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j.jallcom.2017.10.179

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Accepted Manuscript
Phase transitions and related electrochemical performances of Li-Rich layered
cathode materials for high-energy lithium ion batteries
Jianqing Zhao, Xiaoxiao Kuai, Xinyu Dong, Haibo Wang, Wei Zhao, Lijun Gao, Ying
Wang, Ruiming Huang
PII:
S0925-8388(17)33615-0
DOI:
10.1016/j.jallcom.2017.10.179
Reference:
JALCOM 43576
To appear in:
Journal of Alloys and Compounds
Received Date: 7 July 2017
Revised Date:
17 October 2017
Accepted Date: 22 October 2017
Please cite this article as: J. Zhao, X. Kuai, X. Dong, H. Wang, W. Zhao, L. Gao, Y. Wang, R.
Huang, Phase transitions and related electrochemical performances of Li-Rich layered cathode
materials for high-energy lithium ion batteries, Journal of Alloys and Compounds (2017), doi: 10.1016/
j.jallcom.2017.10.179.
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Phase Transitions and Related Electrochemical Performances of Li-Rich Layered Cathode
Materials for High-Energy Lithium Ion Batteries
Jianqing Zhao a, b, d, Xiaoxiao Kuai a, b, Xinyu Dong a, b, Haibo Wang a, b, e, Wei Zhao f,
a
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Lijun Gao a, b, *, Ying Wang d,**, Ruiming Huang c, ***
Soochow Institute for Energy and Materials InnovationS, College of Physics, Optoelectronics
and Energy & Collaborative Innovation Center of Suzhou Nano Science and Technology,
Jiangsu Provincial Key Laboratory for Advanced Carbon Materials and Wearable Energy
Technologies, Suzhou 215006, China
c
Department of Chemistry, Rutgers-Newark, The State University of New Jersey, Newark, New
Jersey 07103, United States
d
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b
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Soochow University, Suzhou 215006, China
Department of Mechanical & Industrial Engineering, Louisiana State University, Baton Rouge,
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Louisiana 70803, United States
Institute of Chemical Power Sources, Soochow University, Zhangjiagang 215600, China
f
Shanghai Haiying Machinery Plant, Shanghai 200436, China
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e
*Corresponding Authors:
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Prof. Lijun Gao, Tel: +86-512-65229905; Fax: +86-512-65229905;
E-mail: gaolijun@suda.edu.cn
Prof. Ying Wang, Tel: +1-225-578-8577; Fax: +1-225-578-9162; E-mail: ywang@lsu.edu
Dr. Ruiming Huang, Tel: +1-973-353-1254; Fax: +1-1-973-353-1264;
E-mail: ayson12345@gmail.com
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Abstract: The present work systematically probes and tracks the phase transition of Li-rich
layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2 (marked as LMNCO) by using an ex-situ chemical
activation that is realized through ion-exchange and post-annealing processes, in order to
lithium-ion
batteries.
Ion
exchanges
H+-Li+
of
and
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understand related electrochemical performances of Li-rich cathode materials for advanced
subsequent
TBA+-H+
(TBA:
tetrabutylammonium) in LMNCO are carried out, resulting in its layered-to-spinel phase
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transition after optimal heat treatments. The resultant compound shows a Li4Mn5O12-type spinel
structure. This converted spinel cathode material can deliver discharge capacities higher than 300
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mAh/g at 0.1 C and 200 mAh/g at 1 C (1 C=250 mA/g), respectively, and also exhibits better
cycling stability and rate capability in comparison with pristine layered LMNCO and other
derivatives. This work offers a feasible route to study all changes of morphologies, crystal
structures, chemical compositions, surface areas and related electrochemical lithium storage
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behaviors during phase transitions of Li-rich layered cathode materials, and thus provides
batteries.
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insights on optimizing electrochemical performances for high-energy and high-power lithium ion
Keywords: Chemical activation, ion exchange, phase transition, Li-rich layered cathode material,
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lithium ion battery
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1. Introduction
The rechargeable lithium ion batteries have been demonstrated as highly effective power
supplies for electric transportation system and portable electronic devices. Performances of
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lithium-ion batteries crucially rely on energy and power densities of electrode materials [1].
Recently, tremendous research efforts focus on developing advanced cathode materials, which
are expected to offer high specific capacity and operating voltage together with outstanding
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cycling stability and rate capability [2],[3],[4]. The Mn-based Li-rich layered oxides have
attracted tremendous research efforts owing to the high lithium storage capacity and working
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potential. These cathode materials marked as Li[LixMnyMz]O2 (M=Co and Ni; x+y+z=1 and
y>0.5) can be cycled over a broad voltage range of 2.0 - 4.8 V vs. Li+/Li and deliver specific
capacities higher than 250 mAh/g, along with other merits including low cost, environmental
friendliness and safety [5-18].
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Li[Li0.2Mn0.54Ni0.13Co0.13]O2 (marked as LMNCO) belongs to aforementioned Li-rich and Mnrich category, which has the desirable theoretical capacity (>300 mAh/g) and high working
voltage (~ 4.0 V vs. Li+/Li) [18]. As reported in literatures [6-8], Li-rich layered LMNCO is the
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product with a structural intergrowth of layered lithium-inactive Li2MnO3 (space group C2/m)
and layered lithium-active LiMn1/3Ni1/3Co1/3O2 (space group R-3m) at a molar ratio of 1:1
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(0.5Li2MnO3·0.5LiMn1/3Ni1/3Co1/3O2). The high capacity of LMNCO cathode material can be
achieved through the electrochemical activation of Li2MnO3 component in the first charge
reaction above 4.5 V vs. Li+/Li. However, such a reaction leads to significantly irreversible
capacity loss and low Coulombic efficiency in the first charge/discharge cycle, and further can
trigger a detrimental layered-to-spinel phase transition during next electrochemical cycling of
activated layered LMNCO. The structural similarity of cubic close-packed oxygen arrays in
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layered and spinel configurations facilitates the layered-to-spinel phase transition in principle
[7],[9],[13],[19],[20]. As a result, transition metal ions migrate to lithium layers in LMNCO and
reside on vacant lithium ion sites permanently. This irreversible phase transformation is
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continued until a hybrid layered-spinel composite material is formed. Accordingly, the voltage
plateau of working cathode is reduced from ~4.0 V (layered) to ~3.0 V (spinel) vs. Li+/Li
[8],[21]. Overall, electrochemical activation of Li2MnO3 component causes structural instability
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and unfavorable phase transition, which accounts for the voltage fading and decreased energy
density of Li-rich layered cathode materials [9].
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On the other hand, it has been reported that the layered-to-spinel phase transformation in Lirich layered cathode materials demonstrates the unexpectedly high-rate capability of hybrid
layered and spinel cathode materials [7],[8],[22], because the spinel phase is lithium-active with
enhanced electronic conductivity and lithium ion diffusivity [7]. The formation of the spinel
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phase in Li-rich layered cathode materials has been demonstrated by high resolution STEM
observations [21], high resolution TEM images with selected area electron diffraction (SAED)
patterns [7], in-situ X-ray diffraction patterns[13], X-ray absorption spectroscopic and Raman
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studies [15],[23], and is also reflected on charge/discharge curves [22], differential capacity plots
[14], and cyclic voltammetric (CV) profiles [7],[24]. The intergrowth of spinel-layered phases
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tends to alleviate the electrochemical inferiority of Li-rich layered cathode materials. However,
structural details of these phases in the cycled electrodes have not been comprehensively
understood yet. The crystal phase of the spinel formed in Li-rich layered oxides is always
reported as either “spinel” or “spinel-like” phase. The effects of such a phase transformation
(whether to improve or deteriorate performances of spinel-layered composite cathodes) are still
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under debate. It is important to explore these effects, in order to understand the fundamental
electrochemical behavior of high-capacity Li-rich layered cathode materials.
As reported in literatures [16, 25], Li-rich layered cathode materials can be chemically
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activated via the protonation (H+-Li+ exchange) in an acidic environment, followed by removing
H+ ions in a post-annealing treatment in air. As a result, the activated cathode material show
considerably increased capacity and associated Coulombic efficiency in the initial cycle.
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Although H+-Li+ ion exchange can induce the formation of the spinel phase, the formation of
spinel domains in the structure of Li-rich layered cathode material is very limited. Consequently,
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the phase transition successively takes place within the hybrid spinel-layered composite cathode
during next electrochemical cycles, leading to structural instability and poor cycling stability
[25]. As referred to other ion-exchange reports [26],[27],[28],[29], alkylammonium hydroxides,
such as tetrabutylammonium hydroxide (TBA+·OH-) and tetramethylammonium hydroxide
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(TMA+·OH-), have been widely employed to exfoliate protonated layered materials into twodimensional (2D) nanosheets. Due to the organic characteristics and larger molecular size of
alkylammonium cations in comparison with protons and lithium ions, TBA+·OH- can be utilized
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for the second ion exchange (TBA+-H+) of protonated Li-rich layered cathode materials, in order
to realize a complete layered-to-spinel phase conversion.
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In our previous work [30], we realized the layered-to-spinel phase transformation of Li-rich
Li[Li0.2Mn0.54Ni0.13Co0.13]O2 by employing ex-situ ion-exchange and post-annealing processes,
and found that the completely-converted material shows a Li4Mn5O12-type spinel structure rather
than commonly-reported LiMn2O4-type spinel. The approach we developed not only allows the
comprehensive study of electrochemical effects resulting from the growth of a spinel phase
within Li-rich layered cathode materials, but also offers a feasible route to precisely identify
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different crystal structures during the formation of newly-formed spinel phase. Herein, we report
a more comprehensive study during the phase transition of Li-rich layered materials with aspects
to all changes of morphologies, crystal structures, chemical compositions, surface areas and
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related electrochemical lithium storage behaviors in details. Electrochemical performances
indicate that introduction of a spinel phase significantly increases the specific capacity of ~100
mAh/g and results in much better high-rate performances as compared with original Li-rich
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layered cathode materials, but reduces the working voltage from 4.0 V to 3.0 V due to the
activation of Mn3+/Mn4+ redox pair. It is also interesting to find that Li+-TBA+ can be carried out
from the converted spinel phase.
2. Experimental
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in the ion-exchanged intermediate material for the possible recovery of the layered structure
2.1 Synthesis of Li-rich layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2 nanoparticles
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Li-rich layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2 nanoparticles were synthesized by using coprecipitation method. Three precursor solutions were simultaneously prepared. 0.08 mol
transition metal precursor at a molar ratio of Mn(CH3COO)2·4H2O : Ni(CH3COO)2·4H2O :
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Co(CH3COO)2·4H2O = 0.54 : 0.13 : 0.13 was dissolved in 50 mL ethanol; the lithium precursor
solution was composed of 0.12 mol LiOH dissolved in 20 mL distilled water; and the surfactant
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solution was 5.4 mmol F127 (EO106PO70EO106) dissolved in 50 mL ethanol. The F127/ethanol
solution and transition metal precursor solution was first mixed together at 40°C under
continuous stirring, and then the lithium precursor solution was dropwise added to precipitate
transition metal ions. The resulting suspension was heated at 80 °C to completely remove the
solvent and then dried in air at 120 °C for 12 h. The dried powder was annealed in air at 300 °C
for 3 h at a temperature ramp of 1 °C/min, followed by sintering at 900 °C for 12 h at a
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temperature ramp of 5 °C/min. Li[Li0.2Mn0.54Ni0.13Co0.13]O2 nanoparticles were obtained after
cooling to room temperature.
2.2 Chemical activation of Li-rich layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2 via ion exchanges and
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post-heat treatments
The chemical activation of Li-rich layered oxide was carried as follows: first, 1 g
Li[Li0.2Mn0.54Ni0.13Co0.13]O2 particles were dispersed in 150 mL 2 M HCl aqueous solution for
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the H+-Li+ ion exchange of Li[Li0.2Mn0.54Ni0.13Co0.13]O2 at ambient temperature. The HCl
solution was replaced every 2 days for 5 times in order to achieve deep protonation. The
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protonated intermediates dispersed in 150 mL aqueous HCl was obtained in a brown suspension.
Secondly, an aqueous tetrabutylammonium (TBA·OH) solution (Sigma Aldrich) with a mass rate
of 20 wt.% was employed to perform TBA+-H+ exchange of protonated particles. The volumetric
ratio of TBA·OH solution over the brown suspension was set to 5 : 1, and these two solution was
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mixed in a vortex stirrer for 30 min. All resulting ion-exchanged particles were collected via
centrifugation and washed with distilled water for several times. Thirdly, the Li+-TBA+ exchange
was carried out in 1 M LiOH aqueous solution to study the reversibility of different ion
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exchanges. Finally, all ion-exchanged derivatives were annealed in air at 500 °C for 3 h at a
temperature ramp of 1 °C/min.
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2.3 Characterizations
Crystallographic structures and phases of Li-rich layered nanoparticles and all derivatives
were analyzed by X-ray diffraction (XRD) on a Panalytical X’pert Diffractometer with Cu Kα
radiation. Morphology and particle characteristics of different samples were examined using a
field emission scanning electron microscopy (FESEM, Hitachi S4800). Detailed structures of
different samples were observed on transmission electron microscopy (TEM, FEI Tecnai G2
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FEG) at an acceleration voltage of 300 kV. Porous structure and specific surface area of powders
was measured by nitrogen adsorption/desorption at 77 K on a Quantachrome AS-1 instrument
using the Brunauer-Emmet-Teller (BET) method. Chemical compositions of specimens were
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determined by inductively coupled plasma atomic emission spectrometry (ICP-AES) on a
SPCTRO CIROS elemental analyzer.
2.4 Electrochemical measurements
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The cathodes were consisted of 80 wt.% Li[Li0.2Mn0.54Ni0.13Co0.13]O2 particles or its
derivatives, 10 wt.% acetylene black as the conductive carbon (Alfa Aesar, 99.5%), and 10 wt.%
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polyvinylidene fluoride (PVDF) as the binder (Alfa Aesar). These electrodes were assembled
into CR2032-type coin cells for electrochemical measurements, with the metallic lithium foil as
the anode and Celgard 2320 membrane as the separator. The electrolyte was 1 M LiPF6 dissolved
in ethylene carbonate (EC), dimethyl carbonate (DMC) and diethyl carbonate (DEC) at a
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volumetric ratio of 1:1:1. Galvanostatic charge and discharge were performed at different current
densities in a voltage range of 2.0 - 4.8 V vs. Li+/Li using an 8-channel battery analyzer (MTI
Corporation). Theoretical capacities of different cathode materials are all set to 250 mAh/g, i.e.,
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current density corresponding to 1 C is 250 mA/g. Cyclic voltammetric (CV) curves of cathodes
were recorded at a scanning rate of 0.1 mV/s between 2.0 and 4.8 V vs. Li+/Li using an
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electrochemical analyzer (CHI 605C).
3. Results and discussion
As shown in Fig. 1, ion-exchange processes and calcinations result in dramatically
morphological changes of different derivatives. The pristine Li-rich Li[Li0.2Mn0.54Ni0.13Co0.13]O2
particles exhibit a distinct aggregation with an even particle size around ~250 nm in Fig. 1a,
while H+-Li+ ion exchange in acidic environment gives rise to distinct layered cake-shaped
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blocks of LHMNCO. We speculate that the multilayered morphology and structure of LHMNCO
(Fig. 1b) is probably attributed to the structural introduction of layered LMNCO. As reported in
literatures [26, 29, 31], the protonation is required for the subsequent H+-TBA+ ion exchange.
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After shaking the intermediate LHMNCO and HCl mixture (i.e., protonated LHMNCO particles
dispersed in the final HCl solution) via violent vortexes in the aqueous TBA·OH solution at a
volumetric ratio of 1:5, the collected LHMNCO TBA shows the interesting morphology of
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nanoflowers (Fig. 1c). Each particle is composed of numerous ultrathin nano-petals. This
phenomenal morphology change is mostly resulted from the cooperative effects of the TBA-
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assisted exfoliation and the turbulence-induced reaction environment. It is interesting that further
Li+-TBA+ ion exchange in a basic solution extensively unfolds petals of LHMNCO TBA
nanoflowers into nanosheet stacks of LHMNCO TBA Li as shown in Fig. 1d. Fig. 1e-g reveal
morphologies of ion-exchanged derivatives after post-annealing treatments, which apparently
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cause aggregations and coarse structures of LHMNCO HT, LHMNCO TBA HT and LHMNCO
TBA Li HT, respectively, after the removal of H+ and TBA+ ions along with other byproducts,
such as H3O+ and OH-.
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Accordingly, phase transitions and structural reconstructions accompanying with morphology
changes from initial layered LMNCO to different converted derivatives have been studied from
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XRD and TEM characterizations in Fig. 2 and Fig. 3, respectively. In agreement with reported
XRD patterns of Li-rich layered materials [7, 8, 17, 18], pristine LMNCO in Fig. 2a shows
typical XRD peaks that are indicative of the intergrowth of monoclinic Li2MnO3 (space group:
C2/m) and rhombohedral LiMn1/3Ni1/3Co1/3O2 (space group: R-3m) in the layered structure [7].
The main layered structure of LiMn1/3Ni1/3Co1/3O2 can be determined from distinct peak splits of
two (006)L2-(012)L2 and (108)L2-(110)L2 doublets at 2θ = 36 - 38º and 2θ = 64 - 66º, respectively.
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The weak (020)L1 reflection at 2θ = 20 - 23º is belong to the layered Li2MnO3 phase, which is the
superslattice within the parent layered structure [17]. The intergrowth of these two layered
components can be further confirmed from the other five peaks at 2θ = 18.7º, 36.9º, 37.9º, 38.5º
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and 44.5º, resulting in diffraction patterns of (001)L1/(003)L2, (200)L1/(101)L2, (113)L1/(006)L2,
(131)L1/(012)L2, and (202)L1/(104)L2. A spinel-like impurity is found in the XRD pattern of
LHMNCO (marked with red asterisks in Fig. 2b), which is accordant with that reported in the
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literatures [7, 22]. As shown in Fig. 2b, the peak merger of (113)L1/(006)L2 and (131)L1/(012)L2
pairs reveals the distortion of layered structure at a certain degree, while the preserved peak splits
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of (108)L2-(110)L2 doublets indicate the retained layered structure in the protonated LHMNCO
intermediate [22]. Furthermore, XRD pattern of LHMNCO TBA powder manifests the growth of
spinel-like phase during TBA+-H+ exchange, since the intensity of one representative XRD peak
at 2θ = 19.3º in LHMNCO for the spinel phase increases. The cubic close packed oxygen arrays
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both in layered and spinel structures is essential to realize the phase transition in Li-rich layered
transition metal oxides by migrating transition metal ions into lithium layers when lithium ion
vacancies exist during ion-exchange processes [13],[20]. LHMNCO TBA Li shows identical
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XRD pattern to that of LHMNCO TBA, indicating limited reversibility of phase transition by
Li+-TBA+ ion exchange. However, the peak splitting of (108)L2-(110)L2 doublets in its XRD
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pattern reveals the prominent layered structure of LHMNCO TBA Li derivative, despite the
spinel phase has been detected. Post-annealing treatment has been demonstrated as an effective
way to remove H+ and TBA+ substituents in air, resulting in the generation of corresponding Li+
vacancies in lithium layers [32]. Such facile process can accelerate the diffusion of transition
metal ions into lithium ion sites, and thus promote the layered-to-spinel phase transition. The
enlarged selected 2θ portions in Fig. 2b at 2θ =16-20º, 34-40º and 62-68º illustrate the phase
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transitions when different ion-exchanged samples are subjected to calcinations in air. The merge
of two separate peaks around 2θ = 19º into one peak occurs for LHMNCO, LHMNCO TBA and
LHMNCO TBA Li after heat treatments, respectively, indicating dramatic phase transitions due
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to the removal of foreign H+ and TBA+ cations. It is surprising to find that the coupled (108)L2(110)L2 pair of LHMNCO TBA has merged to one broad peak for LHMNCO TBA HT at a lower
2θ position, while LHMNCO HT and LHMNCO TBA Li HT still show distinguishing peak
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splits of (108)L2-(110)L2 between 2θ = 62º and 64º. As mentioned before, the peak split of
(108)L2-(110)L2 doublets is characteristic of layered structure, which is distinguished from the
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spinel phase showing the (440)S reflection at the same position. As a result, XRD pattern of
LHMNCO TBA HT can be indexed to the spinel Li4Mn5O12 phase with a Fd-3m space group,
indicating the complete layered-to-spinel phase transition from original Li-rich layered LMNCO
to Li4Mn5O12-type spinel compound after two-step ion exchanges, followed by a post-annealing
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process. In contract, XRD patterns of LHMNCO HT and LHMNCO TBA Li HT both reveal the
coexistence of layered and spinel phases (Fig. 2). It is suggested that a second TBA+-H+ ion
exchange is crucial to realize a complete phase conversion. We speculate that due to the larger
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size of TBA+ cations than protons, TBA+ substituents can increase c-axis of ion-exchanged
layered derivative in comparison with the effect from H+ ions, which will increase the structure
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instability that significantly facilitate migrations of transition metal ions when TBA+ are burned
away in air. On the contrary, Li+ can be partially restored in the lithium layers through the Li+TBA+ ion exchange, resulting in hybrid layered-spinel structure of LHMNCO TBA Li HT.
Fig. 3 shows the structural evolution from the pristine layered LMNCO to the converted
Li4Mn5O12-type LHMNCO TBA HT spinel in TEM and HRTEM observations. In consistence
with SEM image in Fig. 1a, LMNCO nanoparticles have the solid structure with an average
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particle size of ~250 nm. The lattice fringe as shown in Fig. 3b indicates the high crystallinity of
pristine Li-rich layered materials, due to the high synthetic heating temperature at 900 ºC. The
characteristics of overlapped sheets can be observed in Fig. 3c of LHMNCO after H+-Li+ ion
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exchange in an acidic HCl solution. The formation of spinel phase during this process at the
surface of particle is identified in HRTEM image (Fig. 3d), which is consistent with XRD results
in Fig. 2. The continuous TBA+-H+ ion exchange not only tailors LHMNCO TBA to the
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nanoflower-like shape, but also generates porous structure as shown in Fig. 3e. Furthermore, due
to the violent exfoliation effect of TBA+ cations, HRTEM image in Fig. 3f shows disordered
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lattice fringes both at the surface and in the bulk of LHMNCO TBA; the other reason leading to
such a disordered structure possibly results from the partial decomposition of organic TBA+
cations under the attack of high-energy electron beam during HRTEM observations. On the other
hand, fringes with the smaller d-space (d=0.455 and 0.438 nm in Fig. 3f) at the surface of the
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specimen may also be attributed to the partial decomposition of TBA+ substituents. As shown in
Fig. 1c and 1f, monodispersive LHMNCO TBA nanoflowers convert to LHMNCO TBA HT
particles with irregular shapes after post-annealing processes due to the fold of nanopetals.
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Accordingly, the removal of TBA+ cations also contributes to the porous structure of LHMNCO
TBA HT as shown in TEM image (Fig. 3g). According to the XRD results as shown in Fig. 2,
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HRTEM image in Fig. 3h shows the spinel crystal structure of LHMNCO TBA HT with a dspace of (111)S equal to 0.471 nm. In general, ex-situ ion-exchange and heat treatments result in
the complete phase transition from the layered LMNCO to a Li4Mn5O12-type spinel material,
along with intriguing morphological and structural evolutions.
The nitrogen adsorption and desorption isotherms and pore size distributions of pristine Lirich layered LMNCO, ion-exchanged LHMNCO TBA and annealed LHMNCO TBA HT are
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shown in Fig. 4a and 4b, respectively. The corresponding porous characteristics in terms of
surface area, pore volume and relative pore size are summarized in Table 1 together with
Li/Mn/Ni/Co ratios of three samples. It is clear that substitution of Li+ ions within LMNCO by
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H+ protons, followed by continuous replacement with TBA+ cations results in significantly
increased surface area to 11.109 m2/g of LHMNCO TBA, almost four times higher than that of
original LMNCO particles (2.327 m2/g), which can be attributed to the exfoliation effect from
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TBA+ cations as shown in SEM (Fig. 1c) and TEM (Fig. 3e) images [27,31]. Accordingly,
LHMNCO TBA also has a larger pore volume of 8.880e-2 cm3/g in comparison with 1.133e-2
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cm3/g from LMNCO nanoparticles. Both SEM and TEM observations indicate that the pore
volume of LMNCO powder is from special gaps between numerous agglomerated LMNCO
nanoparticles (Fig. 1a and Fig. 3a), while the higher pore volume of LHMNCO TBA mostly
arises from the porous structure of individual LHMNCO TBA nanoflowers (Fig. 1c and Fig. 3e).
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Therefore, agglomerated LMNCO nanoparticles give rise to a relatively higher pore size
distribution of ~6 nm in Fig. 4b as compared with ~4 nm from LHMNCO TBA nanoflowers with
monodisperse characteristic. As aforementioned in XRD characterizations in Fig. 2, the post-
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annealing treatment plays a crucial role in realizing a complete layered-to-spinel phase transition,
and the morphological and structural changes have been observed in SEM (Fig. 1f) and TEM
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(Fig. 3g) images, respectively. As a result, heating LHMNCO TBA in air at 500 °C contributes
to further increased surface area to 13.725 m2/g of LHMNCO TBA HT and two pore size
distributions of ~3 and ~12 nm in Fig. 4d. Its reduced pore volume is probably due to the folded
nanopetals to form an internal porous structure (Fig. 3g) and the obvious aggregation (Fig. 1f)
after post-heat treatment. The chemical compositions of these three samples are compared in
Table 1 in the form of Li/Mn/Ni/Co molar ratios. In comparison with the theoretical ratio of
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Li/Mn/Ni/Co=1.2/0.54/0.13/0.13, the as-prepared LMNCO shows slightly less quantity of the
lithium component that is probably assigned to the lithium loss during heat treatment at high
temperature at 900°C together with the long duration time for 12 h. Such the harsh heat treatment
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is required for the synthesis of Li-rich LMNCO material, in order to achieve its nice integrated
structure and high crystallinity, but leads to unfavorable lithium loss as measured by ICP results.
After two step ion exchanges via H+-Li+ and TBA+-H+, LHMNCO TBA preserves 52.7 % of the
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original lithium ions in LMNCO. As reported in the literature [29], layered transition metal
oxides can be fully exchanged and teared into transition metal oxide nanosheets. In our case, the
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partial ion exchange possibly is attributed to the formation of spinel phase at the surface of
layered derivatives as shown in XRD results (Fig. 2) and HRTEM image (Fig. 3d). The
detectable Ni loss in LHMNCO TBA after the ion-exchange processes may result from the
cationic Li+-Ni2+ disorder in LMNCO. Consequently, a few Ni2+ ions occupied in lithium sites in
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the lithium layer are replaced with H+ and TBA+ cations. The post-annealing process has few
effect on the chemical composition of LHMNCO TBA HT,
resulting in
Li/M
(M=Mn+Ni+Co)=0.75. Such value is much close to the Li/Mn ratio of 0.8 in Li4Mn5O12 spinel
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rather than 0.5 of LiMn2O4 spinel, which is consistent with the XRD result being indexed to
Li4Mn5O12-type spinel for LHMNCO TBA HT. In short summary, ex-situ ion-exchange
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processes along with post-heat treatments offer a feasible approach not only to tailor morphology
and structure of Li-rich layered transition metal oxides, but also to control the phase
transformation between layered and spinel phases. We speculate that the nanoarchitectured
LHMNCO TBA HT spinel material should be favorable to facilely accommodate the electrolyte,
maximize electrochemical active sites and release reaction strain during lithiation and
delithiation; hence, LHMNCO TBA HT cathode with a pure spinel phase and porous structure is
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expected to show enhanced rate capability and cycleability as compared with the pristine Li-rich
layered LMNCO as well as other two hybrid LHMNCO HT and LHMNCO TBA Li HT with an
intergrowth of layered and spinel phases.
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The cyclic voltammetric (CV) measurements of LMNCO, LHMNCO HT, LHMNCO TBA
HT and LHMNCO TBA Li HT cathode materials are carried out, in order to study
electrochemical properties related to phase transitions from Li-rich layered (LMNCO) to either
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layered-spinel (LHMNCO HT and LHMNCO TBA Li) or Li4Mn5O12-type spinel (LHMNCO
TBA HT) phase. Fig. 5 shows resulting CV records in the first three cycles of four different
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samples. LMNCO reveals the typical electrochemical characteristics of Li-rich layered cathode
materials as shown in Fig. 5a. The first anodic peak at 4.17 V in the initial charge curve is
associated with the oxidation of Ni2+ to Ni4+, followed by Co3+ to Co4+, whereas Mn still remains
as tetravalent in LiMn1/3Ni1/3Co1/3O2 component [33]. The second anodic peak at 4.66 V
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corresponds to the electrochemical activation of inert Li2MnO3 component, i.e., the
decomposition of Li2MnO3 to Li2O and lithium-active MnO2, along with the unavoidable
decomposition of electrolyte and the formation of solid electrolyte interphase (SEI) at such a
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high potential >4.5 V [10]. Although the electrochemical activation process leads to the low
Coulombic efficiency in the first cycle, but significantly results in high capacity of Li-excess
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layered cathode materials in the successive cycles. Correspondingly, the reduction of Co4+/Co3+
and Ni4+/Ni3+/Ni2+ redox occurs at 3.66 V in the initial discharge curve. In the second cycle, the
anodic peak at 4.66 V disappears, while an additional cathodic peak at 3.26 V appears, which can
be attributed to the reduction of Mn4+ to Mn3+ from the as-activated MnO2 component. The third
cycle shows the similar profile to the second cycle in less polarization and higher current density,
indicating improved electrochemical reversibility after the electrochemical activation of Li-rich
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layered LMNCO. It is clear to see that CV curves of LHMNCO HT, LHMNCO TBA HT and
LHMNCO TBA Li HT are very similar to each other, but apparently different from that of
LMNCO. Those three materials all show the dominant redox pair around 3.0 V in CV curves,
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together with two minor redox couples located near 4.0 and 4.6 V. The CV performance is
consistent with the typical electrochemical characteristics of reported Ni/Co-doped Li4Mn5O12type spinel in a wide voltage range [34]. The CV responses of LHMNCO HT, LHMNCO TBA
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HT and LHMNCO TBA Li HT support XRD results in Fig. 2, which reveal the Li4Mn5O12-type
spinel structure of newly-formed spinel phase within the original layered structure of LMNCO
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after ion-exchange and post-annealing processes. The appearance of an anodic peak at ~4.6 V in
the initial CV charges of LHMNCO HT and LHMNCO TBA Li HT indicate the existence of
preserved layered structure, in accordance with XRD characterizations (Fig. 2). In contract,
LHMNCO TBA HT reveals much lower current density of such anodic peak in the first CV
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charge. Furthermore, an more intensive anodic peak at 2.95 V is generated, which can be
attributed to the better complete phase transition of LHMNCO TBA HT in comparison with
LHMNCO HT and LHMNCO TBA HT. LHMNCO TBA HT shows identical CV curves of
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second and third cycles, indicating outstanding electrochemical reversibility of this spinel
materials. The anodic peak at 3.07 V and cathodic peak at 2.63 V are attributed to corresponding
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oxidation and reduction reactions of Mn3+/Mn4+ redox pair, which are associated with extracting
and inserting lithium ions on 16c sites in the spinel structure [1]. Two small redox couples at
~4.0 and ~4.6 V probably result from the Co3+/Co4+ and Ni2+/Ni3+/Ni4+redox, respectively.
Fig. 6 exhibits charge and discharge curves of four cathodes in the first five cycles at 0.1 C in
a voltage range of 2.0-4.8 V vs. Li+/Li, which are in well accordance with CV profiles in Fig. 5.
It is noticeable that introduction of a spinel phase with Li-rich layered cathode materials can
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significantly increase the specific capacity. As shown in Fig. 6a-6d, charge/discharge curves of
the fourth and fifth cycle are almost identical for all cathodes, suggesting that the cathodes are
mostly stable after five electrochemical cycles. The pristine Li-rich layered LMNCO delivers a
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specific discharge capacity of 211.3 mAh/g at the fifth cycle with a voltage plateau around 3.7 V,
revealing typical electrochemical behavior of Li-rich layered cathode materials. The XRD result
in Fig. 2a demonstrates the coexistence of layered and spinel phases of LHMNCO HT after ex-
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situ ion-exchange and post-heat treatments, in which the spinel phase is dominant; hence
LHMNCO HT shows a higher discharge capacity of 285. 8 mAh/g but along with a predominant
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voltage plateau at 2.6 V as well as a minor one at 4.4 V. The former voltage stage is due to the
active Mn3+/Mn4+ redox in the spinel component, and the latter probably results from the
Co4+/Co3+ and Ni4+/Ni3+/Ni2+ redox reactions in the reversed layered component. Furthermore,
employing TBA+ cations for the second ion exchange of LHMNCO have contributed to the
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complete phase conversion of Li-rich layered LMNCO, resulting in the spinel LHMNCO TBA
HT with the Li4Mn5O12-type spinel characteristics (Fig. 2) and a mesoporous structure (Fig. 3g
and Fig. 4b). As a result, an unexpectedly high discharge capacity of 343.2 mAh/g is achieved in
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the fifth cycle of LHMNCO TBA HT. There are three voltage plateaus located at 4.6, 4.0 and
2.8 V, respectively, which is distinctly different from the profile of LHMNCO HT, again
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revealing different structural characteristics between LHMNCO TBA HT and LHMNCO HT.
Those three voltage stages can be attributed to reductions from Ni4+/Ni3+/Ni2+, Co4+/Co3+ and
Mn4+/Mn3+ redox pairs in the spinel structure and are favorable to preserve high-voltage
performances of LHMNCO TBA HT. Fig. 3g reveals a porous mesoporous structure of
nanoarchitectured spinel cathode, which would be favorable to accommodate electrolyte and
effectively release reaction strains during lithiation/delithiation. Such a structure may enable to
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absorb lithium ions within the porous structure, resulting in the additional capacity contribution
[35, 36]. On the other hand, the porous structure may partially collapse when adsorbed lithium
ions are extracted during the charge processes, leading to reduced lithium storage capacity in
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corresponding discharge processes. Accordingly, the LHMNCO TBA HT cathode (Fig. 6c)
shows relatively low but gradually increased Coulombic efficiencies in initial cycles as
compared with the other three cathode materials. In contrast, LHMNCO TBA Li HT shows
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similar electrochemical performance to LHMNCO HT with the intergrowth of layered and spinel
structures, delivering a reduced capacity of 259.9mAh/g. This might result from the partially
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recovered layered structure via Li+-TBA+ ion exchange of LHMNCO TBA in LiOH aqueous
solution.
Fig. 7a and 7b show cycling and high-rate performances of LHMNCO HT, LHMNCO TBA
HT and LHMNCO TBA Li HT cathodes in comparison with the pristine Li-rich layered
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LMNCO, respectively. The effects from the introduced spinel phase within the layered LMNCO
cathode material on improving specific capacity, cycling stability and rate capability are more
phenomenal when cycled at higher current densities. As shown in Fig. 7a, LHMNCO TBA HT
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can retain a very high discharge capacity of 197.5 mAh/g with a corresponding capacity
retention of 89.1% after 100 electrochemical cycles at 1C, much better than 58.1 mAh/g and 65.9%
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of pristine layered LMNCO, 116.1 mAh/g and 85.1% of LHMNCO HT, and 77.9 mAh/g and
80.0% of LHMNCO TBA Li HT. Moreover, LHMNCO TBA HT delivers initial capacities of
313.6, 267.2, 203.9, 180.7, 126.3, and 89.4 mAh/g at 0.1, 0.5, 1, 2, 5, and 10 C, respectively, as
exhibited in Fig. 7b. Such the remarkable cyclability and high-rate capability of LHMNCO TBA
HT can be attributed to reconstructed spinel phase and hierarchical mesoporous structure for
facile accommodation and diffusion of lithium ions, and effectively releasing reaction strains in
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the “buffer” structure in high porous characteristics. Overall, generation of a spinel phase within
Li-rich layered cathode materials can considerably increase the specific capacity and rate
capability, but has to sacrifice the working voltage. Doping transition metal cations, such as Ni2+,
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Co3+ and Fe3+ ions, can contribute to the high-voltage performance. Ion-exchange method offers
a desirable way to obtain enhanced electrochemical performance of Li-rich layered cathode
materials.
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4. Conclusions
This work sheds light on fundamental understanding of layered-to-spinel phase transition and
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relevant electrochemical performances of Li-rich layered cathode materials via ex-situ ionexchange processes, followed by post-annealing treatments. Employing TBA+ cations for the
second ion exchange of pronated Li-rich layered oxides is critical to realize a complete phase
transition, resulting in a Li4Mn5O12-type spinel-structured material converted from Li-rich
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layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2. Ion-exchange procedure also tailor the morphology and
structure of solid Li[Li0.2Mn0.54Ni0.13Co0.13]O2 nanoparticles into nanostructured spinel material
with high surface area and mesoporous porosity. In comparison with the pristine Li-rich layered
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cathode material, the final converted spinel cathode material with hierarchical porous structure
reveals significantly increased specific capacity, better cycling stability and rate capability. This
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work opens up a route to maximize electrochemical performance of Li-excess layered cathode
materials for high-power and high-energy lithium ion batteries.
Acknowledgements
This work was supported by US National Science Foundation, the Division of Chemical,
Bioengineering, Environmental and Transport Systems (NSF CBET) [grant number 1438493];
the US Small Business Technology Transfer (STTR) [grant number 1346496]; the Research
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Enhancement Award (REA) program, Louisiana Space Consortium (LaSPACE) funded via the
NASA Space Grant College & Fellowship Program Grant 2011-15 Cycle [grant number
NASA/LEQSF(2010-2015)-LaSPACE]; the National Natural Science Foundation of China
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[grant number U1401248]; the Natural Science Foundation of Jiangsu Province, China [grant
number BK20151227]; the General Financial Grant from the China Postdoctoral Science
Foundation [grant number 2016M601876]. The authors also acknowledge LSU IAM Shared
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Instrumentation Facility (SIF) at Louisiana State University and Suzhou Key Laboratory for
Advanced Carbon Materials and Wearable Energy Technologies, Suzhou 215006, China for
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using facilities, and thank Prof. Kerry Dooley at Department of Chemical Engineering at
Louisiana State University for BET measurements, and Dr. Jibao He for TEM observations at
Tulane University.
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Figure Captions
Fig. 1 SEM images of (a) LMNCO, (b) LHMNCO, (c) LHMNCO TBA, (d) LHMNCO TBA Li,
(e) LHMNCO HT, (f) LHMNCO TBA HT and (g) LHMNCO TBA Li HT.
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Fig. 2 XRD patterns of pristine Li-rich layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2 nanoparticles and
corresponding derivatives in (a) full 2θ range and (b) enlarged 2θ portions between 16-22º, 3440º and 62-68º, respectively.
LHMNCO TBA, and (g and h) LHMNCO TBA HT.
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Fig. 3 TEM and HRTEM images of (a and b) LMNCO, (c and d) LHMNCO, (e and f)
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Fig. 4 (a) Nitrogen adsorption/desorption isotherms and (b) corresponding pore size distributions
of LMNCO and LHMNCO TBA and LHMNCO TBA HT.
Fig. 5 Cyclic voltammetric (CV) curves of (a) LMNCO, (b) LHMNCO HT, (c) LHMNCO TBA
HT and (d) LHMNCO TBA Li HT in the first three cycles at a scanning rate of 0.1 mV/s in a
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voltage range of 2.0-4.8 V vs. Li+/Li.
Fig. 6 Charge and discharge curves of (a) LMNCO, (b) LHMNCO HT, (c) LHMNCO TBA HT
and (d) LHMNCO TBA Li HT in the first five cycles at a current density of 0.1 C in a voltage
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range of 2.0-4.8 V vs. Li+/Li.
Fig. 7 (a) Cycling performances at 1 C and (d) high-rate performances at different current
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densities of LHMNCO HT, LHMNCO TBA HT, LHMNCO TBA Li HT in comparison with the
pristine layered LMNCO in a voltage range of 2.0-4.8 V vs. Li+/Li.
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Fig. 1 SEM images of (a) LMNCO, (b) LHMNCO, (c) LHMNCO TBA, (d) LHMNCO TBA Li,
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(e) LHMNCO HT, (f) LHMNCO TBA HT and (g) LHMNCO TBA Li HT.
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Fig. 2 XRD patterns of pristine Li-rich layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2 nanoparticles and
corresponding derivatives in (a) full 2θ range and (b) enlarged 2θ portions between 16-22º, 34-
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40º and 62-68º, respectively.
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Figure 3. TEM and HRTEM images of (a and b) LMNCO, (c and d) LHMNCO, (e and f)
LHMNCO TBA, and (g and h) LHMNCO TBA HT.
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Fig. 4 (a) Nitrogen adsorption/desorption isotherms and (b) corresponding pore size distributions
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of LMNCO and LHMNCO TBA and LHMNCO TBA HT.
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Fig. 5 Cyclic voltammetric (CV) curves of (a) LMNCO, (b) LHMNCO HT, (c) LHMNCO TBA
HT and (d) LHMNCO TBA Li HT in the first three cycles at a scanning rate of 0.1 mV/s in a
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voltage range of 2.0-4.8 V vs. Li+/Li.
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Fig. 6 Charge and discharge curves of (a) LMNCO, (b) LHMNCO HT, (c) LHMNCO TBA HT
and (d) LHMNCO TBA Li HT in the first five cycles at a current density of 0.1 C in a voltage
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range of 2.0-4.8 V vs. Li+/Li.
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.
Fig. 7 (a) Cycling performances at 1 C and (d) high-rate performances at different current
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densities of LHMNCO HT, LHMNCO TBA HT, LHMNCO TBA Li HT in comparison with the
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pristine layered LMNCO in a voltage range of 2.0-4.8 V vs. Li+/Li.
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Table 1. Porous characteristics and elemental composition of LMNCO, LHMNCO TBA and
LHMNCO TBA HT.
ICP elemental compositions
Samples
Pore volume
(cm3/g)
Pore size
(nm)
Li
2.327
1.133e-02
~6
1.111
11.109
8.880e-02
~4
13.725
5.921e-02
~3 & ~12
Mn
Ni
Co
0.540
0.129
0.128
SC
Surface area
(m2/g)
0.586
0.540
0.119
0.126
0.585
0.540
0.118
0.126
M
AN
U
AC
C
EP
TE
D
Pristine
LMNCO
Ion-exchanged
LHMNCO TBA
Annealed
LHMNCO TBA
HT
RI
PT
Porous characteristics
33
ACCEPTED MANUSCRIPT
Highlights
Ion exchanges of H+-Li+ and TBA+-H+ are performed on Li-rich layered oxides.
The converted spinel phase has a Li4Mn5O12-type spinel structure.
RI
PT
Layered-to-spinel phase transition of Li-rich layered oxides has been realized.
Electrochemical performances of Li-rich layered cathodes are manipulated.
AC
C
EP
TE
D
M
AN
U
SC
The resulting spinel cathode delivers a capacity higher than 300 mAh/g at 0.1 C.
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