with
the hexagonal LiNiO3 unit cell sketched in Ref. 24, except
that in the case of Li3Mn9Ni19O3 the nickel layer is partially occupied by Li and manganese ions, and the monoclinic
c axis is inclined by 13-90° with respect to the hexagonal unit
cell (Fig. 10). As far as the mixed cation layer is concerned,
the lithium ion is assumed to be slightly off center within the
hexagonal nickel-manganese environment (Fig. 11), as was
Table Iti. Monoclinic (space group C/2c) and rhombohedral (space group Rm) target phases.
Overall stoichiometry
at the two-phase target surface
Estimated phase stoichiometry°
e=
Li1 39Mn339Ni943O3
(Sputter surface)
LiLII(M;Nill)l?303
= 2/3Li338(Mn Ni3)L3303
Li3 33(Mn9Ni3..333O3
Li319Mn933Ni939O3
(Sputter surface)
L1L3I(MnYNil)W7303
= 2/3Li199MnINi3U)L3303
Li1 33(Mn3Ni33)8 8303
Li333Mn833Ni348O3
(Back side)
Li1 33(Mn9Ni39)3 7303
= 2/3Li333(Mn8Ni3)33303
Li1 33(MnNi3)395O3
Li3 39Mn833Ni388O3
(Back side)
Unit cell parameters a, b, c, in (A)
Li1 33(Mn9Ni3)3 7503
= 2/3Li381(Mn 1i39)L3303
Li3 33(Mn3Ni33509503
4.946(1)
b = 8.539(2)
c = 9.658(4)
e = 2.8717(3)
a = 4.9454(4)
b = 8.531(1)
c=
9.6648(9)
13 = 100.02(3)°
c = 14.2244(3)
= 99.96(1)°
e = 2.8777(7)
e = 4.9353(7)
c=
e = 2.8763(3)
a = 4.9389(8)
6 = 8.544(2)
c = 9.658(2)
a = 2.881(1)
c = 14.239(2)
6 = 8.527(2)
c = 9.677(3)
14.222(4)
= 99.97i2)°
I = 99.89(2)°
c = 14.228(6)
° For the Li/(Mn + Ni) ratio of each phase, cf. Fig. 16. Only the overall Mn/Ni ratio of a two-phase region is known, the Mn/Ni ratios
of each phase may be different.
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4153
J. Electrochem. Soc., Vol. 145, No. 12, December 1998 The Electrochemical Society, Inc.
Ch = 3/2 Cm sin(18O°-3)
• Oxygen
• Lithium
3/2Cm Sifl
o Nickel,
Manganese
Fig. 8. Monoclinic unit cell of Li2Mn2Ni,_ 03 (space group C2/c)
derived from the LI2NiO3 structure.' The iIexagonal nickel-manganese environment around an arbitrary lithium ion is highlighted
for better comparison with Fig. 9.
observed for Li2NiO3, in order to minimize the coulombic
repulsion between the Li ions and the Ni4 and
ions.
When comparing the monoclinic unit cell of Li2Mn.dNiI903,
which
contains eight formula units [=
8
(three formula units), the following transformations Eq. 1—6
are determined from Fig. 9-11
am =
b,,,
[1]
= 3 a,
Fig. 10. Geometrical relationship between the monoclinic (space
group C2/c) and the hexagonal (space group Rim) unit cell within
the series Li,(Mn,Ni,..,)3_002. The indexes m and h denote mono-
clink and hexagonal, respectively. The relation of the a parameters is illustrated in Fig. 11.
[2]
a,, = (am/3+ b,,,J3)/2
[3]
C = 3/2 Cm Sifl 13
[4]
VOlm = 4V,,
5]
13 = ir — arctan
am=J3ah
X 3/2
Li133(MnNi1_9)0.,7O2], with the hexagonal LiNiO2 unit cell
(ch/ah
[6]
where the indexes m and h stand for the monoclinic and
hexagonal unit cells, respectively; a, b, c, and 13 are cell
parameters, and V denotes the volume. The above trans-
but hold true for the entire series LrN2O2 and
Li(MnNi1_)2O2 wherever a monoclinic unit cell (C2/c)
is adopted for x 1.33 and the transition metal ions are
octahedrally coordinated by oxygen ions. It appears that
the monoclinic distortion of the rhombohedral structure
occurs as soon as the lithium concentration on the pre-
dominantly filled transition metal layers reaches a
threshold value at a composition of Li1 22Ni0 7802 8 and, in
formations are not limited to LiNiO2 and Li2Mn21Ni18O3,
• Lithium
• Oxygen
Nickel, Manganese
-c
C)
II
E
ah=l/3bm
Fig. 9. The monoclink unit cell of Li2MnNi,_703 from Fig. 8 has
Fig. 11. Geometrical relationship between the a parameters and
b parameters of the monoclink (space group C2/c) and the hexag-
been extended by 150% in the c direction. Far better orientation, the
same six-membered ring as highlighted in Fig. 8 has been marked.
onal (space group Rim) unit cell within the series
Only that part of the extended monoclink unit cell is shown which
corresponds to the hexagonal unit cell of UMnNi1. 03. This monoclinic (pseudo-hexagonal) cell becomes identical witl the hexagonal
Li,,MnNi,_)2_0O2. The indexes m and h signify monoclinic and
hexagonal, respectively. The highlighted nickel-manganese six-
cell when = 900.
and 9.
membered ring around lithium is the same one as marked in Fig. 8
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4154
J. Electrochem. Soc., Vol. 145, No. 12, December 1998 The Electrochemical Society, Inc.
the case of the manganese-substituted compounds, at
Li1,15(MnNi2_)0,8502, as will be shown below. We assume
that substituting the larger Li ions (r = 0.76 A 20) for the
smaller M2 and M4' ions on the octahedral sites belonging to the same layer induces a sufficient stress to cause
a hexagonal-to monoclinic distortion.
Comparison of thin-film and bulk phases.—All of the
films prepared in air or 02 atmosphere gave single-phase
XRD patterns indicating phases of the series
Li(Mn,Ni1 _y)2_r02. As shown in Fig. 12, the film annealed at
800°C under 02 having the stoichiometry Li126Mn037Ni03702
(= 2/3 Li199Mn957Ni056O3) crystallized in a monodlinic lattice.
In comparison to an XRD pattern of a sample showing R3m
symmetry, features that are typical of the monoclinic phase
based on space group C2/c comprise the peak conglomerate
at about 20° < 20 < 23°, the (015)h peak splitting into (133)m
and (222)m, and the (113)h peak splitting into (331)m and
(117)m. Only the (002)m peak of the conglomerate is clearly
visible whereas the peaks (110)m, (hhl)m, and (021)m are poor-
ly resolved. Thus, the last three peaks were not included in
the lattice refinement below. Also, it was confirmed in a
grazing angle scan, which generally identifies substrate
peaks by drastically diminishing their intensity, that the
(220) Au-substrate peak at 20 64.6° was very small and did
not affect the intensity ratio of (314)m/(331)m.
The lattice refinement resulted in am = 4.9400(7) A, bm =
8.558(6) A, Cm = 9.671(2) A, and 3 = 100.04(5)° which can
be converted into a pseudo-hexagonal cell having a =
2.852 A and c5 = 14.284 A according to Eq. 1—6. This very
small ah parameter is consistent with the high average
valence state of manganese and nickel of +3.7, because the
a5 parameter is a measure of the average M-M (M = Li,
Mn, Ni) intralayer distance on the predominantly filled
transition metal layers illustrated in Fig. 11. The average
tuted for the smaller Ni3 ions (r = 0.56 A), Mn3 ions (r=
0.58 A), Ni4 ions (r = 0.48 A), and Mn' ions (r = 0.53 A).
We believe that the successful preparation of the monoclinic single phase under "normal" annealing conditions
was alleviated by employing the unique characteristics of
sputter deposition. In the as-deposited films, the Li, Mn,
Ni, and 0 atoms were not only homogeneously distributed
over the entire sample, but are also assumed to have already developed a short-range order similar to their local
environment in the crystalline state after annealing. Thus,
only slight atomic translations within the chemically
homogeneous and X-ray amorphous grains shown in Fig. 5
were necessary to form the crystalline monoclinic lattice.
In contrast, the typical preparation of bulk samples from
powder mixtures involves a much higher activation ener-.
gy. On a grain-size level, these powders are highly inhomogeneous where, e.g., -'-1 i.m grains of Li2CO3 are neighbored by —1 p.m grains of MnO2 and —'-1 p.m grains of NiO.
Interdiffusion of all three cations on a micron-scale and
the perturbation caused by significant amounts of evaporating CO2 (H20 or NO when using LiOH or LiNO5,
respectively) certainly increases the activation energy of
reaction and crystallite formation compared to that of
sputter-deposited films.
The XRD patterns of films annealed in N2 atmosphere
(Pos = i0- bar) revealed that each film was composed of a
Li(Mn9Ni1_9)2O2 phase and coprecipitated Li20 (Fig. 13),
which could react to form Li2CO3 and/or LiOH during the
XRD measurement in air. These subsequent reactions do not
present a problem for the interpretation of XRD results, but
they must be avoided when determining the oxygen concentration by RBS, because the formation of Li2CO3 or LiOH
implies the absorption of CO2 or H2O, respectively. To minimize this problem, we limited the air exposure time to less
pensates the adverse effect of the monovalent, larger Li'
ion which enters the transition metal layers when the lat-
than 30 s when loading the samples into the RBS equipment. Furthermore, a result of Li20 coprecipitation is that
not all the analyzed lithium and oxygen belonged to the
Li,(Mn9Ni89)2O2 phase. Thus, we make the following
assumptions for the calculation of the x parameter: Although the oxygen partial pressure of about iO bar was
LiNiO2, however, the relatively large c5 parameter (a meas-
as has been observed in lithium nickel oxides even when
Ref. 2) of monoclinic Li8 26Mn027Ni03702 is attributed to the
sublattice, however, appears to be unlikely after annealing
below 800°C. This is supported by Rietveld refinements of
distance shortens when the average valence states of manganese and nickel increases, i.e., when the average radii of
manganese and nickel shrink. This trend has already been
observed for LiNi2_O2 5,8,9,24,25,27 and obviously overcom-
tice becomes monoclinic. As compared to hexagonal
ure for the mean interlayer distance as seen in Fig. 9 and
expanded height of the Li-Mn-Ni layers where a significant concentration of larger Li' ions (r = 0.76 A) substi-
low enough to account for Li2O and O loss upon annealing,
annealed under 02, 5,7,25,29 the formation of an oxygen-defect
neutron and X-ray diffraction measurements5'7'25'26 on lithi-
um nickel oxides that had a higher cation disorder (lower
400
350
350
300
300
U)
0. 250
U
200
-0.250
150
150
0
100
50
0
10 20 30 40 50 60 70 80 90
29 (Degrees)
Fig. 12. XRD pattern of a 3 m thick Li26Mn037Ni037O2 film on
gold foil substrate (5) annealed at 800°C under °2 for 3 h. Miller
indexes are based on the monoclinic space group C2/c. The poorly resolved peaks (110), (111), and (021) were not included in the
lattice refinement. The peak at 20 40° was confirmed to stem
from the substrate by annealing an undeposited gold foil un4er the
same con4itions. Refined lçittice parameters: a — 4.9400(7) A, b =
8.558(6) A, c = 9.671 (2) A, and = 100.04(5)°.
20
30
40
50
60
70
20 (Degrees)
Fig. 3. XRD powder pattern of a 2.6 m thick rhombohedral
Li0 93Mn0 Ni061O2 film on gold foil substrate (5) annealed at
800°C under N2 for 3 h. The stoichiometry of Li0 93Mn0 46Ni0,6102
was calculated using Eq. 7—13. Miller indexes are given for
Li093Mn046Ni061O2, and the Li20 peaks are also marked. The
peak at 20 40° was confirmed to stem from the substrate by
annealing an undeposited gold foil under the same conditions.
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4155
J. Electrochem. Soc., Vol. 145, No. 12, December 1998 The Electrochemical Society, Inc.
LifM ratio) than our films and wherein the number of
cations still equaled the number of anions which in turn
formed a fully occupied oxygen sublattice. Moreover, we
believe that the Li20 and 02 loss appears to be a facile route
for lithium manganese nickel oxides to adjust their cation
ratios to an undisturbed oxygen sublattice. As a consequence, the stoichiometry of the actual Li-Mn-Ni-0 phase
in the films annealed under N2 was computed on the basis
of the overall film composition LiOMnbNiCOd (including
Li20) according to
[LiaMnbNicOd]annealed = mLi(Mn8Ni1_8)2.102
+ nLi2O [7]
2.8
0 104
E 103
0 102
0 101
a4C
100
where only a, b, c, and d could be analyzed by EBS,
£ it 000
i
PIXE, and PIGE. The values for the four variables m, n,
x, and y were obtained after solving the following system
of equations
b/c = y/(1 — y)
[8]
a=mx+2n
b + c = m(y + 1 — y)(2
—
(50 (j 0(50
14.3
[9]
x) = m(2
—
x)
d=2rn+n
[]
[11]
Because the as-deposited films had as many cations as oxygen anions, namely, Li1 28Mn036Ni03002 and Li1 28Mn834Ni04602,
and annealing in N2 produced an overall film stoichiometry
of LiaMnbNi,Od where a + b + c> d, we conclude that the
Li20 coprecipitation had to be accompanied by corresponding 02 evolution
[Li5(Mn8Ni1 —o)2—xo°2]as deposited = [LiaMnbNicOd]000ealed
+ n/2 021 [12]
or inserting Eq. 7 into Eq. 12
[Li,0(Mn8Ni1 —)2 X02]a5 deposited
Averaged Valence State of M
3
3.2
3.4
3.6
3.8
= inLijMn8Ni1 8)2_X02
+ nLi2O + n/2 021 [13]
As mentioned above, films annealed in air or 02 also fol-
low the notation Li(Mn8Ni18)2_02 where x is smaller
than the as-deposited stoichiometric lithium parameter x,
i.e., Eq. 13 must generally hold true, with the amendment
in this case that Li20 does not coprecipitate, but rather
evaporates. We did not find any indication of Li20 in the
XRD patterns, including that of the Li896Mn052Ni05202 film
which was prepared at 800°C in air and which had to lose
13 atom % Li20 based on its as-deposited stoichiometry. If
precipitated, this large amount of Li20 would have been
detectable in the 2 p.m thick film as seen in the XRD pat-
U
14.26
14.22
2.9
2.88
o 2.86
2.84
0.8
1
1.2
1.4
1.6
1.8
LI I M
Fig. 14. Unit cell parameters of the rhombohedral (space group
R3m) and the monoclinic (space group C2/c) Li(Mn,,Ni,_,J2_0O2
film phases in hexagonal setting as a function of the concenfration
ratio Li/M where M = Mn + Ni. The hIM ratio is proportional to
the averaged valence state of the pseudofransition metal M. For
comparison, monoclinic unit cells are converted into hexagonal
cells according to Eq. 3—5. The annealing conditions are given for
each data paint. Annealing times at the final temperature were
3 h for all films. The films were sputtered from the target with
the sputter-face stoichiomefry (0) Li,19Mn031Ni03002 or (A)
hi,,8Mn039Ni04302 and produced films with (0) y = 0.43 or (A)
y = 0.5, respectively. The rhombohedral phase Li,,,Mn042Ni04702
after the TGA run (7) and the phases of the target surfaces are
also included.
tern of Li893Mn846Ni06102 prepared under N2 (cf. Fig. 13).
Concluding from Fig. 13, Li20 developed a long-range
order with a particle size of at least 900 A according to the
Scherrer equation.'3 The true value is probably higher due
to strain-broadeningt3 of the measured peak because our
films were attached to a substrate during preparation. Apparently, the oxygen from the gas atmosphere during the
annealing process in air helps Li20 to evaporate rather
than to coprecipitate whereas too high an oxygen partial
pressure, e.g., pure oxygen, appears to arrest 02 evolution
from the films thereby giving LiX(Mn.dNi, 9)2_X02 film phas-
es with higher x values (Fig. 14). As a general trend in air
and N2 atmosphere, we find that the higher the annealing
temperature and the lower the oxygen partial pressure
during annealing, the lower is x in the actual lithium manganese nickel oxide phase. It is important to note that only
thin films can be compared to each other, as their sur-
face/mass ratios (—1 cm2/103 g = i0 cm2/g) were more
than 500 times greater than those of the targets (—55 cm2/
30 g = 1.83 cm2/g), and a high surface/mass ratio significantly enhances the percent-mass loss of a heated body
containing volatile species.
Plots of the unit cell parameters of the Li(IVIn,Ni1_8)2_02
phase in a hexagonal setting as a function of the concentration ratio Li/M = x/2 — x, where M = Mn + Ni, are shown
in Fig. 14. Because Li/M is proportional to the average
valence of M, we notice a virtually linear dependence of the
unit cell parameters on the average valence state of M, itgardless of the Mn/Ni ratio and the annealing conditions.
Moreover, we see from Fig. 14 that the unit cell parameters
of the target surfaces blend into the trend that was observed
for the film phases and fill the mandatory two-phase gap
between the rhombohedral R3m phase and the monoclinic
C2/c phase. As a result, the two-phase region is limited by
the phase boundaries near Li/M = 1.35 and 1.65, i.e., the
rhombohedral single phase R3m appears at
LiX(IVIn.dNilY)2X02 for x < 1.15 and the monoclinic single
phase C2/c for x> 1.25, at least for 0.38 y 0.5. The value
y = 0.38 stems from the sputter target surface with the overall stoichiometry Li,,91V1n031Ni05002 = Li1 19(Mn0 36Ni8 62)0.8,02.
With an increasing Li/M ratio, we notice from Fig. 15 an
increasing departure of the c/a ratio of the hexagonal
phases from the cubic rocksalt structure value of 2J6,
which is accompanied by an appreciable amount of transition metal ions in the lithium layers.2'5'7272° Obviously,
the manganese and nickel ions did not significantly mix
with the lithium ions from the lithium layers (cf. Fig. 8 and
9), neither in the hexagonal structure nor in the monoclinic structure. This may be understood in view of the high
lithium concentration in all of our films.
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4156
J. E/ectrochem. Soc., Vol. 145, No. 12, December 1998 The Electrochemical Society, Inc.
Averaged Valence State of M
2.8
5.04
3.2
3
3.4
3.8
Overall stoichiometry of
0
.
5.00
.5! 4.96
0
3.6
Table IV. Phase distribution within the bulk of the sputter targets
and X-ray density of the targets.
hexagonal
.
0
0 a
.
Emono °
o
Monoclinic phase
(space group C2/c)
.
0
o
hex
4.92
the two-phase target
mono-
cUnlc .
.,
0.8
I ,l cubic)
1.2
1
1.4
.1
1.6
,,l
1.8
Li I M
Fig. 15. Deviation of the hexagonal unit cells of Li,Mn2_,O2 (M
Mn7Ni.7 and 0.38 y 0.5) characterized by the unit cell parameter ratio c/a from the cubic rock salt structure where c/a would
become 2J. In the cubic rocksalt structure a significant number of
transition metal ions would reside in the lithium layers.2'5'7'27'28
The X-ray density of the Li(Mn9Ni1..5)2_O2 series
(Fig. 16), which was obtained from the unit cell volume of
known stoichiometry (Fig. 14), shows a linear dependence
on the parameter x. This implicitly supports the validity of
our oxygen analysis by BBS. If the measured (Mn + Ni)/O
ratio had been significantly lower than the ratio of the
actual lithium manganese nickel oxide film due to the
presence of the SiO and Li2SiO3 tarnishing layers on the
Si substrate (see above), we would have in effect substituted oxygen ions for heavier Mn and Ni ions in calculating the film stoichiometry of Li(Mn5Ni..8)2.,O7. Because
the determination of the unit cell volume was independent
of the BBS measurements, this should have resulted in a
noticeable decrease in the calculated X-ray density with
increasing temperature and increasing oxygen partial
pressure, which is not observed in Fig. 16. The X-ray den1.25 can
sity within the two-phase region 1,15 x
change due to Gibbs' phase rule: (i) LiM2O2 (M =
behaves at the very least as a pseudo-ternary
system Li-M-O so that there remains one degree of freedom after choosing pressure and temperature. Therefore,
the unit cell dimensions of all phases present can vary
with x. (ii) If Li(Mn9Ni15)O2 behaves as a quaternary system, two degrees of freedom will remain. So, the unit cell
parameters may again very with x. However, the limited
number of data points allowed only an estimation of the
Mn5Ni1.5)
Li119Mn833N1848O2
80 atom %
40 atom %
Li1 25(Mn Ni18)871O2 Li125(Mn Ni1_5)07102
p 4.11 g/cm
20 atom %
Rhombohedral phase
(space group R3rn)
p 4.4 g/cm
p=
E
4.80
X-ray density at the monoclinic/monoclinic-hexagonal
phase boundary. On the other side, data points for x 1.14
could be fit by a straight line which suggests that the
hypothetical Mn05Ni85O is cubic and has an X-ray density
of about 6.55 g/cm3. From Fig. 16, we could estimate the
density and the Li/(Mn + Ni) ratio of the two-phase sputter targets. The results are listed in Table IV.
The comparison of the x-ray densities (Fig. 17) of the
parent series LiMn2_O7 and LrNj2xO7 over the full range
0 x 1.33 with our U(Mn5Ni15)2_O2 (0.38 y 0.5)
phases clearly demonstrates a closer structural relationship of the latter to LiNi2O7 rather than to LiMn2,O7.
We therefore expect Li(Mn5Ni1.5)2O2 to adopt cubic
symmetry below x 0.6, as is the case for LiNi2_O2 for
x u 0.62.728 The most important result inferred from
Fig. 17, however, is the linear dependence of the X-ray
density of both series LiNi2_O2 and Li,,Mn7.O7 on the x
parameter, irrespective of the crystal systems that are
adopted over the range 0 x 1.33. As a consequence, the
fitted curve for Li(Mn5Ni1_5)2O2 within the range 0.93
x 1.26 where 0.38 y 0.50 (Fig. 16) can be reasonably
extrapolated to 0 x < 1.33. Hence, the least-square function for the X-ray density
(Fig. 16)
Px-r.y = (—1.920 X + 6.549) g/cm3
I
[15]
where V is the unit cell volume EAI in hexagonal setting,
and y represents the stoichiometric parameter of the tran-
,
hexagonal
' mono-
:
chnic -
,
Fit = (6.549 -1,920 x) 9/cm3
mono
....... -
fR=O.999
,l,___.,,l,,,, I,,
0.9
1
monoci.
NiO (cubic)
El-..
LiNiO2
1.1
1.2
a value of 6.55 g/cm3 for Mn0,5Ni0,50. The boundaries of the two-
phase region at x 1.15 and x = 1.25 correspond to Li/M =
1.65 (Fig. 14), respectively. Gibbs' phase rule
predicts a sloping X-ray density in the iwo-phase region (see text).
0.
Li2NiO3
o-
(morioci)
0
>.
0 4.00
x
(ortholhombic)
-
Li,Mn03
(monad.)
3.00
0
1.3
Fig. 16. X-ray density of Li,,M2O2 (M = Mn8Ni1..7 and 0.38
y 0.5). The straight-line fit relates to x 1.14 and would yield
(hexagona')
0
C 5.00
0
x in LiXM2..X02
1.35 and Li/M
7.00
hexagonal
0..
hex
4.00 -
[14]
becomes valid over the entire range x < 1.33 (Fig. 17) and
yields the x parameter according to
x (3'7.4ly + 6.549V — 744.l8)/(l8.'72y + 1.920V — 257.88)
6.00
4.40
4.20 -
p = 4.25 g/cm3
For the Li/(Mn + Ni) ratio of each phase cf. Fig. 16. Only the
overall Mn/Ni ratio of a two-phase region is known, the Mn/Ni
ratios of each phase may be different.
E
Ar.
4.60
4.11 g/cm
60 atom %
p 4.4 g/cm
4.16 g/cm3
cubic
C,,
p
Li1 15(Mn Ni17)081O2 Li185(Mn Ni1_9)8,8102
X-ray density of target
4.88 .,
Li123Mn837N1840O2
0.2 0.4 0.6 0.8
1
1.2 1.4
x in LiXM2XO2
Fig. 17. X-ray density of Li,M2_,,02 (M = Mn8,Ni1_7 and 0.38
y
0.5) in comparison with the series of the y parameter end
members Li,,Mn2_,02 and Li,,Ni2_,O2. The crystal systems for each
Ax range of Li,,Ni2,.,O are given on top of the figure. Areas
between the vertical dashed lines indicate the mandatory Iwo
phase regions of Li,,Ni2_02 that have not yet been reported.
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4157
J Electrochem. Soc., Vol. 145, No. 12, December 1998 The Electrochemical Society, Inc.
sition metals. This work proves the Eq. 15 holds true at
least for 0.38 y 0.50. Based on LiXM2.202 with M being
a pseudotransition metal element, e.g., M = MnNi1,
Eq. 15 provides a powerful tool for determining the stoi-
ysis or XRD measurements, we relied on two fixed conditions for the quantitative interpretation of the entire TGA
experiment: the overall cation stoichiometry of the final
requires PIGE measurement in the thin film case.
could be met by the quantitative thermogravimetric analy-
chiometric x parameter of lithium which otherwise
Thermogravimetric (TGA) studies .—As analyzed by
XRD, the fine powder that was initially filled into the TGA
balance consisted of a monoclinic phase (Li2NiO3-related,
space group CZ'c), a rhombohedral phase (LiNiO2-related,
space group R3m), unreacted NiO, and unreacted Li2CO2.
From the weight change during the first run (see below), the
amount of urireacted NiO and Li2CO2 in the fine powder
was calculated to be less than 3 wt %. The overall cation
stoichiometry could be approximated by the composition of
the starting chemicals, Li2 42Mn245Ni042. After heating the
mixture in atmospheres according to the sequence
air—N2—air--N2 (Fig. 18), we foun4 a rhombohedral phase
[a =
2.872(1)
A, c =
14.261(5)
A] and
small amounts of Li20 and Li2PtO3 in the final XRD powder pattern. The latter product originated from the reaction
between Li2CO2 and Pt in air at elevated temperatures, as
was also reported in the literature.29 The overall cation stoi-
chiometry of the final TGA reaction products was determined by ICP analysis to be Li142Mn942Ni247. The actual
amount of lithium that exclusively belonged to the rhombohedral Li(MnNi1.9)2_O2 phase was calculated from Eq. 15
using p =
0.421(0.42
+ 0.47)
= 0.47 which yields
Li2 12Mn942Ni94702. Because the nominal overall starting
cation stoichiometry Li242Mn242Ni042 was close to the analy-
sis of the final products, Li1 45Mn042Ni047, the lithium loss
due to evaporation had to be negligible. We also assume that
Mn/Ni = 0.42/0.47 =
TGA runs.
0.89
was valid throughout all of the
Because the treatment sequence, air—N2--air--N2, was car-
ried out on the sample without intermediate chemical anal100
07 it
products, Li145Mn042Ni042, and the stoichiometry of the final
rhombohedral phase Li2 22Mn942Ni24702. These requirements
sis only if less than 1 atom % of the lithium ions were
assumed to have been present as Li2PtO3 that formed dur-
ing the first TGA run in air and remained unchanged in
subsequent runs.
In the first run (Fig. 18a),the powder was heated in air
Between 30-280°C, moisture which was absorbed during
air exposure was lost, and between 280 and 800°C the following weight loss occurred
Li2CO, + 2NiO +
O2
—*
2LiNiO2
+ co9t
Li2CO3 + Pt + 024 -÷ Li2PtO2 + C021
[16]
[17]
Li2CO2 -+ Li20
[18]
C02t
The small weight loss during the 1 h dwell time at 800°C is
attributed mostly to these reactions. After cooling to room
temperature, we obtained Li1 24Mn9 36Ni240O2. The oxygen
stoichiometry was inferred from the observation above
that these preparation conditions produced a single-phase
or a two-phase mixture of the series Li(MnNi2_9j202 in
thin-film and bulk phases.
In the following run, heating in N2 (Fig. 18b) resulted in
severe oxygen loss starting above 680°C, possibly by a reaction mechanism similar to Eq. 13. This is supported by the
fact that coprecipitated Li20 was found in the XRD powder
pattern after the final run of the TGA sequence which was
also carried out under N2. According to the proposed reaction mechanism, the quantitative analysis gave 82 atom %
rhombohedral Li197Mn944Ni04902 and 18 atom % Li20 when
cooled to room temperature.
The third run was performed in air (Fig. 18c) where most
of the lost oxygen was reabsorbed above 280°C under concurrent Li20 reinsertion into the lithium-manganese-nickel oxide phase. We believe that the mass loss between 560-
640°C was due to CO2 evaporation competing with 02
absorption. The CO2 evaporation stemmed from decomposing Li2CO2, which probably formed between Li20 and
air below 560°C. Further 02 absorption was observed when
cooled to room temperature thus resulting in 95 atom %
Li2 29Mn939Ni04302 and 5 atom % Li20.
In the final treatment under N2 (Fig. 18d), the powder
sample behaved as in the previous N2 run. The main features comprised the strong 02 loss under concurrent Li20
coprecipitation beginning at 680°C. The room-temperature
phase distribution was 83 atom % rhombohedral
Li2 08Mn044Ni949O2 and 17 atom % Li20.
1uu
c)
The reason why the stoichiometry Li1 11Mn942Ni047O2,
calculated from Eq. 15 and the unit cell parameters, could
not be exactly matched by our quantitative interpretation
of the TGA lies in the experimental error of the ICP
97.5
analysis. If a higher amount of Li2PtO2 than the aboveestimated one is assumed to have been present,the results
are worse. Nonetheless, the TGA results clearly demonstrate the high reversibility of oxygen release/reinsertion
95
ci)
under simultaneous Li20 coprecipitation/reinsertion in
the series Li(Mn9Ni1_9)2O2 as a function of the oxygen
partial pressure.
XPS measurements—The XPS measurements were con-
97.5
ducted to obtain additional information on the Mn/Ni
_____________
95
120
280
440
600
760
Temp (°C)
Fig. 18. Thermogravimetric analysis (TGA) of lithium manganese
nickel oxide (a) in air, (b) under flowing N2, (c) in air, and (ci) under
flowing N2. Heating and cooling rates were set to 10°C/mm and the
maximum temperature (800°C) was held for 1 h prior to cooling.
ratio as well as on the Mn and Ni valence states at the surof the films. The latter are of interest for the interpretation of the initial voltages in freshly prepared lithium
batteries!2
The Ni (2p) photoelectron emission spectrum is given in
Fig. 19. The peaks located at 854.3 eV/860.8 eV and at
871.9 eV/878.3 eV are attributed to the spin-orbit splitting
of the Ni (2p) components, Ni (2P312) and Ni (2P112), respectively. The satellites at about 860.8 and 878.3 eV originated
from shake-up lines39 which are commonly observed with
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4158
J. Electrochem. Soc., VoL 145, No. 12, December 1998 The Electrochemical SocIety, Inc.
4.0
Also shown in Fig. 21 is the Li (is) photoemission spectrum around 54.0 V. This peak could be fitted by a single
peak which therefore gave the binding energy of Li (is) in
the Li-Mn-Ni-O phase at the surface of the film, as the
854.3
071 .9 oV
Nil'
860.8ev
878.3eV
3.5
amount of Li2CO3 and LiOH was found to be small, as concluded from the C (is) and 0 (is) spectra (not shown). Only
3.0
upon extended exposure to air a very strong new Li (is)
peak appeared at 55.7 eV (not shown) which matched the
Ni 2py
Ni 2p1/2
0.5
value for Li2CO3 reported in the literature.30'34
For the determination of the Li/Mn/Ni ratio at the surface
of the film, we used the Li (is) peak at 54.0 eV (Fig. 21), the
0
'
890
880
'
870
'
860
850
entire Ni (2p) area (Fig. 19), and the entire iVin (3p) area
(Fig. 20) yielding an overall cation stoichiometry of about
Li117Mn039Ni044. The stoichiometric factor of lithium, howev-
er, cannot be considered to be exact due to the slight overlap
of the Li (is) and Mn (3p) areas in Fig. 21, but it gives a good
Binding Energy (eV)
Fig. 19. X-ray photoemission spectrum (XPS) in the Ni (2p) region
of the film with the bulk stoichiometry Li;,1oMno.39Ni0.O2. The fit
estimate of the stoichiometric conditions at the surface.
envelope results from adding up the asymmetric Gaussian envelopes
of the deconvoluted XPS peaks. The dots below the spectrum
approaches reasonably closely the cation bulk stoichiometry
indicate the difference in counts belween the measure and the fitted data.
common observation that the bulk and surface may differ
paramagnetic ions like Ni2(3d8). For transition metal spec-
tra with prominent shake-up lines, the entire 2p region has
to be included for quantitative analysis, because the total
amount of the respective ion species is equal to the integral
number over all Ni (2p) states. As can be seen from Fig. 19,
the Ni (2p1) main peak at 854.3 eV was fitted easily by a
single peak which agreed well with data reported for
Ni2. 31-33 The average valence state of Mn and Ni in the
bulk of Li1 10Mn039Ni051O2 was calculated to be +3.22, indi-
cating that some Ni3* had to be present in the film sample.
According to the XPS results, however, there was no indication for such species at the film surface. In order to final-
ly match the bulk stoichiometry of the film, the nickel
valence states had to increase with film depth.
For the quantitative analysis of Mn, we used the area
under the Mn (2p) and Mn (2piiz) peaks at 641.8 and
653.5 eV (Fig. 20), respectively. Valence states of Mn, however, are better resolved in the Mn (3p) emission spectrum
(Fig. 21), because the 3p electrons belong to the outer shell
and therefore better reflect the binding state than the 2p
electrons. This was already demonstrated by Topfer et al.31
Therefore, we deconvoluted the Mn (3p) area into three
peaks located at 48.0, 49.4, and 50.8 eV which corresponded to Mn2', Mn3, and Mn4 according to Ref. 31. The presence of Mn2 at the surface of the film, however is uncer-
Overall, the cation surface stoichiometry of Li1 17IVIn039Ni044
of the film, Li1 10M.n039Ni051, and is in accordance with the
slightly in stoichiometry.
Summary
The series Li(Mn0Ni1_5)3_O3 for 0.38 < y 0.50 and x
1.33 showed a very close structural relationship to its
parent series LrN2.O2. Chemical preparation below
800°C in atmospheres with partial pressures higher than
10 atm 02 evidently led to phases wherein the total
number of cations equaled the number of oxygen anions,
irrespective of the adopted crystal system, which in turn
depended on the stoichiometric parameter x. In contrast
to LjrNj3O3, phases of the series Ljr(MflyNji_y)2O2 for
x 1.0 could be easily prepared in air or flowing 03 at
700-800°C. Starting from the as-deposited film stoichiometry of Li0(Mn4Ni1_5)3_0O3, higher annealing tempera-
tures and lower oxygen partial pressures tended to decrease the stoichiometric parameter of lithium, but the
rule "cations/anions = 1" did not appear to be violated
under the above-mentioned conditions probably because
of the system's self-regulation mechanism: by reversibly
coprecipitating/reinserting Li30 while releasing/taking
up 03 the Lir(MngNii_v)2x03 phases were able to adjust
the number of cations to an undisturbed oxygen sublattice. Whereas films prepared under low oxygen partial
pressures such as flowing N3 contained Li30 as an ad-
ditional phase, the oxygen pressure in air was high
enough to keep all of the lithium within the actual
tain owing to the very small full width at half-maximum
(fwhm) of the deconvoluted peak, which is only slightly
Li(Mn5Ni1_5)2O2 phases.
0.9 eV at the used pass energy of 11.75 eV. The calculation
of the mean valence state of manganese yields about Mn3'3
and is barely affected by the very small area contribution
underneath the peak at 48 eV.
linear function of the ratio Li/(Mn + Ni), which in turn
was proportional to the averaged valence state of the
larger than the instrumental line-broadening of about
The refined lattice parameters of rhombohedral
Li(Mn5Ni15)3O3 for x < 1.15 (space group R3m) were a
3.0
00
C
0 2.5
0
0
S5
C
w
C
2.0
2Pit51XP;:J8ev
0
'A
C
'A
C
58
660
655
650
645
640
635
Binding Energy (cv)
Fig. 20. X-ray photoemission spectrum (XPS) in the Mn (2p) region of the film given in Fig. 19.
56
54
52
50
48
46
Binding Energy (cv)
Fig. 21. X-ray photoemission spectrum (XPSI in the Li (Is) and
the Mn (3p) regions of the film given in Fig. 19. The presence of
Mn2 is uncertain due to the small fwhm of the deconvolutecJ line at
48 eV.
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J. Electrochem. Soc., Vol. 145, No. 12, December 1998 The Electrochemical Society, Inc.
pseudotransition metal M (M = MnNi1_). Monoclinic
Li(Mn5Ni1_)2_O2 (space group C2/c) nucleated as soon as
lithium exceeded the threshold concentration x = 1.15
(0.38 y 0.50). Further increasing the concentration to
x> 1.25, i.e.,> 25% of the sites in the predominantly filled
transition metal layers were occupied by Li ions, converted the two-phase system into a monoclinic single
phase. A direct comparison of rhombohedral and monoclinic lattice parameters could be achieved by transforming the monoclinic unit cell into the pseudo-hexagonal
cell. Based on the linear relationship between X-ray
density and stoichiometric parameter x over the entire
range 0 s x 1.33, the lithium concentration of
Li(MnNi1_)2_O2 thin films for 0.38 y 0.50 can now
be determined with confidence without requiring ionbeam analysis.
Quantitative and qualitative analysis of XPS measurements on a Li1 10Mn039Ni051O2 thin film proved that the
cation stoichiometry of the surface was almost the same as
in the bulk of the film. However, the Li2CO3 formed on the
surface upon extended exposure time to air deprived the
Li-Mn-Ni-O phase of lithium oxide. Although Mn4 and
Mn were found at the surface of Li110Mn039Ni051O2, there
was no evidence for nickel valence states other than Ni2.
Judging from the bulk stoichiometry of the film, which
required at least Ni3 in addition to Ni2, the discrepancy
between bulk and surface stoichiometry indicated that the
average nickel valence state increased with film depth.
Acknowledgments
This research was supported by the Department of
Energy's Division of Materials Sciences, Division of Chemical Science, Office of Energy Research 'Ibchnology Transfer
4159
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