Hollow microspherical layered xLi2MnO3·(1-x)LiNiO2 (x=0.3–0.7) as cathode material for lithium–ion batteries

Article abstract of DOI:10.1016/j.jallcom.2019.03.005

Li-rich layered Li₂MnO₃ is of great attraction for high energy lithium ion batteries. However, its cycling is still needed for improvements. Here we report a hollow microsphere-structured xLi₂MnO₃·(1-x)LiNiO₂ (x = 0.3–0.7) that is synthesized by using in-situ template-sacrificial strategy. Powder X-ray diffraction (XRD) and scanning electron microscope (SEM) characterizations prove that the xLi₂MnO₃·(1-x)LiNiO₂ (x = 0.3–0.7) are based on monoclinic Li₂MnO₃ with α-NaFeO₂ layered structure in which Li⁺ ions are orderly arranged in the transition metal layers, and the hollow-microspheres have diameters of ～3 μm. Electrochemical results show that the optimal ratio of Li₂MnO₃/LiNiO₂ is 0.6/0.4. As a consequence, the stabilized discharge capacity of 0.6Li₂MnO₃·0.4LiNiO₂ (0.6LLMNO) is ～210 mAh g^?1 after the first few cycles. This shows that appropriate amount Ni substitution for Mn in Li₂MnO₃ helps to improve the specific capacity and cycling stability.

Full text of DOI:10.1016/j.jallcom.2019.03.005

Accepted Manuscript
Hollow microspherical layered xLi MnO ·(1-x)LiNiO (x=0.3–0.7) as cathode material
2
3
2
for lithium–ion batteries
Kai Zhang, Lei Zhang, Junjie Liu, Xiongwei Wu, Chunjiao Zhou, Wenqi Yan,
Congshan Zhou, Lijun Fu, Yuping Wu
PII:
S0925-8388(19)30809-6
DOI:
https://doi.org/10.1016/j.jallcom.2019.03.005
JALCOM 49769
Reference:
To appear in: Journal of Alloys and Compounds
Received Date: 8 December 2018
Revised Date: 20 February 2019
Accepted Date: 1 March 2019
Please cite this article as: K. Zhang, L. Zhang, J. Liu, X. Wu, C. Zhou, W. Yan, C. Zhou, L. Fu, Y. Wu,
Hollow microspherical layered xLi MnO ·(1-x)LiNiO (x=0.3–0.7) as cathode material for lithium–ion
2
3
2
batteries, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/j.jallcom.2019.03.005.
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ACCEPTED MANUSCRIPT
Hollow microspherical layered xLi2MnO3·(1-x)LiNiO2 (x=0.3-0.7) as cathode
material for lithium–ion batteries
Kai Zhang1, Lei Zhang2, Junjie Liu1, Xiongwei Wu1,*, Chunjiao Zhou1, Wenqi Yan3,
Congshan Zhou2,*, Lijun Fu 3, Yuping Wu1,3,4,
*
1 College of Science, Hunan Agricultural University, Changsha 4101289, China
2 College of Chemistry and Chemical Engineering, Hunan Institute of Science and
Technology, Yueyang 414006, China
3 State Key Laboratory of Materials-oriented Chemical Engineering & School of
Energy Science and Engineering, Nanjing Tech University Nanjing, Nanjing 211816,
China
4 DKJ New Energy Tech Co. Ltd., Shaoxing 312000, China
Email: zhoucongsh@126.com, wuyp@fudan.edu.cn
Abstract
Li-rich layered Li2MnO3 is of great attraction for high energy lithium ion batteries.
However, its cycling is still needed for improvements. Here we report a hollow
microsphere-structured xLi2MnO3·(1-x)LiNiO2 (x=0.3-0.7) that is synthesized by
using in-situ template-sacrificial strategy. Powder X-ray diffraction (XRD) and
scanning electron microscope (SEM) characterizations prove that the
xLi2MnO3·(1-x)LiNiO2 (x=0.3-0.7) are based on monoclinic Li2MnO3 with α-NaFeO2
layered structure in which Li+ ions are orderly arranged in the transition metal layers,
and the hollow-microspheres have diameters of ~3 ꢀm. Electrochemical results show
that the optimal ratio of Li2MnO3/LiNiO2 is 0.6/0.4. As a consequence, the stabilized
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discharge capacity of 0.6Li2MnO3·0.4LiNiO2 (0.6LLMNO) is ~210 mAh·g-1 after the
first few cycles. This shows that appropriate amount Ni substitution for Mn in
Li2MnO3 helps to improve the specific capacity and cycling stability.
Keywords: Lithium ion batteries; Li-rich layered material; hollow-microspheres; high
capacity.
1. Introduction
The secondary batteries have been upgraded several times from traditional lead-acid
batteries to nickel-metal hydrogen batteries, and then to the latest lithium-ion
batteries.[1-5] Lithium-ion batteries are widely used owning to their high energy
density, long cycle life and no memory effect, which are not only the dominant power
source for portable electronics but also satisfy the requirements of the rapid
development for large-scale applications such as electric vehicles (EVs) and hybrid
electric vehicles (HEVs).[6-8] Among the materials for the lithium-ion batteries, the
cathode material not only participates in the electrochemical reaction as the electrode
but also act as the source for Li+ ions. However, most commercial lithium-ion
batteries use layered LiCoO2 or LiNi1/3Co1/3Mn1/3O2 as the cathode, which has limited
practical capacity.[9-11] Therefore, there is an urgent need to develop new cathode
materials with higher operating voltage, higher capacity as well as better cycling
stability.
The cathode for lithium-ion batteries needs stable structure in wide voltage range, so
that more Li+ ions can be intercalated/de-intercalated during continuous
charge-discharge cycles to maximize the capacity, energy density, and cycle life for
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lithium-ion batteries.[12] At present, layered Li2MnO3 has drawn much attention due
to its high capacity. If Li+ ions completely intercalate/de-intercalate between 2-4.8 V,
the ideal capacity of Li2MnO3 can reach 459 mAh·g-1.[13-16] But it is well known
that the valence state of Mn is +4 and Li2MnO3 is electrochemically inactive. During
the initial charge of Li2MnO3 at 4.5 V oxygen is released and Li+ ions de-intercalate to
form irreversible "Li2O" and electrochemically active MnO2.[17] With the release of
oxygen from the material, Li+ ions migrate from the octahedral sites of the transition
metal layers to the tetrahedral sites of the lithium layers, and octahedral vacancies in
transition metal layers are formed. At this time, the lattice structure is unstable,
leading to transition metal ions to migrate and occupy the octahedral vacancies until
the end of oxygen release. This is the rearrangement of transition metals in the
material, which usually leads to the capacity fading.
Previous studies have shown that surface coating can improve these defects to some
extent.[18-22] However, it can simply increase the Li+ ions diffusion kinetics in the
surface area rather than stabilizing the lattice structure in the bulk material. In this
regard, researchers have rationally designed Li-rich compounds doping with different
cations.[23-25] The monoclinic Li2MnO3 and rhombohedral LiMO2 (where M
represents transition metals ions) form layered xLi2MnO3·(1-x)LiMO2 compound that
shows higher capacity and better cycling stability.[26-30] As previously reported,[31]
doping Fe+3 and Cr+3 in Li[Li0.2Mn0.4Co0.4-xMx]O2 can reduce oxygen loss during the
initial charging process. However, since the doping of Fe+3 causes the decrease of the
metal-oxygen covalence which results in a greater irreversible capacity loss during the
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initial cycle. Cr+3 doping increases the charge capacity below 4.5 V due to the three
electron redox of Cr+3/Cr+6 but the capacity above 4.5 V decreases. Similarly Co3+
doping in Li[Li0.2Mn0.6-xNi0.2-yMx+y]O2 increases the capacity at ~4.5 V and hence the
discharge capacity, which through the overlap of the Co+3/4+ bands with the top of the
O-2 bands to increases metal-oxygen covalence.[32,33]
In the traditional Li-rich layered cathode materials, the electrons of d-orbital in the
transition metal ions directly participate in the redox reaction. Meanwhile, the
intercalation/de-intercalation process of Li+ ions preserves the charge balance. If the
electronegativity of transition metal ions is lower, the electrons in the d-orbital are
more likely to be lost, which result in lower operating voltage. Therefore, the amount
of d-orbital electrons involved in the redox reaction play a crucial role in deciding
charge-discharge capacity.[34-37] It is well known that Ni ions exhibit higher
electronegativity and ionization energy than Mn ions under the same condition, and
the Ni-O band tends to exhibit higher covalence characteristics under high-voltage.
[38, 39] Therefore, the substitution of some Mn ions by Ni ions in Li2MnO3 can
improve the capacity and cycling stability. In this work, we adopt a facile in-situ
template-sacrificial
strategy
to
synthesize
hollow-spherical
layered
xLi2MnO3·(1-x)LiNiO2 (x=0.3-0.7), the samples were synthesized method as similar
with reported previously.[40] The influence of Ni ions substitution for Mn ions in
Li2MnO3 on the structure, morphology, and electrochemical performance of the
as-prepared Li-rich layered material was investigated.
2. Experimental
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2.1. Material synthesis
MnCO3 microspheres were synthetized via a precipitation method. Briefly, two stocks
of MnSO4·H2O (10 mmol) and NH4HCO3 (100 mmol) were separately dissolved in
700 ml deionized water, then 70 ml ethanol and the NH4HCO3 stock solutions were
poured into the MnSO4·H2O stock solution under continuous stirring at room
temperature for 2 h. The reaction mixture became cloudy. The as-obtained
milky-white MnCO3 microspheres were thoroughly filtered and washed with
deionized water and ethanol for three times. Then MnCO3 microspheres were dried in
a vacuum oven at 80 °C overnight. Finally, MnCO3 microspheres were calcined at
400 °C for 5 h in the air at a heating rate of 1 °C·min-1 to obtain MnO2 microspheres.
The MnO2 microspheres were used as template precursors to synthesize
xLi2MnO3·(1-x)LiNiO2 (x=0.3-0.7). According to the desired stoichiometry ratio, the
MnO2 was added into the ethanol solution of LiOH·H2O and Ni(NO3)2·6H2O, and the
ethanol was evaporated slowly at 50 °C under stirring. The mixture was manually
ground again, and then calcined at 800 °C for 12 h in the air at a heating rate of
2 °C·min-1 to obtain xLi2MnO3·(1-x)LiNiO2 microspheres (an excess of 5%
LiOH·H2O was added in order to prevent the loss of Li in the synthesis). For
comparison, xLi2MnO3·(1-x)LiNiO2 (x=0.3-0.7) samples were named as xLLMNO
(x=0.3-0.7).
2.2. Measurements
The crystal structure of xLLMNO (x=0.3-0.7) was characterized by X-ray diffraction
(XRD) at the scan rate of 10 °·min-1 in the range of 10° ~ 80°, which was performed
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on a Bruker D4 X-ray diffractometer with Ni-filtered Cu Kα radiation (40 kV, 30 mA).
The surface morphology of xLLMNO was observed by scanning electron microscope
(SEM) of a Philip XL30 microscope operated at 10 kV. Specific surface areas and
pore size distribution of the pristine xLLMNO were measured by the multipoint
Brunauer-Emmett-Teller (BET) method using Micrometrics ASAP 2020 Plus analyser.
The cathode component consisting of 80 wt.% active material (xLLMNO), 10 wt.%
carbon black (CB) and 10 wt.% polyvinylidene fluoride (PVDF) was mixed in
N-methyl-2-pyrrolidone (NMP) to form slurry by magnetic stirring, and the slurry
was coated onto the current collector of aluminum foil. The coated film was vacuum
dried at 80 °C overnight, and used as the cathode after cutting according to the
required size. The mass loading of the active material as about ~2.5 mg·cm−2. The Li
metal was used as anode and the Celgard 2400 film was used as the separator. The
electrolyte was 1 M LiPF6 solution dissolved in ethylene carbonate (EC), diethyl
carbonate (DEC) and dimethyl carbonate (DMC) solvent (1:1:1 by volume). All cells
were assembled in a glove box filled with argon gas of high purity. The galvanostatic
charge-discharge tests were performed by Land Tester (CT2001A) with cut-off
voltage between 2.0-4.8 V (vs.Li/Li+) and current density of 0.1 C (1C=200 mAh·g-1).
The cyclic voltammetric curves (CV) were obtained between 2.0 to 4.8 V at a scan
rate of 0.1 mV·s-1. The electrochemical impedance spectroscopy (EIS) was measured
in the frequency range of 100 kHz to 0.1 Hz with a potentiostatic signal amplitude of
5 mV. All electrochemical measurements were implemented at room temperature.
3. Results and discussion
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3.1. Structure characterization and morphology
In order to understand the lattice structural change caused by Ni ions substitution, the
MnCO3, MnO2 and xLLMNO were analyzed by XRD. Comparing the patterns of
MnCO3 and MnO2 with the standard PDF card, the positions of all diffraction peaks
are well consistent with those of the standard ones. All diffraction peaks of MnCO3
can be ascribed to the rhombohedral phase (Fig. 1a). At the same time, the MnO2 peak
intensity is sharp and the half-width of the peak is narrow (Fig. 1b), indicating that the
microspheres template precursors have high purity and good crystallinity which are
favorable for the synthesis of xLLMNO. After 12 hours of calcification with
LiOH·H2O and Ni(NO3)2·6H2O at 800 °C, the MnO2 template precursors developed
into xLLMNO.
Fig. 2a shows that the main peaks for XRD of xLLMNO have high crystallinity, and
ꢀ
well matched to the phase of the α-NaFeO2 (R3m space-group). Furthermore, with the
increase of Li2MnO3, the super-lattice peaks between 20° ~ 23° becomes gradually
clear (Fig. 2b), which indicates that proved the xLLMNO is a monoclinic (C2/m
space-group) Li2MnO3–like compound.[41] According to the previous report,[42]
Li2MnO3 can be expressed as Li(Li1/3Mn2/3)O2, in which Li+ ions and Mn ions with a
1:2 ratio are orderly arranged in the transition metal layers, and each LiO6 octahedra
is surrounded by six MnO6 octahedron, which result in an orderly arrangement of
LiMn6 in the form of hexagon. It can be well understood that the xLLMNO is based
on monoclinic Li2MnO3 with α-NaFeO2 layered structure compound, in which Li and
M occupy Na and Fe sites, respectively.[43,44] Fig. 2c and 2d shows that the peak
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separations of (006)/(102) and (018)/(110) gradually become evident, which also
proves that a highly ordered layered structure of hexagon is formed as the Li2MnO3
content increases. Consequently, we can conclude that the Li-rich layered compound
of xLLMNO was successfully synthesized by the in-situ template sacrificial strategy.
The morphology of microspheres template precursors and xLLMNO can be evidently
observed by SEM (Fig. 3a, b). MnCO3 was calcined at 400 °C to form MnO2
according to the following equation:
2MnCO3+O2=2MnO2+2CO2
(1)
The release of CO2 during the calcification process produces a lot of micropores on
the surface of MnO2, and the surfaces of each microsphere is smooth and compact
(arrow mark). The particle size distribution shows that the average diameter of
MnCO3 and MnO2 are ~3 ꢀm (Fig. 3c).
Fig. 4 shows the SEM images and the particle size distribution histogram for the final
products of xLLMNO. The microspheres with diameters of ~3 ꢀm are still retained
after the calcification. However, due to the Kirkendell effect during high-temperature
calcination, the hollow sphere structure is formed by the imbalance diffusion rate of
different ions at the reaction interface. This can be divided into two stages: the first
stage is mainly to generate Kirkendall voids through the inside diffusion process; the
second stage is the diffusion of rapidly diffusing species along the surfaces of the core
materials,[45] so the surface of xLLMNO becomes rougher than that of MnO2
precursor. During the synthesis process of the sacrificied template, Mn ions and Ni
ions diffuse rapidly outward while the O diffuses slowly inward, which result in the
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formation of hollow xLLMNO microspheres (Fig. 5).[40]
At the same time, Fig. 6 shows that the high-magnification SEM images and
elemental mappings of xLLMNO. It can clearly observe that Mn, Ni and O elements
are uniformly distributed for the hollow-microspheres. Meanwhile, with the Mn
content increases, the surface morphology of the microspheres tend to be smoother.
More importantly, the hollow-microspheres structure will also present a significant
improvement to its electrochemical performance, and the main reasons are: 1) the
extra Li+ ions in its inner wall of cavities are activated to provide more reversible
capacity during cycling; 2) the larger specific surface area has better rate capability; 3)
the hollow-microspheres structure is more stable, which can buffer the volume change
caused during the charge-discharge process. [39, 46-48]
Fig. 7 shows nitrogen adsorption–desorption isotherms. It is clear that the
microspheres exhibit typical IUPAC type IV curves, which suggest the existence of
mesopores.[35] Compared with 0.3LLMNO (2.765 m2·g-1), 0.4LLMNO (3.642
m2·g-1), 0.5LLMNO (6.459 m2·g-1) and 0.7LLMNO (7.118 m2·g-1), 0.6LLMNO
(11.369 m2·g-1) has the largest specific surface area, which can shorten the diffusion
path of Li+ and improve electrochemical performance. At the same time, Fig. 7b
shows the pore distribution curves. It is obvious that the mesoporous size is about
0.5-3.5 nm, which can promote the penetration of the electrolyte into the electrode
particles and promote rate performance. [12]
3.2. Electrochemical performance
The initial charge capacity in voltage range of 2.0 to 4.8 V at 0.1 C for all xLLMNO
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samples can be divided into two sections. The first section below 4.5 V corresponds to
+
ꢀ
de-intercalation of partial Li ions from the structure of space group R3m together
with the oxidation reaction of the Ni2+ and Ni3+ to Ni4+, which is the main origin of the
reversible capacity.[49] In the section above 4.5 V, it is ascribed to Li2MnO3 voltage
platform, which corresponds to de-intercalation of another part of Li+ ions from the
layered Li2MnO3 lattice (C2/m space-group) and rearrangement of the transition metal
ions, and further oxygen removal from Li2MnO3 to form irreversible "Li2O". At the
same time, in order to maintain the charge balance, the transition metal ions on the
surface of the material migrate to occupy the Li+ site inside the bulk phase, so that the
de-intercalation Li+ cannot be completely intercalation in the original site, which leads
to produce large irreversible capacity loss. [50] As can be seen from Fig. 8a, the
second voltage platforms above 4.5 V in initial charge for 0.3LLMNO, 0.4LLMNO,
0.5LLMNO are not apparent. Meanwhile, accompanied with the gradually increase of
Li2MnO3 content, the initial charge curve of 0.6LLMNO and 0.7LLMNO have two
distinct charge voltage platforms. This is mainly due to that xLLMNO are gradually
formed based on monoclinic Li2MnO3 with α-NaFeO2 layered structure as the content
of Mn increases. In addition, it is noted that as the Ni content decreases, the
irreversible capacity in the initial charge also decreases (Table 1). The 0.3LLMNO
sample shows the most serious capacity loss of 62 mAh·g-1 in the initial charge, but
the irreversible capacity for 0.7LLMNO is only 34 mAh·g-1. Such a high irreversible
capacity can be attributed to major lattice structure evolution, including oxygen
removal from layered structure and rearrangement of the transition metal ions.[51] By
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comparing the columbic efficiency of xLLMNO in the 100 cycles (Fig. 8b), it can be
found that the efficiency is significantly improved in the first few cycles and quickly
reached to ~100%, indicating that additional capacity generated by Li2MnO3
activation is almost completely reversible after the first few cycles. The higher
efficiency of xLLMNO in the subsequent cycles could be explained by that any
Li2MnO3 was hardly activated, which can significantly reduce the transition metal
ions occupying Li+ ions vacancies.[52] Accordingly, a moderate amount of Ni ions
substitution for Mn ions in the Li-rich layered material can stabilize the lattice bulk
structure and decreases irreversible capacity during the initial charging process.
Cycling stability between 2.0 to 4.8 V with current density at 0.1 C is shown in Fig.
9a. Among all xLLMNO, the discharge capacity shows an increasing trend with the
increase of Li2MnO3. After activation, the discharge capacity of 0.6LLMNO is
increased from 173 to ~210 mAh·g-1 in the first few cycles, which is mainly because
that oxidation of Ni2+ and Ni3+ to Ni4+ provides sufficient reversible capacity while
the intercalation/de-intercalation of Li+ ions in the inner wall of cavities are
activated.[46,53] However, the discharge capacity of 0.7LLMN after activation is
~130 mAh·g-1, which is due to too low Ni content to provide sufficient reversible
capacity. Fig. 9b shows the rate performance of xLLMNO, which is tested at different
current densities from 0.1 C to 5 C and then back to 0.1 C with 5 cycles at each rate.
The 0.6LLMNO has the best rate performance and the discharge capacity at 0.1 C, 0.2
C, 0.5 C, 1 C, 2 C, and 5 C and back to 0.1C are 173, 153, 142, 119, 103, 55 and 214
mAh·g-1, respectively. Based on the above discussion, there are several reasons why
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0.6LLMNO shows the highest capacity and the best rate performance. On the one
hand, a appropriate amount of Ni ions doping to Li2MnO3 can stabilize the ordered
layered structure and reduce the rearrangement of the transition metal ions to destroy
the layered structure, and the oxidation of Ni2+/Ni3+ to Ni4+ can provide higher
reversible capacity.[54] On the other hand, the hollow-microspheres structure also has
a significant contribution to increasing capacity which because the extra Li+ ions in its
inner wall of cavities are activated to provide more reversible capacity during cycling.
Through the above discussion, the 0.6LLMNO sample exhibits the best
electrochemical performance, which proves the ratio of Li2MnO3/LiNiO2 with 0.6/0.4
is the optimal proportion to prepare the Li-rich layered materials.
CV curves for the initial cycle and the 20th cycle of xLLMNO with scan rate of 0.1
mV·s-1 are shown in Fig. 10. The range of 3.85 ~ 4.15 V is ascribed to the oxidation
of Ni2+/Ni3+ to Ni4+ and that at ~4.5 V corresponds to oxygen removal from Li2MnO3
to irreversibly form "Li2O".[17,49,50] With the decrease of Ni ions, the oxidation
peak at ~4.5 V of 0.6LLMNO and 0.7LLMNO gradually becomes obvious (Fig. 10a).
In the subsequent cycling to the 20th cycle (Fig. 10b), the peak at ~4.5 V of Li2MnO3
disappears completely. Meanwhile, the oxidation peaks of 0.6LLMNO and
0.7LLMNO in 3.85 ~ 4.15 V for the oxidation of Ni2+/Ni3+ to Ni4+ is higher than those
in the initial scan, which is in consistent with the increase of capacity as demonstrated
by the above cycling tests.
EIS is a method for studying electrode kinetics and other electrode characteristics,
such as charge transfer resistance, double-layer capacitance, surface film resistance,
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and Li+ ions diffusion coefficient. The EIS curves consist of semicircles at the
high-to-medium frequency and straight lines in the low-frequency region. The
semicircular at the high-frequency region is related to the resistance of Li+ ions
diffusion transfer on the surface film (Rsf), which involves the spontaneous
electrochemical reaction between the active cathode materials and electrolyte. The
semicircular at the medium frequency region is the charge transfer resistance (Rct),
which is caused by the Li+ ions intercalation/de-intercalation process on the surface of
active material. At the same time, the inclined line in the low-frequency shows the
Warburg impedance (ZW) resistance caused by Li+ ions diffusion in the solid phase of
electrode materials.[55,56]
In order to further understand the charge transfer resistance during the
charge-discharge process, we compared the impedance values of xLLMNO after the
1st and the 20th cycles (Fig. 11. and Table 2). It is well known that small charge
transfer resistance facilitates the reversible intercalation/de-intercalation of Li+ ions
and leads to better rate performance.[57] The charges transfer resistance of
0.3LLMNO, 0.4LLMNO, 0.5LLMNO, 0.6LLMNO, 0.7LLMNO at the beginning are
334, 266, 181, 163 and 247 Ω, and 0.6LLMNO shows the smallest charge transfer
resistance. After the 20th cycle, the charge transfer resistances of all samples are much
lower than those after the 1st cycle. From Table 2, it can be seen that the charges
transfer resistance of 0.3LLMNO, 0.4LLMNO, 0.5LLMNO, 0.6LLMNO, 0.7LLMNO
was 94, 76, 68, 36 and 102 Ω. This confirms that 0.6LLMNO have the best
electrochemical performance during charge–discharge. Here, the value of the charge
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transfer resistance for xLLMNO after the 1st cycle is much larger, but it decreases in
the subsequent cycles. The exact reason of this result and the change in surface
kinetics of xLLMNO are not clear. However, it can be concluded that after long-term
cycling, the electrochemical properties of xLLMNO have not been seriously
deteriorated, and the reversible capacities of 0.6LLMNO and 0.7LLMNO increase
after the first few cycles.
Through the above discussion, the relationship between the electrochemical
performances between Ni content of these Li-rich cathode materials can be
summarized. If the Ni content is too high in xLLMNO, more transition metal ions
occupy the Li ions sites in the initial cycle to generate more irreversible capacity. In
contrast, the xLLMNO with low content of Ni leads to the oxidation of Ni2+/Ni3+ to
Ni4+ cannot provide sufficient reversible capacity meanwhile Li2MnO3 cannot be
activated effectively. Therefore a appropriate content Ni ions doping to Li2MnO3 is
favorable to kinetically activate, which can not only stabilize the ordered layered
structure while reduce the rearrangement of the transition metal ions to avoid destroy
of the layered structure but also provide sufficient reversible capacity.
4 Conclusions
In this work, we adopt a facile in-situ template-sacrificial strategy to synthesize
hollow-spherical xLLMNO (x=0.3-0.7) of layered structure. As the content of
Li2MnO3 increases, the super-lattice peak between 20° ~ 23° becomes gradually
evident. It represents monoclinic (C2/m) Li2MnO3–like component which prove the
orderly arrangement of Li+ ions in the transition metal layers. The discharge capacity
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of 0.6LLMNO is increased after the first few cycles and maintained at ~210 mAh·g-1.
Through the above discussion, the ratio of Li2MnO3/LiNiO2 for 0.6/0.4 is the optimal
proportion to synthesize Li-rich layered materials. As a result, an appropriate Ni
content can not only stabilize the ordered layered structure but also improve
electrochemical performance. This work provides basic insights to optimize Li-rich
layered oxides, which will advance the rational design and synthesis of Li-rich
layered cathode materials for lithium-ion batteries.
Acknowledgment:
The research was supported by the funding of National Materials Genome Project
(2016YFB0700600), National Natural Science Foundation Committee of China
(Distinguished Youth Scientists Project of 51425301, U1601214, 51573013,
51773092 and 51772147), Research Foundation of State Key Lab (ZK201805),
Double First-Class Construction Project of Hunan Agricultural University (Grant no.
SYL201802008).
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Figure captions
Fig. 1. (a) XRD patterns of MnCO3 (JCPDS Card No.: 86-0173), (b) MnO2 (JCPDS
Card No.: 89-5171).
Fig. 2. (a) Full XRD patterns of xLLMNO and local magnified XRD patterns at
(b) 20° ~ 23°, (c) 37.5° ~ 39° and (d) 63.5° ~ 66°.
Fig. 3. SEM images of (a) MnCO3, (b) MnO2 and (c) the particle size distribution
histogram.
Fig. 4. The low-magnification SEM images of (a) 0.3LLMNO, (b) 0.4LLMNO, (c)
0.5LLMNO, (d) 0.6LLMNO and (e) 0.7LLMNO and (f) the particle size
distribution histogram.
Fig. 5. SEM of (a) 0.3LLMNO, (b) 0.4LLMNO, (c) 0.5LLMNO, (d) 0.6LLMNO and
(e) 0.7LLMNO hollow-microspheres.
Fig. 6. The high-magnification SEM images and elemental mappings of (a)
0.3LLMNO, (b) 0.4LLMNO, (c) 0.5LLMNO, (d) 0.6LLMNO and (e)
0.7LLMNO.
Fig. 7 (a) Nitrogen adsorption−desorption isotherms and (b) pore size distribution of
the pristine xLLMNO.
Fig. 8. (a) Initial charge-discharge curves of xLLMNO at 0.1 C and (b) Columbic
efficiencies of xLLMNO in the 100 cycles.
Table 1 Initial charge-discharge capacity and irreversible capacity of xLLMNO at
0.1C.
Fig. 9. (a) Cycling stability and (b) rate performance of xLLMNO.
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Fig. 10. Cyclic voltammetric curves of xLLMNO at (a) the initial cycle and (b) the
20th cycle.
Fig. 11. Electrochemical impedance spectroscopy of xLLMNO (a) after the 1st cycle
and (b) the 20th cycle.
Table 2 Compare the Rct of xLLMNO after the 1st cycle and the 20th cycle.
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Fig. 1.
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Fig. 2.
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Fig. 3.
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Fig. 4.
29