Evolutions of Li1.2Mn0.61Ni0.18

Article pubs.acs.org/cm

Evolutions of Li1.2Mn0.61Ni0.18Mg0.01O2 during the Initial Charge/ Discharge Cycle Studied by Advanced Electron Microscopy Adrien Boulineau,* Loïc Simonin, Jean-François Colin, Emmanuel Canévet, Lise Daniel, and Sébastien Patoux CEA-LITEN, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France

ABSTRACT: The evolutions of the structure occurring into the lithium rich cobalt free layered cathode material Li1.2Mn0.61Ni0.18Mg0.01O2 upon the first electrochemical cycle were investigated by the means of high angle annular dark field (HAADF) imaging in a scanning transmission electron microscope and electron diffraction in a transmission electron microscope. They are coupled with electron energy loss spectroscopy (EELS) experiments in order to probe the chemical evolutions occurring during the first charge/discharge cycle. In the pristine material, the analysis of the HAADF images and electron diffraction patterns confirmed the ordering between the cations (Li or Ni with Mn) and the existence of disoriented domains stacked along the c axis. Moreover, the partial solid solution of Ni into Li2MnO3 leading to a composite material is evidenced. Upon the first charge, a loss of material is shown to have occurred, and the presence of a defect spinel phase due to the transfer of transition metal cations to the interslab is clearly established. It is localized at the edge of the particles. This defect spinel phase apparition is confirmed by EELS experiments and identified as (Li)Mn(2−x)NixO4. After the first discharge, the spinel phase is still present, and structural discrepancies from one crystal to another are observed. Also, it seems that all the domains would not have the same behavior upon discharge. KEYWORDS: Li-rich layered oxides,lithium-ion battery, electron diffraction, HAADF STEM imaging, EELS

1. INTRODUCTION To enable the Li-ion battery technology to meet the needs of the automotive industry, cathode materials with higher capacity are expected. Among all the materials available, the lithium rich layered nickel manganese oxides, the formula of which can be written as Li[Li(1/3−2x/3)NixMn(2/3−x/3)]O2 with 0 < x < 0.5, are of particular interest. These cobalt-free materials are low cost, and their high Mn content leads to interesting practical capacities, up to 250 mAh·g−1 or more.1,2 However, these materials exhibit complex structural properties. Indeed, the structure of the layered oxides with formula LiMO2 is based on the α-NaFeO2 type structure (R3m ̅ space group). It can be described as layers of MO6 octahedra (so-called “slabs”) alternatively stacked with layers of LiO6 octahedra (so-called “insterslabs”). In the lithium rich layered manganese oxides, Li(1+x)M(1−x)O2 (M = Mn, Ni,...), extra lithium ions are substituting the transition metal ions into the slabs. Due to their differences of ionic radii, an ordering between Li+ and Mn4+ ions occurs into the slabs. However, the microstructure of © 2012 American Chemical Society

such materials is controversial and is still under debate. Thus, as illustrated in Figure 1a, some research groups are considering the material as a “composite” one being built of two kinds of slabs, that is, MO2 (M = Mn, Ni,...) and Li1/3Mn2/3O2 slabs, as encountered in LiMO2 (M = Mn, Ni,...) and Li2MnO3 materials, respectively.3−6 This point is even more complex because the structures of these two components themselves are still under debate.9,10 On the other hand, others are considering the material as being homogeneous, built only with one kind of slab LiyNixMn(1−x−y)O2, see Figure 1b.7,8 Upon cycling, such materials suffer from an important irreversible capacity. During the first charge, a long and irreversible domain occurs where the potential is almost constant at 4.5 V vs Li+/Li. Lu and Dahn11 hypothesized that this high irreversible capacity could be explained by the Received: April 12, 2012 Revised: August 20, 2012 Published: August 30, 2012 3558

dx.doi.org/10.1021/cm301140g | Chem. Mater. 2012, 24, 3558−3566

Chemistry of Materials

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of 200 kV. EELS spectra were collected in STEM mode with a Gatan Imaging Filter using a dispersion of 0.2 eV per channel and a 2 mm aperture. In such conditions, the energy resolution is 1.2 eV according to the zero loss peak full width at half-maximum. Acquisitions were done in cumulative mode with total acquisition time between 0.5 and 1.0 s. Concerning the pristine material, the powder was ground in ethanol and a droplet of this suspension was then deposited on a lacey carbon coated grid. After electrochemical test, the cells were disassembled in a glovebox under argon and the cathodes were washed with dimethyl carbonate (DMC) several times and then dissolved in N-methyl-2pyrrolidone (NMP). Finally, a droplet of the powder dispersed in DMC was then deposited on a lacey carbon coated grid. For the observations, the grids were transferred from the glovebox to the microscope using a vacuum transfer holder in order to protect the sample from air exposure.

Figure 1. Two discussed microstructures for Li-rich layered oxides Li[Li(1/3−2x/3)NixMn(2/3−x/3)]O2: case of a composite material (a) and case of a one component material (b).

3. RESULTS AND DISCUSSION Figure 2 shows the first electrochemical cycling curve of the material typical for such a lithium-rich layered oxide.1,13−15

simultaneous lithium and oxygen release during the “plateau”. This was reinforced by the evidence of oxygen gas formation by Armstrong et al. using mass spectrometry.5 Such a phenomenon is still not clear, and structural evolutions related to it are expected. A better understanding is needed for improving these materials. This article deals with the examination of the evolutions occurring in Li1.2Mn0.61Ni0.18Mg0.01O2 during the first electrochemical cycle. The pristine material is compared with two samples charged up to 4.4 and 4.8 V vs Li+/Li and one discharged to 2.5 V vs Li+/Li. The combination of high resolution transmission electron microscopy (HRTEM), high angle annular dark field scanning transmission electron microscopy (HAADF-STEM) imaging, electron diffraction, and electron energy loss spectroscopy (EELS) was used to probe the chemical and structural changes involved during the first electrochemical charge/discharge cycle.

Figure 2. Charge−discharge cycling curve from a Li1.2Mn0.61Ni0.18Mg0.01O2/Li cell. A, B, C, and D marks point out the studied materials.

2. EXPERIMENTAL SECTION

Indeed, according to the literature, the first sloppy part for 1.2 > x > 0.9 where the voltage is increasing continuously is usually attributed to the oxidation of the transition metal ions. A Li2O loss from the structure and a release of O2 is then expected during the following long plateau until x = 0.3.16−18 The points placed on the curve depict the materials studied. Besides the pristine material (point A), the others were obtained from three cycled cells. The first one (point B) was stopped during the charge on the sloppy part at 4.4 V vs Li+/Li, the second cell (point C) was stopped after the complete charge (4.8 V vs Li+/ Li), and the third one (point D) was stopped at the end of the first discharge (2.5 V vs Li+/Li). 3.1. Pristine Material. Prior to doing complex characterization, the pristine material was characterized from a chemical point of view using inductively coupled plasma mass spectrometry (ICP-MS). The Li, Mn, Ni, and Mg mass percentages were found to be 10.2, 39.1, 12.2, and 0.3, respectively, in very good agreement with the theoretical ones that are 9.8, 39.6, 12.5, and 0.3. The synchrotron X-ray pattern of the pristine material was analyzed thanks to a Rietveld refinement. The result given in the Figure 3a shows that the material crystallizes in the C2/m space group with the following cell parameters: a = 4.9614 Å, b = 8.5891 Å, c = 5.0392 Å, and β = 109.31°. Furthermore, no impurity has been detected. Due to the complexity of the structure, which contains crystallographic sites that could be shared by more than two cations, it is impossible to obtain reliable results on atomic occupancies from a Rietveld

2.1. Sample Preparation. Li1.2Mn0.61Ni0.18Mg0.01O2 was prepared using an all solid state method. Lithium, manganese, nickel, and magnesium carbonates were carefully mixed in a stoichiometric way using an agate ball mill and hexane as lubricant. The as obtained mixture was heated at 1000 °C for 24 h and rapidly cooled to room temperature. Cycled materials were obtained using coin cells prepared as follows. First, a slurry was obtained by mixing 80 wt % active material, 10 wt % polyvinilidene fluoride (PVDF), and 10 wt % carbon black in Nmethyl pyrolidinone (NMP) solvent. This slurry was coated on an Al current collector using a 100 μm doctor blade and dried at 50 °C overnight. Electrodes (14 mm diameter) were punched and dried under vacuum at 80 °C for 48 h. Finally, coin cells were assembled in an Ar filled glovebox using Li-metal as counter electrode and LiPF6 1 M in EC/PC/DMC in a 1:1:3 vol. ratio as electrolyte. 2.2. Sample Characterizations. Electrochemical cycling was performed with coin cells using a Biologic VSP galvanostat operating in a galvanostatic mode at C/10 (31.6 mA·g−1) between 2.5 and 4.8 V vs Li+/Li. X-ray diffraction was carried out at the European Synchrotron Radiation Facility (ESRF, Grenoble) at the BM20 beamline. The sample was filled in a 0.4 mm capillary glass, and the acquisition was carried out in transmission mode at 0.4959 Å. The refinement was performed using the Fullprof software.12 Electron diffraction experiments were performed using a JEOL 2000FX for recording the selected area electron diffraction patterns and a JEOL 3010 microscope for the nano-beam electron diffraction patterns. HRTEM images were recorded using the JEOL 3010 microscope. STEM images and EELS spectra were recorded using a FEI Cs-corrected Titan microscope running at an accelerating voltage 3559



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Table 1. Results of the Rietveld Refinement in the C2/m Space Groupa C2/m

Sgr

Bragg R-factor = 13.8

a = 4.9614(1) Å, b = 8.5891(2) Å, c = 5.0392(1) Å, β = 109.31(2)° atom

site

x

y

z

occ.theo

occ.refined

Mn Ni Li Li Mn Ni Li Li O O

4g 4g 4g 2b 2b 2b 2c 4h 4i 8j

0 0 0 0 0 0 0 0 0.231(2) 0.255(1)

0.169(2) 0.169(2) 0.169(2) 0.5 0.5 0.5 0 0 0 0.679(6)

0 0 0 0 0 0 0.5 0.677(2) 0.220(1) 0.233(8)

0.75 0.25 0.0 0.6 0.3 0.1 1.0 1.0 1.0 1.0

0.7500 0.1762(2) 0.0738(2) 0.5262(1) 0.3000 0.1738(2) 1.0000 1.0000 1.0000 1.0000

a

Only the occupancies related to Ni/Li exchange occurring into the slab have been refined.

Table 2. Results of the Rietveld Refinement in the R3̅m Space Groupa R3̅m

Sgr

Bragg R-factor = 8.35

a = 2.8638(1) Å, b = 14.266(5) Å atom

site

x

y

z

occ.theo

occ.refined

Li Ni Ni Li Mn O

3b 3b 3a 3a 3a 6c

0 0 0 0 0 0

0 0 0 0 0 0

0.5 0.5 0 0 0 0.258(2)

1.0 0.0 0.2 0.2 0.6 1.0

0.962(2) 0.038(2) 0.162(2) 0.238(2) 0.600 1.0000

a

Only the occupancies related to Ni/Li exchange between the slab and the interslab have been refined.

Figure 3. Refinements of the synchrotron X-ray pattern for the pristine material in both (a) C2/m and (b) R3̅m space groups. A wavelength of 0.4959 Å was used.

mass, Mn and Ni atoms are expected to be brighter than Li, O, and Mg, which are not heavy enough to produce any contrast and, thus, are not visible. The HAADF-STEM observation has the advantage that atomic columns are directly identified, when using a small electron probe and a sufficient collection angle for annular detector,21,22 while complicated simulations of images are needed for the analysis of conventional TEM images. The image reveals a well crystallized material with a very regular stacking of the transition metals layers separated by ≈4.75 Å, which is in agreement with the distance between slabs and with the c parameter value. When looking at the image, it is noticeable that the slabs are not always stacked in the same way. Indeed, three stacking vectors from one slab to the next one can be found. Transitions can occur either without any displacement perpendicularly to the stacking direction or with a translation of 1/3 of the distance between two rows of atoms in a same slab. Therefore, domains are clearly visible in between each change of stacking vectors. The two enlarged parts of Figure 4, parts b and c, revealed the dumbbell character of the atomic columns. Indeed, a splitting of 1.43 Å exists between two atoms, and each dumbbell is separated by 2.85 Å. These distances are in accordance with the projected bond length between two adjacent transition metal (TM) ions and two (TM) ions separated by a Li ion when the cell is viewed along the [11̅0]mon. direction. Thus, each slab in Figure 4b can be considered as built by a succession of atomic columns containing either transition metal ions (TM) or lithium ions

refinement without the use of a second set of data such as neutrons. However, refinements were carried with the goal of getting a qualitative value of Ni/Li exchange between the metal and the lithium slab, and of Ni/Li exchange in the metal slab. As refining both exchanges at once in a C2/m space group would have led to the simultaneous refinement of 4 correlated occupancies for both Ni and Li, a two-step strategy was used. In a first step, as depicted in Figure 3b, the Ni/Li exchange between slabs and interslabs in the R3̅m space group where only one crystallographic site exists for each slab was refined. In a second step, the exchange between Li and Ni into the metal slab was refined in the C2/m space group. Tables 1 and 2 report the starting structural parameters, which correspond to “ideal” LiMO2 and Li2MnO3 structures adapted to our composition, and the refined atomic positions and occupancies. These results show that 3.8% Ni are present in the lithium interslabs and that the exchange between Ni and Li into the slabs, which is about 7.4% is relatively important. However, this is in accordance with previous results obtained on similar materials,19,20 and it does not make the Li2MnO3-like order disappear. The HAADF-HRSTEM image (also called “Z-contrast”) presented in Figure 4 is taken along the [11̅0] zone axis with the pristine material. In such an image, the contrast is proportional to ≈Z2 (Z being the atomic number). Thus, it couples both structural and chemical information at an atomic resolution. In our material, due to their much higher atomic 3560



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means that atomic columns of pure Li do not exist and these two slabs must be seen as a succession of almost identical atomic columns containing mostly TM ions. Thus, the material has to be seen as domains of ordered slabs stacked with changes of stacking vector among which disordered slabs can be observed with a random occurrence. To understand the origin of these changes of stacking vectors leading to domains and to determine whether or not the structure of the material is the same in each domain built of ordered slabs, electron diffraction patterns were collected with crystals of the pristine Li1.2Mn0.61Ni0.18Mg0.01O2 material. Figure 5a presents an HRTEM image where several domains can be seen. Figure 5b is the associated electron diffraction pattern that reflects the structure of the whole crystal. Two sets of spots with different intensities are visible. Considering only the most intense ones, the pattern can be indexed using the [11̅0] monoclinic zone axis, but along the c* direction and between the spots (hhl) and (hh(l+1)), two extra spots are not indexed. The three patterns presented in Figure 5c were collected by nano-beam electron diffraction within the domains noted as 1, 2, and 3 in the Figure 5a. Thus, they reveal the local structure of each domain. As the global one in Figure 5b, they can be indexed considering the same monoclinic structure, but three different orientations must be taken into account. Indeed, the patterns 1, 2, and 3 can be indexed using the [110], [11̅0], and [100] zone axes, respectively. Thus, all the domains built of ordered slabs possess the same crystallographic structure. Thus, the translations of the slabs, as observed in Figure 4, are a consequence of the stacking of the domains with disorientations of about ±60° around the stacking direction. The global pattern in Figure 5b results of the sum of these three patterns, the two less intense extra spots being indexed as (0h̅0) and (hh̅0) along the zone axes [100] and [110], respectively. Moreover, the diffusion running on some spot along the c* direction is related to the random stacking of these domains that present aleatory thicknesses.9 When observing with a lower magnification such images as Figure 6a, one can remark a fluctuation of the contrast; bright and dark regions alternate along the stacking direction. As the

Figure 4. (a) HAADF-HRSTEM image evidencing the way the slabs are stacked. The crystal is mainly oriented along [11̅0], with domains oriented along [110] and [100] being also observable. (b) Enlargement of a perfect stacking of ordered slabs. (c) Enlargement where two disordered slabs can be observed among two ordered ones. The superimposed white line gives a help for visualizing how the slabs are stacked one with other. Note that the slabs stacking direction is from the bottom to the top of the picture.

(Li) with an ordered TM−TM−Li−TM−TM−Li... sequence. The second enlargement, in Figure 4c, presents two slabs among four, along which the intensity is almost constant. This

Figure 5. TEM image showing the stacking of domains (a) and its associated electron diffraction pattern (b). Nano-beam electron diffraction patterns collected in 1, 2, and 3 (c). 3561



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Figure 6. (a) STEM HAADF image from a crystal of the pristine material mainly oriented along [11̅0], with domains oriented along [110] and [100] being also observable. (b) Bragg filtered image obtained selecting only the reflections induced by the ordering between Li and TM ions enhancing the composite nature of the material containing Li2MnO3 parts (blue) and LiNi0.45Mn0.525Mg0.025O2 ones (green). Reflections characteristic of the ordered slabs are circled in blue and those related to the disordered ones are pointed with green arrows on the Fourier transform in the insert.

contrast is closely related with the average atomic number, a variation in composition can be expected. A careful observation reveals a link between both kinds of slabs previously pointed out (i.e., ordered and disordered ones) and the brighter and darker regions. Indeed, Li1/3(TM)2/3 and (TM) slabs are encountered in the darker and brighter parts, respectively. This is consistent with the HAADF imaging contrast rule. To enhance the contrast between these two kinds of slabs, we applied a Bragg filter on the image by selecting only the reflections induced by the ordering between Li and Mn on the Fourier transform (i.e., the blue circled reflections in Figure 6b; see ref 9 for more explanations). In this way, the contrast of the ordered slab (described in the C2/m space group) is accentuated while that of the disordered ones (described in the R3̅m space group) is attenuated. Finally, calculating the proportion of each kind of slab on this representative image leads to a proportion of about 55% of Li1/3(TM)2/3 slabs and 45% of (TM) slabs. Such a result is indicative of a reduced mixing between Mn and Ni. In other words, our material microstructure is in good agreement with the 0.6·Li[Li1/3Mn2/3]O2−0.4·LiNi0.45Mn0.525Mg0.025O2 notation that is the formula of the material Li1.2Mn0.61Ni0.18Mg0.01O2 when rewritten as a two-component material. In this way, according to previous studies,3,6,23 this material should thus be considered as a composite rather than a solid solution material, as stated in refs 7 and 8. 3.2. During the Charge. Crystals were first studied after a cut off voltage at 4.4 V vs Li+/Li, a potential that corresponds to the end of the slope. A typical STEM-HAADF image from a representative crystal is presented on Figure 7. In comparison with the pristine material, the particles look different, presenting variations of contrast that traduce the presence of heterogeneities. One can indeed observe the presence of etched areas that appear as elongated dark areas. These areas present two different orientations: either parallel to the slabs when located at the edge of the particle or perpendicular to the slabs when located in the core. It is relevant to note a degraded

Figure 7. HRSTEM-HAADF images from a particle in the halfcharged material (4.4 V vs Li+/Li) mainly oriented along [11̅0], with domains oriented along [110] and [100] being also observable.

crystallinity into these darker areas located close to the edges. We assume a loss of Li2O occurred into these parts that are similar to those observed when Li2O is electrochemically removed from Li2MnO3.2 The fact that the phenomenon is much more obvious near the edge supports this analysis. Figure 8 presents a conventional TEM micrograph from a crystal after the complete charge (up to 4.8 V vs Li+/Li) with two corresponding electron diffraction patterns having the central rows of reflection in common. The electron diffraction patterns presented in Figure 8 cannot be interpreted considering only the layered structure of the pristine material. Indeed, electron diffraction patterns highlight extra spots located between the diffuse scattering lines. The origin of 3562



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Figure 8. Conventional TEM micrograph from a crystal after the first charge with its corresponding electron diffraction patterns obtained one with the other by a rotation around the c* direction. Yellow and red italic indexes stand for the layered and the spinel structures, respectively.

these extra spots can be explained using a spinel structure in addition to the layered one as detailed elsewhere.24 Such a formation of a new phase can be the result of the diffusion of transition metal ions (TM) from the slab to the interslab space for stabilizing the delithiated structure. Indeed, viewed along the [111]cub. direction, the (Li)Mn(2−x)NixO4 spinel framework can be considered as a layered one constituted by [(M3/4□1/4)OhO2] slabs and [(Li1/2)Td(M1/4□3/4)OhO2] interslabs. This description is consistent with the electron diffraction patterns. Indeed, according to the indexations, the spinel plan (111) is merged with the (001) plan of the layered one structure. As a result, the spinel cell parameter is close to 8.2 Å, which is consistent with the expected value.25,26 Note that cracks are observed at the edge of the crystal and could be the consequence of a strain increasing due to the lithium deintercalation possibly accompanied by an in plane reorganization into the slabs and the spinel formation. This spinel formation is supported by the HRSTEM-HAADF image presented in Figure 9 in which a crystal is viewed along the [010]mon. zone axis. Thus, the image is evidencing the stacking of the layers from the bottom to the top of the image. Contrary to the pristine material, a transition from one phase to another exists in the crystals when going from the bulk to the edge. Indeed, transition metal ions can be seen into the interslab space at the edge of the crystal, while they are only located into the slabs in their bulk. As shown by the enlargement of the two-phases domain, the spinel phase appears progressively from the core to the shell. This observation is consistent with a diffusion of transition metal ions from the slabs to the interslabs that is leading to the spinel structure by a second order transition. According to the HAADF image, the TM ions in the spinel insterslab are present in all the atomic columns instead of occupying half of them and being ordered with vacancies. Indeed, in a perfect spinel seen

Figure 9. (a) HRSTEM-HAADF images of a particle taken from the fully charged material. (b) Detail of the transition between the two phases. The layered structure and the defect spinel one projected along the [010]L and [1̅10]S zone axis, respectively, are superimposed.

along [11̅ 0], pure columns of TM and pure columns of Li must be observed in octahedral and tetrahedral sites, respectively. Here, it is not the case, and this spinel phase must be consequently considered as possessing a defect spinel structure presenting a disordered interslab. 3.3. After the First Complete Discharge. The study of the material after the first discharge (2.5 V vs Li+/Li) revealed the presence of heterogeneities between the crystals examined. To be representative of the whole powder, four kinds of crystals must be considered. Figure 10 presents the representative set of crystals coming from the electrode after a charge and a complete discharge and their associated electron diffraction patterns. From a structural point of view, all of these four crystals are different from one another. Also, it is relevant to note that the third one (Figure 10c) seems to have been much more sheared in comparison with the others. As witnessed by the quasi-inexistence of extra spots between the diffuse lines in its electron diffraction pattern, the first crystal (Figure.10a) is largely composed of the layered structure. On the other hand, their high intensity in the 3563



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compared with Figure 4b, one can remark that the spots induced by the presence of domains oriented along the [100] direction (i.e., spots indexed as 020 and 040) in the pristine material are no longer observable in such crystal after the first discharge. Thus, it seems that the three kinds of domains would not have the same behavior upon discharge. Such small discrepancies from one crystal to another could arise from changes in connectivity, which would result in (partial) electrically isolated particles. Further work is underway to better understand these observed changes. Finally, the HAADFSTEM image presented in Figure 11 confirms the irreversible formation of the defect spinel phase that is still present after the complete discharge at the edge of the crystals.

4. EELS ANALYSIS Typical EELS spectra collected on the pristine, half charged, fully charged, and fully discharged materials are presented in Figure 12. The EELS profiles of O−K, Mn−L2,3, and Ni−L2,3 edges located around 530, 640, and 850 eV, respectively, are

Figure 10. (a), (b), (c), and (d). Crystals with their associated electron diffraction patterns highlighting variations of structure from crystal to crystal at the end of the first discharge. The arrows highlight the diffuse scattering lines induced by the random stacking of the domains along the c axis. The stars are positioned where the rows of extra spots are located when the spinel structure is present.

pattern related to the other crystals (Figure 10b, c, and d) reveal the presence of the spinel structure in a higher proportion among the layered one in these three crystals. The third crystal, which appears as the most sheared, shows a peculiar electron diffraction pattern. Indeed, supplementary diffuse lines can be clearly observed (Figure 10c). Into this pattern, four diffuse scattering lines (instead of two) are located between the rows of spots related to the layered structure. These two extra diffuse lines can easily be attributed to a double diffraction phenomenon of the incident beam on the two first diffuse lines. Moreover, contrary to the previous electron diffraction, the spots related to the different disoriented domains cannot be observed over the diffuse lines. The same double diffraction phenomenon is also observed (with less intensity) into the last electron diffraction pattern related to of a fourth kind of crystal (Figure 10d). Nevertheless, when

Figure 11. HRSTEM-HAADF images from a particle in the fully discharged material. The layered structure and the defect spinel one projected along the [010]L and [1̅10]S zone axis, respectively, are superimposed. 3564



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third peak, labeled γ, located up to 30−50 eV above the first peak is induced by multiple scattering within oxygen shells. Its displacement toward the higher energy losses can be observed from the half charge to the fully charged material. It has been shown that its position can be correlated to the interatomic distances between oxygen atoms, varying inversely with the O− O bond length in transition metal oxides.31 This is in accordance with a structural change from the layered structure to the spinel one where shorter first-neighbor O−O distances exist. Indeed, the shortest O−O distance is 2.72 Å into the layered structure, while it is 2.54 Å into the spinel one. The fact that this peak position does not change during the discharge supports the irreversibility of the spinel formation. Finally, all these evolutions in the EELS spectra support the structural evolutions we have pointed out during cycling using STEM-HAADF imaging and electron diffraction experiments.

Figure 12. Representative EELS spectra obtained near the edge of crystals from the four studied materials. Detail of the O−K edge is given on right.

5. CONCLUSION We confirm in this article the complexity of our lithium-rich layered oxide constituted of stacked disoriented domains. The partial solid solution of Ni into Li2MnO3 that is leading to a composite material that can be written as 0.6·Li[Li1/3Mn2/3]O2−0.4·LiNi0.45Mn0.525Mg0.025O2 is also evidenced. During the first charge, the formation of a defect spinel phase (Li)Mn(2−x)NixO4 localized at the edge of the crystals is demonstrated using imaging, diffraction, and EELS techniques. The layered structure is not recovered after discharge. It is an irreversible phenomenon resulting from the diffusion of transition metal ion from the slabs to the interslabs according to a second order transition. This defect spinel formation does not preclude lithium insertion/deinsertion; it will result in a small lowering of the average reaction voltage and a small decrease of the capacity. Also, the small structural discrepancies observed among crystals after the first discharge reveal the presence of heterogeneities within the electrode, while it is still homogeneous after the charge. Moreover, it seems all the domains do not behave in a same way during discharge and crystals seem to be variably sheared from one to others. We think that some particles may become more or less electrically isolated upon cycling avoiding (or limiting) any electrochemical reaction. The dependence of layered oxides electrochemistry on structure needs to be deeply investigated for a better understanding of the behaviors of the material.

directly linked to the electronic structure of the materials and can be used for probing the chemical and structural evolutions occurring upon cycling. While the evolutions of the oxidation states of the manganese and the nickel on the same material upon cycling have already been studied elsewhere by our group,27 we will focus on the analysis of the O−K edge for discussing the structural changes involved in the material. The O−K edge, the threshold of which is located around 530 eV, is composed of three main peaks. The first one, also called prepeak and labeled α, is attributed in manganese oxides to transitions from 1s core state to oxygen 2p states hybridized with manganese 3d orbitals.28 This peak intensity can be consequently related to the transition metal oxidation state. Also, it has been shown that it was function of the lithium content in lithiated manganese oxides.29 The second peak, labeled β, is usually attributed to the oxygen 2p states hybridized with the transition metal 4s and 4p states.30 As reported in Table 3, we can observe an increasing of the Table 3. Peak α and Peak β positions and Their Ratio for the Pristine Material and the Cycled Ones after a Charge at 4.4 V, 4.8 V, and a Discharge until 2.5 Va

pristine charge 4.4 V charge 4.8 V discharge 2.5 V a

peak α position (eV)

peak β position (eV)

peak α/peak β ratio

529.6 528.2 527.6 528.6

540.1 540 540.2 540.2

0.48 0.58 0.77 0.49

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AUTHOR INFORMATION

Corresponding Author

*Phone: +33 (0)4 38 78 27 88. E-mail: adrien.boulineau@cea. fr.

Peak ratios were calculated integrating peaks over 5 eV windows.

Notes

intensity ratio of the peak α/peak β from the pristine to the charge materials followed by its decreasing during the discharge. This is consistent with the deintercalation of lithium during the charge and its intercalation during the discharge. Moreover, according to the increasing of the splitting between these two peaks during charge followed by its decreasing during discharge, an oxidation of the transition metals during charge followed by a reduction during discharge is expected. This is consistent with our previous study on the same material, in which in situ synchrotron X-ray diffraction was combined with in situ X-ray absorption spectroscopy. Indeed, it has been shown that 26% of the manganese that are Mn3+ in the pristine material was first oxidized into Mn4+ from x = 1.2 to x = 1.0. After which, the oxidation of Ni2+ follows until x = 0.8.27 The

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank the ESRF for access to the synchrotron radiation source and particularly Carsten Baehtz at the BM20 beamline for his help in recording the pattern.

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REFERENCES

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Evolutions of Li1.2Mn0.61Ni0.18

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