Solution-based combinatorial samples of lithium manganese nickel oxide were synthesized and studied by X-ray diffraction in order to map out the entire pseudoternary system. The samples were made by heating to 800 °C in an oxygen atmosphere, and two cooling methods were tested: quenching and a slower cooling rate. This article focuses on the areas of the Gibbs triangle that lie between the layered and spinel structures as well as the single-phase layered region. The layered region is much larger than previously reported in the literature, and lattice parameters throughout the solid solution are reported. The complex coexistence region includes a two-phase region as well as two 3-phase regions. Supporting evidence demonstrating the 3-phase regions and careful identification of their corners is included here. The 3-phase regions also undergo a transformation during slow cooling that can be attributed to high entropy phases no longer being favored. In order to confirm that the findings for the small combinatorial samples are relevant to bulk samples, we present XRD patterns from a few bulk samples synthesized in oxygen as well as some heated in air. The results show that though the boundaries in the phase diagram move with synthesis conditions, the main features of the coexistence regions remain the same.

Lithium-rich Li[Li x M 1−x ]O 2 (M = Ni, Mn, Co) materials have been claimed to be two phase by some researchers and to be one phase by others when all the available lithium is extracted electrochemically. To clear up this confusion, the Li-rich samples [Li[Li 0.12 (Ni 0.5 Mn 0.5 ) 0.88 ]O 2 and Li-[Li 0.23 (Ni 0.2 Mn 0.8 ) 0.77 ]O 2 with different particle sizes were synthesized for in situ X-ray diffraction experiments. In situ X-ray diffraction measurements revealed two-phase behavior of 10 μm particles and one-phase behavior for samples with submicrometer particles. The phase separation in samples with large particles agrees with literature proposals of oxygen release from a surface layer and the observation of distinct surface and bulk phases. The small particle samples are so small that they are entirely composed of the surface phase found in the large particle samples. These results strongly suggest that the size of particles can significantly affect the structural evolution testing and electrochemical performance of the Li-and Mn-rich materials. It is proposed that the surface phase continuously grows during charge−discharge cycling, which leads to voltage fade in large particle samples. Meanwhile, in situ X-ray diffraction measurements were also performed for the layered Li−Ni−Mn−Co oxides with varying nickel contents, including NMC811 (LiNi 0.8 Mn 0.1 Co 0.1 O 2 ), NMC442 (LiNi 0.42 Mn 0.42 Co 0.16 O 2 ), [Li[Li 0.12 (Ni 0.5 Mn 0.5 ) 0.88 ]O 2 , and Li[Li 0.23 (Ni 0.2 Mn 0.8 ) 0.77 ]O 2 . Samples with higher nickel content showed much faster contraction of unit cell volume as a function of cell voltage, which suggests that the core−shell structures with a nickel-rich core (e.g., NMC811) and a Mn-rich shell (e.g., Li 1.23 Ni 0.154 Mn 0.616 O 2 ) should not crack during charge−discharge cycling.

The capacity fade mechanisms of LiCoO 2 /Si-alloy:graphite pouch cells filled with a 1M LiPF 6 EC:EMC:FEC (27:63:10) electrolyte were studied using galvanostatic cycling, electrochemical impedance spectroscopy on symmetric cells, gas-chromatography and differential voltage analysis. Analysis of the gas generated during the first cycle indicated that FEC reacts at the negative electrode following a 1-electron reduction pathway and other pathways that do not lead to the formation of gaseous products. An analysis of the electrolyte showed that FEC is continuously consumed during the first 80% of the first charge (formation cycle). Typical cells charged and discharged at 40 • C showed a gradual capacity loss for the first 250 cycles followed by a sudden capacity drop associated with a large polarization growth. Analysis of the electrolyte showed that this sudden failure is associated with the depletion of FEC. The capacity loss as well as the consumption of FEC prior to the sudden failure was fitted using a model that includes a time dependence and a cycle number dependence. The time dependence was associated with the thickening solid electrolyte interface at the surface of the negative electrode particles (Si-alloy and graphite) and the cycle number dependence was associated with the solid electrolyte interface repair at the surface of Si-alloy particles during repeated expansion and contraction. Cost reduction and an increase in the volumetric energy density of Li-ion cells can possibly be attained using silicon-based negative electrodes.1-3 While these benefits have been known for many years, commercial implementation remains limited due to silicon's inherent challenges.Si-based electrodes can suffer from particle pulverization upon cycling (e.g. micron sized silicon), 1,4-7 can have loss of electronic contact between particles after repeated expansion and contraction during cycling, 2,8-10 can have high irreversible first cycle coulombic efficiency 11,[12][13][14][15] and can have low coulombic efficiency (compared to graphite electrodes) during cycling. 1,16,17 The root cause of these challenges is the large volume expansion 1,18 of the material during lithiation and subsequent contraction during delithiation. Several issues have been solved through the use of either nanosized Si or nanostructured Si. For instance the use of nanosized Si particles solves some of the pulverization problems encountered by micron sized Si particles. 1,4,11,[19][20][21] However, nanosized Si still suffers from inter-particle contact loss during cycling as well as low coulombic efficiency due to large specific surface area. 1,16In a recent publication, Chevrier et al. 16 presented a rational way to design commercially relevant Si-based negative electrodes. They showed that confining nano-domains of silicon in an alloy matrix suppresses particle pulverization and greatly reduces the surface area of Si exposed to the electrolyte. This approach yields materials with lower first cycle irreversible capacity (IRC) loss, better coulombic e...

The first charge–discharge cycling behaviors of two sets of Li–Ni–Mn–Co type positive electrode materials were compared. The samples in each set have similar Ni–Mn–Co ratios but different Li-to-total metal ratio (Li/M). The samples that were Li-rich with a Li­[Li x M(1–x)]­O2 structure showed a typical 4.5 V “oxygen loss” plateau and a typical irreversible capacity loss near 25%. Surprisingly, other samples with lower Li/M ratios that still exhibited a 4.5 V “oxygen loss” plateau exhibited an irreversible capacity loss as low as 4.0% of their first charge capacity. XRD analysis revealed that all samples were single-phase layered oxides. A separate and a detailed XRD analysis combined with dQ/dV analysis showed that the reduced irreversible capacity loss was not caused by the admixture of a spinel phase. ICP-OES results and the oxidation state versus atomic occupancy rules suggested the presence of metal site vacancies in the pristine materials with low IRC, which were confirmed by densities measured with a helium pycnometer. The results presented here show that the small irreversible capacity is a consequence of (a) metal site vacancies, leading to Li­[□ q M(1–q)]­O2 structures, where □ is a metal site vacancy, which leads to (b) no Li atoms in the transition metal layer. These materials still have Li/M > 1, so they are “Li-rich”, but they are “traditional layered materials” with no Li in the transition metal layer. This study identifies a new route for fabricating high capacity Li-rich positive electrode materials with small irreversible capacity loss.

Lithium-rich layered Ni–Mn–Co oxide materials have been intensely studied in the past decade. Mn-rich materials have serious voltage fade issues, and the Ni-rich materials have poor thermal stability and readily oxidize the organic carbonate electrolyte. Core–shell (CS) strategies that use Ni-rich material as the core and Mn-rich materials as the shell can balance the pros and cons of these materials in a hybrid system. The lithium-rich CS materials introduced here show much improved overall electrochemical performance compared to the core-only and shell-only samples. Energy dispersive spectroscopy results show that there was diffusion of transition metals between the core and shell phases after sintering at 900 °C compared to the prepared hydroxide precursors. A Mn-rich shell was still maintained whereas the Co which was only in the shell in the precursor was approximately homogeneous throughout the particles. The CS samples with optimal lithium content showed low irreversible capacity (IRC), as well as high capacity and excellent capacity retention. Sample CS2-3 (the third sample in the 0.67Li1+x (Ni0.67Mn0.33)1–x O2·0.33Li1+y (Ni0.4Mn0.5Co0.1)1–y O2 CS2 series) had a reversible capacity of ∼218 mAh/g with 12.3% (∼30 mAh/g) irreversible capacity (IRC) and 98% capacity retention after 40 cycles to 4.6 V at 30 °C at a rate of ∼C/20. Differential capacity versus potential (dQ/dV versus V) analysis confirmed that cells of the CS samples had stable impedance as well as a very stable average voltage. Apparently, the Mn-rich shell can effectively protect the Ni-rich core from reactions with the electrolyte while the Ni-rich core renders a high and stable average voltage.

Li-rich positive electrode materials (e.g. Li 1.2 Ni 0.13 Mn 0.54 Co 0.13 O 2 ) are potential candidates for high energy density Li-ion batteries. [1][2][3] They are capable of delivering reversible specific capacities up to 250 mAh/g 4,5 at an average discharge potential of ∼3.5 V vs Li metal. 6Understanding the structure of Li-rich materials is essential to improve their properties and performance. Li-rich materials are layered transition metal (TM) oxides comprised of alternating layers of metal atoms (Li or TM) and oxygen atoms. 7 Compared to non-Li-rich layered transition metal oxides such as LiCoO 2 , the Li/TM ratio is greater than 1 for Li-rich layered transition metal oxides and usually Li atoms occupy the TM layer in addition to the Li layer.8 Li 2 MnO 3 is a typical example of such a material in which 1 4 of the Li atoms occupy sites in the TM layer.Li-rich layered transition metal oxides can be defined as O3 structures with A-B-C-A-B-C stacking but the arrangement of the atoms in the TM layer is different from that in other layered materials. 9 The presence of Li + ions with large ionic radii (0.74 Å) and small sized Mn 4+ ions (0.54 Å) in the TM layer causes an in-plane ordering resulting in a √ 3a × √ 3a superstructure or superlattice 10 and changes the symmetry from R-3m to C2/m. The superlattice ordering in the TM layer results in superstructure Bragg peaks in the range of ∼ 20• to 35

LiCo 1-2x Mg x Mn x O 2 (0 ≤ x ≤ 0.05) materials were prepared from Co 1-2x Mg x Mn x (OH) 2 (0 ≤ x ≤ 0.05) co-precipitated precursor materials by mixing precursor materials with stoichiometric amounts of Li 2 CO 3 and heating to 900 • C for 10 h. All precursor and lithiated materials were characterized by Scanning Electron Microscopy, X-ray Diffraction (XRD), Inductively Coupled PlasmaOptical Emissions Spectroscopy and electrochemical testing. In situ XRD was performed on LiCo 1-2x Mg x Mn x O 2 (x = 0, 0.02, 0.05) electrodes while cycling to study the effects of substitution on phase transitions and unit cell variations. Increasing Mg/Mn substitution in the material was found to slightly increase the 1 st charge capacity, decrease the 1 st discharge capacity and increase the 1 st cycle irreversible capacity (3.6 V-4.7 V). Cells with even 1% Mg/Mn doping were shown to have markedly improved cycling performance, and results suggest that the improvements stem from suppressing the cell impedance growth, not from the suppression of the O3-O6-O1 phase transitions. Lithium ion (Li-ion) batteries are used to power cell phones, laptops, other portable electronics, electric vehicles and grid storage. With an ever-increasing demand for energy, battery manufacturers and researchers constantly look for ways to increase the energy density while maintaining a long lifetime. LiCoO 2 (LCO) positive electrode materials, proposed by Mizushima et al. in 1980, 1 have been utilized since the commercialization of Li-ion battery technology by Sony in 1991. 2 Significant improvements have been made since then, but commercially available batteries never push LCO electrodes past 4.48 V (vs Li/Li + ), corresponding to the deintercalation/intercalation of ∼0.7 Li (per LCO) or a capacity of ∼190 mAh/g (out of a theoretical capacity of ∼274 mAh/g). In order to unlock more capacity and increase the energy density, the LCO electrodes have to cycle to voltages above 4.5 V (vs Li/Li + ). However, pushing cells with LCO electrodes to an even higher voltage result in a dramatic decrease in long term cycling performance. [3][4][5][6][7][8][9] There are multiple causes for this drop in performance, including structural instability of highly delithiated LCO,[6][7][8]10,11 electrolyte oxidation 5-7,9,10,12,13 and Co dissolution. 3,5,7,8,14,15 To further understand the structural instability that arises from high voltage cycling, X-ray diffraction (XRD) has been instrumental in studying the unit cell lattice parameter changes and phase formations. 10,11,16,17 Theoretical work has also confirmed the majority of these occurrences and helped shed light in understanding phase formation at low lithium content. 18,19 As Li deintercalates from LCO, the material undergoes a series of phase changes. The first is an insulatormetal transition, resulting in a 2-phase region. 10,16 As the material continues to deintercalate, LCO will undergo an order-disorder transition around 0.5 Li.16,18 Both of these transitions occur reversibly and are not detrimental to cycli...

NaNi0.5Mn0.5O2 is a promising sodium-ion battery cathode material that has been extensively studied. However, the air sensitivity of this material limits practical application and is not well understood. Here, we present a detailed study of NaNi0.5Mn0.5O2 powders stored in different atmospheres (oxygen, argon, and carbon dioxide), either dry or wet. X-ray diffraction and Fourier transform infrared measurements were used to characterize the materials and their surface species before and after controlled-atmosphere storage. It was determined that the exposure of untreated NaNi0.5Mn0.5O2 powders to moisture results in desodiation and material degradation, leading to poor cycling. This effect was found to be caused by reactive surface species. From these results, a simple ethanol washing method was found to significantly reduce the air-reactivity and improve the electrochemical performance of NaNi0.5Mn0.5O2 by removing surface impurities formed by air exposure.