Improved electrochemical performances of Li-rich nickel cobalt manganese oxide by partial substitution of Li+ by Mg2+

Journal of Power Sources 359 (2017) 27e36

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Improved electrochemical performances of Li-rich nickel cobalt manganese oxide by partial substitution of Liþ by Mg2þ bastien Sallard a, *, Denis Sheptyakov b, Claire Villevieille a, ** Se a b

Paul Scherrer Institute, Electrochemistry Laboratory, CH-5232 Villigen PSI, Switzerland Paul Scherrer Institute, Laboratory for Neutron Scattering and Imaging, CH-5232 Villigen PSI, Switzerland

h i g h l i g h t s

g r a p h i c a l a b s t r a c t

 Mg content at 1% does not form spinel under cycling and the potential drop is mitigated.  Mg contents at 2.5% or higher show the inverse phenomena.  Li2MnO3 is active at 25  C.

a r t i c l e i n f o

a b s t r a c t

Article history: Received 20 December 2016 Received in revised form 8 May 2017 Accepted 9 May 2017

Li-rich nickel cobalt manganese oxide (NCM) materials with partial substitution of lithium by magnesium is synthesized by the Pechini process. Different synthesized Li100-2xMgx-NCM materials (x ¼ 0, 1, 2.5, 5, and 10) are investigated by X-ray diffraction (XRD), neutron diffraction, and scanning electron microscopy to determine the role of Mg in the structure and its impact on the morphology and electrochemical properties. The chemical composition, crystal structure, and particle morphology are compared with those of the reference (x ¼ 0) material. Mg substitution has a significant impact on the electrochemical properties. In comparison with the reference sample, the x ¼ 1 sample exhibits a mitigation of the voltage drop owing to more stable structure during cycling, leading to a specific discharge of 210 mAh g1 after 100 cycles at C/10 rate. However, compositions with x  2.5 exhibit larger voltage drop in discharge, due to the faster formation of a spinel-like structure during cycling. © 2017 Elsevier B.V. All rights reserved.

Keywords: Batteries Cathode material X-ray diffraction Electrochemistry Li-rich NCM

1. Introduction Lithiated nickel cobalt manganese oxides (LiNi1-x-yCoxMnyO2, NCM) are considered as the most promising insertion material for the next generation of high-energy Li-ion batteries [1]. They were

* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (S. Sallard), [email protected] (C. Villevieille). http://dx.doi.org/10.1016/j.jpowsour.2017.05.028 0378-7753/© 2017 Elsevier B.V. All rights reserved.

first introduced in 2001 by Ohzuku et al. [2], as high-voltage materials able to cycle above 4 V vs. Liþ/Li, with a high specific charge between 160 and 180 mAh g1. High-energy NCM (HE-NCM) materials, also termed Li- and Mn-rich nickel cobalt manganese oxides, with formula a (Li2MnO3)$b (LiNi1-x-yCoyMnxO2) were introduced in 2007 by Thackeray et al. [3,4]. The main advantages of HE-NCM materials are a specific charge between 250 and 290 mAh g1 and an average discharge voltage close to 3.5 V vs. Liþ/ Li, which lead to a gravimetric energy density of 900 mWh g1 (higher than the corresponding values < 600 mWh g1 of LiFePO4 and LiCoO2). HE-NCM materials are composed of two phases,

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monoclinic Li2MnO3 (space group C2/m) attributed to a superstructure and rhombohedral LiMO2 (space group R-3m). During the first charge, HE-NCM is activated at high voltage (> 4.5 V vs. Liþ/Li), i.e., the Li2MnO3 phase is irreversibly oxidized to form Li2O, from which O2 gas is finally released [5e8]. The commercialization of HE-NCM as cathode materials in Liion batteries is hindered by the instability of their electrochemical performance, i.e., by the fading of the specific charge (a common issue for most battery systems) and, more problematic, by a drop of the effective potential during the discharge [9]. One of the possible explanations for this behavior comes from the instability of the carbonate-based electrolytes commonly used [10,11]. The side reactions that occur at high voltages at the cathode/electrolyte interface may be responsible for the capacity fading but not for the potential drop. Structural changes in the HE-NCM material during cycling are the main causes for the latter effect. Different mechanisms are involved in this case, such as the dissolution of transition metal ions in the electrolyte and their partial deposition on the anode, the migration of transition metals from the metal slabs to the lithium interslabs, the formation of spinel phases, etc [9,12e17]. Different approaches have been proposed to limit the degradation of the electrochemical performance of HE-NCM, mainly involving coating or modifying the electrolyte. The coating approaches [18e24] involve applying protective layers (by atomic layer deposition, ALD, or coating in solution) in order to stabilize the electrode/electrolyte interface. Even though good results have been achieved by this method, the gains are only temporary, corresponding to the time for the coating to be leached by the HF present in the electrolyte. Moreover, the effectiveness of the coating is highly dependent of its thickness and uniformity, and so far no improvement in the voltage drop has been reported [19]. The use of electrolyte additives [25e29] can significantly stabilize the solid polymeric interface (SPI) to prevent its further growth and consequently limit the increase in the cell resistance during longterm cycling, either in half-cell or in full-cell configuration. Additives react faster at the surface of the electrodes than ethylene carbonate (EC) and dimethyl carbonate (DMC). The degradation of EC and DMC is then reduced compared to standard electrolytes, but similarly to the coating approach no significant improvement in the potential drop has been reported. From a structural perspective, different doping strategies have been published, involving cationic transition metals as partial replacement of Mn, Ni, and/or Co [30e33]. The reported improvements mostly concern better redox kinetics and possibly a reduced fading of the specific charge. Partial substitution of O2 anions by F was also proposed as a way to stabilize the layered structures with the formation of strong M-F bonds [34]. Interesting results have been published but further 7Li and 9F NMR studies showed that F ions were not incorporated in the crystal structure and incomplete LiF coating was achieved at the surface [35,36]. Another recent approach to stabilize the HE-NCM structure is the partial replacement of the mobile Liþ with bigger alkali ions such as Naþ or Kþ [37e39]. The idea is to take advantage of the much larger ionic radius of Naþcompared to that of Liþ and M2þ (M ¼ Ni, Mn, Co) metals, in order to increase the segregation between the alkali and the transition metal ions. Even if the possibility of the Liþ/Naþ exchange occurring during cycling is still debated, the main results are: i) a significant stabilization of the specific charge, ii) the limitation of the growth of the spinel phase, iii) the stabilization of the cell resistance, and iv) the reduction of the potential drop. Similar results have been reported for K-doped HE-NCM by Li et al. [39]. The authors proposed that the less mobile Kþ ions hinder the Mn migration. Another interesting idea proposed by Zhang et al. [40,41] is to dope HE-NCM with large phosphate, or sulfate or silicate anions introduced in the metal layers, in order to slow down

the structural changes during cycling. Recently a new concept to stabilize the structure of layered overlithiated metal oxide materials with the partial substitution of the Liþ with Mg2þ has been reported [42,43]. Mg2þ possesses an ionic radius similar to the one of Liþ [44], and it is considered as quasiimmobile, due to the strong coulombic interactions between the Mg2þ in the Liþ interlayer and the O2 present in the metal-oxygen layers [45,46]. The coulombic repulsion between Mg2þ and Myþ is expected to be strong enough to hinder any movement of the transition metal during cycling from the metal slab to the Liinterslabs (Fig. 1) [46]. The Mg-doped HE-NCM [43] shows a slight stabilization of its specific charge and a partial mitigation of the potential drop over cycling. In this paper, we present Mg-HENCM materials, with Mg2þ used as partial substituent of Liþ. The goal is to understand the role of the Mg in the electrochemical performance and voltage drops but also on the structure especially during the ageing. 2. Experimental 2.1. Synthesis In a typical synthesis, lithium nitrate (98%, Aldrich, 3.022 g), nickel nitrate hexahydrate (97%, Aldrich, 2.079 g), cobalt nitrate hexahydrate (98%, Aldrich, 1.086 g), manganese nitrate tetrahydrate (99%, Aldrich, 5.072 g), and citric acid monohydrate (99.5%, Aldrich, 32 g) were dissolved in demineralized water (250 mL) at room temperature. The solution was heated up to 100  C under stirring until a dried powder was obtained. The powder was directly calcined under ambient atmosphere at 300  C, ground, and finally calcined in a muffle oven at 450  C (with heating up, plateau, and cooling to room temperature steps of 6, 12 and 6 h, respectively), then ground and calcined again at 850  C following the same procedure. The chemical composition of the reference HE-NCM is then Li1.171(Ni0.191Co0.099Mn0.539)O2 and was confirmed by ICP-MS. Li100-2xMgx NCM materials were obtained by replacing the desired amount of LiNO3 by magnesium nitrate tetrahydrate in a 2:1 ratio. Mg-substituted NCM are hereafter labelled Mgx-NCM. As example, the chemical composition of Mg1-NCM is Li1.148Mg0.012(Ni0.191Co0.099Mn0.539)O2. 2.2. Scanning electron microscopy The scanning electron microscopy (SEM) images were recorded by a Carl Zeiss UltraTM 55 (Germany) apparatus at a 3 kV voltage using the in-lens detector. The powders were sputtered with gold by Ar plasma. 2.3. X-ray diffraction measurements X-ray diffraction (XRD) patterns of the pristine materials were obtained from the powder, measured with a zero-Si background sample holder at room temperature over a 9.5e95.0 (2q) angular range with a 0.033 step size for total duration of 11 h, using a PANalytical Empyrean diffractometer equipped with a Cu Ka source. The ex situ patterns were measured with capillaries (previously sealed in Ar filled glove box with wax) using the same equipment with a 9.5e90.0 (2q) angular range and a 0.033 step size for a total duration of 11 h. 2.4. Neutron diffraction measurements The powder neutron diffraction experiments were carried out on the HRPT/SINQ beamline at the Paul Scherrer Institute. A neutron wavelength l ¼ 1.494 Å was used, and the diffracted

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Fig. 1. Concept of potential drop mitigation for the HE-NCM by partial Li/Mg substitution.

neutrons were recorded over a range of internal scattering angles of 5e165 at a temperature of 300 K. 2.5. Cell preparation and electrochemical measurements An Al foil was coated with slurries composed of HE-NCM, Super C carbon black, and Kynar Flex HSV 900 PVDF (polyvinylidene fluoride) binder in 80:10:10% weight ratios, respectively. After drying, the electrodes were cast, and half-cells were assembled in an Ar-filled glove-box using LP30 as electrolyte (1 M LiPF6 in a 1:1 EC-DMC solvent, BASF), Li metal (Alfa Aesar) as counter electrode, and a glass fiber separator. The cells were electrochemically cycled from 2.5 to 4.8 V vs. Liþ/Li by cyclic voltammetry (scan rate 50 mV s1) and galvanostatic cycling at C/10 rate. 3. Results 3.1. Structural characterization In order to assess the role and position of Mg inside the structure, we analyzed a combination of X-ray and neutron diffraction data. All HE-NCM materials are usually represented as a combination of two distinct crystal phases: a rhombohedral (R-3m) and a monoclinic (C2/m) [3,4] in agreement with our data. The rhombohedral (R-3m) phase corresponds to “NCM” such as the transition metal layers alternating with lithium layers. The monoclinic phase, attributed to Li2MnO3 phase, is often considered as a superstructure. Fig. 2a shows the full 2q range of the XRD patterns. The patterns of HE-NCM and Mg1-NCM are very similar, with peaks located at the same positions and approximately the same intensities. Increasing the amount of Mg leads to minor structural modifications for Mg1-NCM (Li100-2xMgx NCM), but tends to form an

additional phase for samples with x  2.5 (the corresponding peaks are marked with * in Fig. 2a and b). The formation of this phase is better evidence for Mg10-NCM. The additional phase detected for Mg2.5-NCM, Mg5-NCM and Mg10-NCM compounds was identified as a cubic MnMg6O8-type structure (JCPDS 011-0031) considered here as an electrochemical inactive impurity. Another important feature to note at this stage is the appearance of secondary peaks belonging to the Li2MnO3 phase (not visible in the reference sample and starting to appear for Mg2.5-NCM) (labelled with þ in Fig. 2a and b). The same observation can be made for the neutron diffraction patterns presented in Fig. 2d, with the Mg10-NCM pattern differing significantly from that of the reference sample. From this basic analysis, we can infer that while the substitution of Li by Mg leads to form and stabilize higher amounts of the Li2MnO3 phase, it also leads to the formation of the previously reported cubic phase MnMg6O8-type, which presumably results in a lower specific charge at high Mg contents. As we are substituting a monovalent by a divalent cation, even if the radius of these cations is almost the same and the total cationic charge of the materials remains unchanged upon substitution, the impact of the substitution in term of steric effects and repulsion between the layer/interlayer is not obvious and must be assessed. To this purpose, we performed some Rietveld refinements based on the neutron diffraction data, in order to identify the position of Mg inside the structure. A typical Rietveld refinement plot is presented in Fig. 3. As can be seen in Fig. 3, the Mg10-NCM sample was refined using three different phases: NCM (R-3m), Li2MnO3 (C2/m), and MnMg6O8. The unit cell parameters of each phase are consistent with those reported in the literature and in their respective JCPDS cards (056-0147, 018-0737, and 011-0031). The relative ratios of the

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Fig. 2. XRD patterns of Mgx-NCM and pristine HE-NCM; (a): full spectra; (b and c) magnification for specific 2q ranges. Neutron diffraction of Mgx-NCM and pristine HE-NCM; (d): full spectra; (e and f) magnification for specific 2q ranges. The stars and the cross denote the peaks of the MnMg6O8-type phase and the Li2MnO3 phase respectively.

Fig. 3. Rietveld refinement of the neutron diffraction pattern of Mg10-NCM. The star denotes an unidentified minor impurity.

Li2MnO3, MnMg6O8, and NCM phases calculated from the Rietveld refinement were 64.4/24.5/11.1 wt%, respectively. Based on the refinement data, we found out that Li2MnO3 can be doped by Mg in both the lithium sites (0% (Li3 site) and 4% (Li1 site)) and in the transition-metal sites (~13% of the (Mn) site) leading to the following chemical composition (Li1-xMgx)2(Mn1-yMgy)O3 (Table S1, Supplementary Information). 3.2. Morphological characterization SEM pictures of the as-synthesized samples are shown in Fig. 4 and Fig. S1 (Supplementary Information). HE-NCM and Mg1-NCM

exhibit the same morphology, i.e., irregular round-shaped primary particles with a diameter between 50 and 200 nm. For x ¼ 2.5, 5, and 10, two different particle morphologies are visible: i) the same one as described above, and ii) irregular truncated-interlocked octahedra of 400e600 nm diameter. The i/ii population ratio obviously shows a constant decrease when x increases. To identify a possible phase segregation, EDX mappings were performed on Mg1-NCM and Mg10-NCM materials (Fig. S2, Supplementary Information). The Mg1-NCM material presents an homogeneous distribution of the elements, whereas the Mg10-NCM material shows an obvious elemental segregation. We found out two different chemical composition between i) particles composed of Mn,Mg, Co, Ni and O elements and, ii) particles composed of the Mn, Mg and O elements (please remember than Li cannot be tracked by EDX). It appears reasonable for the Mg10-NCM material to consider that the former morphology corresponds to the NCM phase (11.1 wt%) and the latter one corresponds to Li2MnO3 (64.4 wt%) and MnMg6O8 (24.5 wt%) phases. Those results are in agreement with the SEM pictures (Fig. 4 and Fig. S1, Supplementary Information). 3.3. Electrochemical characterization 3.3.1. Specific discharge retention The evolution of the specific discharge is plotted in Fig. 5 (the corresponding coulombic efficiency is plotted in Fig. S3, Supplementary Information): as expected, the higher is the Li substitution level, the lower the specific discharge. 3.3.1.1. Activation process (1st cycle). The values of specific charge, discharge, and irreversibility obtained along the first cycle for the different materials are reported in Table 1. The materials can be divided in two different groups: i) 1  x  2.5, and ii) 5  x  10. For

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Fig. 4. SEM micrographs of HE-NCM (a), Mg1-NCM (b), Mg2.5-NCM (c), Mg5-NCM (d), and Mg10-NCM (e) materials.

Table 1 Specific charge (Sp. ch), specific discharge (Sp. dsc.) and irreversibility (Irr.) along the first cycle for different Mgx-NCM samples. Sample

Sp. ch. [mAh g1]

Sp. dsc. [mAh g1]

Irr. [mAh g1]

Irr. [%]

HE-NCM Mg1-NCM Mg2.5-NCM Mg5-NCM Mg10-NCM

318 304 303 270 117

254 249 254 194 82

64 55 49 77 34

20 18 16 28 29

point out a higher amount of Li2MnO3 in Mgx-NCM with x  5 materials which is directly linked to be low reversibility.

Fig. 5. (Top) Evolution of the specific discharge of the different Mgx-NCM samples; (bottom) magnification of the specific discharge in the 180 to 260 mAh g1 range.

the former, the irreversible specific charge (attributed mostly to the activation of Li2MnO3) is directly linked to the amount of Mg and contributes to around 16e18% of the total specific charge of the 1st cycle. Since the irreversibility is lower for low Mg-content samples (x  2.5), we conclude that a small substitution of Liþ with Mg2þ improves the efficiency of the first oxidation. However, for the latter case 5  x  10, we found that the irreversible charge is close to 30%. This result is consistent with the structural analyses that

3.3.1.2. Following cycles. Mg1-NCM material exhibits a specific discharge close to that of the reference sample (245 mAh g1 for the 2nd cycle, ca. 220 mAh g1 after 20 cycles, and finally ca. 215 mAh g1 after 50 cycles). The specific discharge of the Mg1NCM material is around 10e13 mAh g1 lower than the one of HENCM. An important point to note is that the specific discharge becomes constant after 65 cycles at ca. 210 mAh g1 with no more fading, whereas the one of HE-NCM regularly decreases at a rate of 0.025 mAh g1 cycle1. This is an indication that the Mg substitution leads to the stabilization of the specific discharge. Increasing the Mg content leads to significant changes in the electrochemical properties. In the case of Mg2.5-NCM, the evolution for the first 15 cycles is similar to that of the reference sample, with the specific discharge decreasing from 250 mAh g1 (2nd cycle) to 230 mAh g1 (15th cycle). The specific discharge then continuously decreases until the 60th cycle, where it reaches 217 mAh g1.

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Interestingly, from the 60th cycle to the 100th, the specific discharge increases to reach a value of 223 mAh g1 at the 100th cycle (increase rate of þ0.073 mAh g1 cycle1 from the 60th to 100th cycle). The Mg5-NCM material presents a rather low 2nd cycle specific discharge, close to 190 mAh g1. The specific discharge then increases to reach ca. 200 mAh g1 after 5 cycles and then decreases from the 5th to the 22nd cycle, where it reaches a value of 195 mAh g1. After the 22nd cycle, the specific charge continues to increase to finally reach a value close to 210 mAh g1 (the increase rate of the specific discharge was found to be þ0.345 mAh g1 cycle1 from the 35th to the 75th cycle). From the 85th cycle until the 100th cycle, the specific discharge is stable around 210 mAh g1. Finally, the specific discharge measured for Mg10-NCM continuously increases from the 2nd to the 100th cycle. The gain is relatively high at the beginning, from 82 to 110 mAh g1 between the 1st and 10th cycle, followed by a slower increase from the 11th cycle to the 100th cycle where 142 mAh g1 is obtained (increase rate of þ0.380 mAh g1 cycle1 from the 25th to the 100th cycle). Such increase is attributed to the delayed activation of the Li2MnO3 phase at least for the first ten cycles (Fig. S4, Supplementary Information). This result is in agreement with the very large particles identified by SEM. Indeed, the bigger the particle size, the more difficult is to delithiate the monoclinic phase and to release the gaseous O2 [47]. Note that the low specific discharge of the Mg10NCM is attributed to the presence of the inactive MnMg6O8 phase (24.5 wt%, see part 3.1). 3.3.2. Evolution of the voltage drop (galvanostatic cycling) 3.3.2.1. 1st cycle. Fig. 6a shows the normalized galvanostatic curves for the 1st cycle. Two distinct regions can be identified: (I) a first one in which the potential is constantly evolving in a monotonic

increase (typical of insertion/extraction mechanisms), followed by (II) a long potential plateau called activation plateau, located at around 4.45e4.55 V vs. Liþ/Li [7,8]. Interestingly, the length of this potential plateau increases with the Mg content (in agreement with the amount of Li2MnO3 phase in Mgx-NCM). The potential plateau is reached at ca. 35% of the total charge for the reference sample HENCM and at ca. 25% of the total charge in the case of Mg10-NCM. It can also be noticed that higher the Mg content in the materials is, higher is the polarization for the 1st cycle, especially in region (I). On discharge, the samples with 0  x  2.5 exhibit a monotonous decrease of potential with similar polarization, whereas the polarization is higher for the samples with 5  x  10, and an additional process appears between 3.0 and 3.5 V vs. Liþ/Li (indicated by a red arrow). For sake of clarity, the derivative curves of all the samples for the first discharges are displayed in Fig. S5 (Supplementary Information).

3.3.2.2. 2nd cycle. In order to determine if the Li2MnO3 phase was properly activated along the 1st cycle, we plotted in Fig. 6b the normalized galvanostatic curves of the 2nd cycle. Once again the cycling behavior for the materials can be divided in two categories, for low Mg contents (0  x  5), the curves in charge and discharge look all similar; on discharge we observe a slightly higher polarization for samples Mg2.5-NCM and Mg5-NCM. However, for higher Mg contents (x > 5) the Mg10-NCM sample shows a different reaction mechanism. During the charge, the profile is composed of four regions, out of which the regions (II) and (III) are similar to the reference sample HE-NCM, except for a polarization higher by ca. 200 mV. The region (I) shows a difference in shape and polarization, whereas the region (IV) cleary shows that the Li2MnO3 phases was not fully activated along the first cycle since a small potential

Fig. 6. Normalized galvanostatic curves for the 1st cycle (a), 2nd cycle (b), 50th cycle (c) and 100th cycle (d) of HE-NCM (black) and Mgx-NCM materials (x ¼ 1, red; x ¼ 2.5, blue; x ¼ 5, grey and x ¼ 10, green). The vertical doted-dashed lines symbolize the frontier between the regions (i) and (ii) or (I), (II), (III) and (IV). The oblique doted line in Fig. 6a illustrates the shift of the frontier between the regions (I) and (II) for each sample depending of the relative length of the activation plateau during the 1st oxidation. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

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plateau can be identified. At the beginning of the discharge, we can notice an additional phenomenon, at around 4.5 V vs. Liþ/Li, where a small potential plateau appears (indicated by a blue arrow).

3.3.2.3. Long-term cycling. After the 50th cycle (Fig. 6c), the potential in charge follows the same trend for all materials even if some small differences in overpotential can be found at the beginning of the charge (region (I)). However, clear differences emerge in discharge for the Mg10-NCM sample. Indeed, a change in slope is observed at ca. 3.5 V vs. Liþ/Li (between region (i) and (ii)), giving rise to a potential plateau from 0.2 to 1 in normalized specific charge (region (ii)). For the 100th cycle, the precited behaviors observed for the 50th cycle are similar except that all the features are more pronounced, such as a higher overpotential observed on charge (mainly in region I). In discharge, similar to Mg10-NCM, Mg5NCM exhibits too a change in slope at ca. 3.3e3.4 V vs. Liþ/Li (between region (i) and (ii)), leading to a long potential plateau (region ii). To sum up the findings regarding the effect of Mg on the potential, we plotted the evolution of the average potential observed along cycling for all the samples in Fig. 7. We compared the fading rate per cycle for Mg1-NCM (0.80 mV cycle1) and the reference sample HE-NCM (1.31 mV cycle1) proving a mitigation of the potential drop of ca. 40%. However, the materials with x > 2.5 show a more pronounced voltage drop than the reference one (Fig. 7). This phenomenon is attributed to the appearance of the spinel-like phase evidence after long-term cycling in Fig. 6d with the appearance of a potential plateau (evidence also in the derivative of the galvanostatic curves, Fig. S5, and in the cyclic voltammetries, Fig. S6, Supplementary information) [37,38]. To sum up, all these results suggest that a low Mg content hinders the transition into spinel-like phase leading to a stabilization of the voltage drop during cycling; on the other hand, high amounts of Mg promote the formation of the Li2MnO3 thus of the spinel-like structure during cycling, which is normally observed in the reference sample HE-NCM [48,49].

Fig. 7. Evolution of the average discharge potential vs. Li/Liþ of the HE-NCM (black) and Mgx-NCM materials (x ¼ 1, red; x ¼ 2.5, blue; x ¼ 5, grey and x ¼ 10, green). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

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3.4. Structural evolution during cycling We followed the structural evolution during cycling of the two extreme cases Mg1-NCM and Mg10-NCM, and compared their behavior to the reference sample HE-NCM. The analysis was done at two different stages of lithiation, after 1 cycle and after 100 cycles. Ex situ XRD spectra collected after one cycle (Fig. 8a) show different features depending of the material analyzed. The peaks attributed to Li2MnO3 phase (19 < 2q < 25 ) vanish in the HE-NCM diffractogram as described by Li et al. [9], whereas they are still present in the Mg1-NCM and Mg10-NCM ones, in agreement with the literature for Mg-doped HE-NCM materials [43]. Regarding the main peak (003) of the rhombohedral (NCM) phase is similar in HENCM and Mg1-NCM samples. Whereas a shoulder localized at the right edge of the main peak (002) is appearing for the sample Mg10NCM (blue arrow in Fig. 8b) and it was previously attributed to the spinel phase [9,12,26,37,38]. Note that the (002) peak of Mg10-NCM phase is shifted to a higher angle compare to the (003) peak of the HE-NCM and Mg1-NCM (Fig. S7, Supplementary information). After 100 cycles (Fig. 8 c), we notice that the changes in the diffractograms are less pronounced for the Mg1-NCM sample than for the reference HE-NCM sample and the Mg10-NCM sample. The most important point is the apparition of another phase in the HENCM reference sample (labelled with * in Fig. 8c) which is attributed to the spinel-like structure LiMn2O4 (JCPDS 088-1749) and LiMnO2 (JCPDS 009-0109), generally associated to the voltage drop observed during cycling [9,12,37,38]. This spinel-like phase is not visible for the Mg1-NCM sample, which leads to a better electrochemical stability, as pointed out in the electrochemistry section. The XRD refinements reveal that the Mg1-NCM phase presents the smallest changes in unit cell parameter and lattice volume compared to HE-NCM and Mg10-NCM (Table 2). The full width at half maximum (FWHM) for the (003) peak of the HE-NCM increases from 0.238 at the 1st cycle to 0.304 at the 100th cycle, whereas the corresponding FWHM increase for Mg1-NCM is more limited, from 0.211 to 0.238 which again points out a better electrochemical stability of the Mg1-NCM [9] [39]. Note that the XRD data supporting the in situ formation of spinel-like phase(s) and obtained by standard apparatus [9,37,38], i.e. excluding synchrotron data [12], are usually described as a peak at ca. 18.6 or a broadening on the right of the main peak (003). The difference with our current results, i.e. minor diffraction peaks 15.5 < 2q < 23 (Fig. 8c), are attributed to the combination of different parameters as the material pristine state, typical of the chemical composition and of the synthetic procedure used, and the cycling conditions (temperature, current density, potential windows). The Mg10-NCM sample presents globally the most pronounced changes in its crystalline structure (see lattice parameters reported in Table 2) which is in agreement with the highest potential drop reported above. The (002) peak of the Mg10-NCM sample is fully split into two sub-peaks and slightly shifted to higher angles (Fig. 8d and Fig. S7, Supplementary Information) [38] [50]. Additionally, the splitting observed in the main peak (located at 2q ¼ 18.5 ) was attributed by Amalraj et al. [50] to the formation of a spinel-like phase. As the background of the diffractogram was corrected and that the contribution of the carbon Super C was removed, we cannot exclude also the presence of an amorphous phase (19 < 2q < 25 ) attributed to the spinel-like phase as well. Some peaks belonging to Li2MnO3 phases are still visible, even if drastically diminished which is in agreement with delayed activation of the Li2MnO3 phase. Finally, the cubic phase MnNixMg(6-x)O8 (labelled þ in Fig. 8c) detected in the pristine Mg10-NCM is still visible and unchanged (in agreement with its lack of electrochemical activity).

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Fig. 8. Ex situ XRD spectra of HE-NCM, Mg1-NCM, and Mg10-NCM; after 1 cycle with (a) full spectra and (b) magnification of the same spectra for two different angle ranges; after 100 cycles with (c) full spectra and (d) magnification of the same spectra for two different angle ranges. The cross þ denotes the peaks originating from the cubic MnMg6O8structure like impurity and the stars * correspond sto the spinel-like phases.

Table 2 Evolution of the lattice parameters of the HE-NCM Mg1-NCM and Mg10-NCM. Data obtained from the refinement of the ex situ XRD data (Fig. 8). Sample

a (Å)

b (Å)

c (Å)

b ( )

HE-NCM after 1 cycle HE-NCM after 100 cycles Mg1-NCM after 1 cycle Mg1-NCM after 100 cycles Mg10-NCM after 1 cycle Mg10-NCM after 100 cycles

4.949 4.988 4.964 4.990 4.920 4.886

8.581 8.649 (6) 8.607 8.649 8.525 (5) 8.755 (231)

5.046 5.069 5.064 5.083 5.028 5.046

109.082 109.137 109.072 109.115 109.221 108.326

4. Discussions By combining the results on electrochemical and structural properties of these materials, we demonstrated two different reaction mechanisms, depending on the Mg content. The most promising material appears to be Mg1-NCM, as we have shown that the small Mg content limits the voltage drop without losing the favorable specific charge of these promising 5 V cathode materials. We also found out that this material present the smallest structural changes, particularly no spinel-like phase is formed after 100 cycles, which can explain the better cycling stability of Mg1-NCM compared to the reference HE-NCM. The materials with high Mg content, such as Mg10-NCM, show a very low initial specific charge (even if the specific charge increases during cycling) which cannot be explained only with the absence of lithium. The structural investigations show the presence of another cubic phase (electrochemically inactive) and an increase in the Li2MnO3 phase content. During cycling, we observed the appearance of a long potential plateau attributed to a spinel-like phase

Volume (Å3) (3) (3) (1) (4) (19) (9)

202.5 206.6 204.5 207.3 199.1 204.9

originated from the in situ transformation of the high Li2MnO3 content. Even though these electrochemical results are less promising for Mgx-NCM materials with high Li substitution levels (x  2.5), they illustrate an effective approach to obtain electrochemically active phase Li2MnO3/spinel, which is normally quite inactive at 25  C. Indeed, Li2MnO3 needs to be activated either by temperature (45e50  C) or by decreasing the particles size (< 20e30 nm) [47,51,52]. Mg-doping of the Li2MnO3 (Table S1, Supplementary Information) appears to promote similarly to Fedoping, the room temperature electrochemical activation of the sub-microscopic Li2MnO3 [53]. Finally, Table 3 summarizes the evolution of the phase composition of the HE-NCM, Mg1-NCM and Mg10-NCM material during the cycling. 5. Conclusion A new class of HE-NCM materials with partial substitution of Liþ by Mg2þ was synthesized. The lamellar structure was mostly kept, however, structural changes were amplified with the increase of

S. Sallard et al. / Journal of Power Sources 359 (2017) 27e36 Table 3 Qualitative evolution of the phase composition of the HE-NCM, Mg1-NCM, Mg10NCM during cycling based on the XRD data. Material

phase

HE-NCM Mg1-NCM Mg10-NCM

Li2MnO3 LiMO2 Li2MnO3 LiMO2 Li2MnO3 LiMO2 MgMn6O8

Ageing state Pristine

1 cycle

100 cycles

þ þþþþ þþ þþþ þþþ þ þþ

e þþþþ þ þþþþ þþ þþ þþ

e þþþþ e þþþþ þ þþþþ þþ

the Mg content, up to the appearance of an additional cubic phase for Mgx-NCM, (x > 2.5). The amount of Mg2þ introduced also had a direct influence on the electrochemical properties. Molar substitution of  2.5% Liþ by Mg2þ is detrimental for applications in Li-ion batteries, due to a decrease in the specific charge and to a dramatic increase of the potential discharge drop (attributed to the formation of the spinellike phase). Interestingly, the introduction of low amounts of Mg, as in Mg1-NCM, results in a mitigation of the potential drop (40% less than the reference HE-NCM sample) and of the specific discharge fading (above 210 mAh g1 after 100 cycles), also with no spinellike formation. Further work is currently focusing on i) the optimization of small Mg contents and ii) the investigation of the electrochemical properties in full-cell configuration. Acknowledgments The authors would like to thank the entire technical and support staff of the HRPT/SINQ PSI beamline. The authors are also grateful to Prof. Dr. Petr Novak for fruitful discussions. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.jpowsour.2017.05.028. Glossary DMC EC HE-NCM PVDF SEM XRD

dimethyl carbonate ethylene carbonate high-energy nickel cobalt manganese oxide polyvinylidene fluoride scanning electron microscopy X-ray diffraction

References [1] J. Yan, X. Liu, B. Li, Recent progress in Li-rich layered oxides as cathode materials for Li-ion batteries, RSC Adv. 4 (2014) 63268e63284. [2] T. Ohzuku, Y. Makimura, Layered lithium insertion material LiCo1/3Ni1/3Mn1/ 3O2 for lithium-ion batteries, Chem. Lett. (2001) 642e643. [3] S.H. Kang, P. Kempgens, S. Greenbaum, A.J. Kropf, K. Amine, M.M. Thackeray, Interpreting the structural and electrochemical complexity of 0.5Li2MnO3$0.5LiMO2 electrodes for lithium batteries (M ¼ Mn0.5-xNi0.5-xCo2x, 0  x  0.5), J. Mater. Chem. 17 (2007) 2069e2077. [4] M.M. Thackeray, S.H. Kang, C.S. Johnson, J.T. Vaughey, R. Benedek, S.A. Hackney, Li2MnO3-stabilized LiMO2 (M¼ Mn, Ni, Co) electrodes for lithium-ion batteries, J. Mater. Chem. 17 (2007) 3112e3125. [5] N. Tran, L. Croguennec, M. Menetrier, F. Weill, P. Biensan, C. Jordy, C. Delmas, Mechanism associated with the “plateau” observed at high voltage for the overlithiated Li1.12(Ni0.425Mn0.425Co0.15)0.88O2 system, Chem. Mater 20 (2008) 4815e4825. [6] E. Castel, E.J. Berg, M. El Kazzi, P. Novak, C. Villevieille, Differential electrochemical mass spectrometry study of the interface of xLi2MnO3$(1-x)LiMO2 (M¼ Ni, Co, and Mn) material as a positive electrode in li-ion batteries, Chem. Mater 26 (2014) 5051e5057.

35

[7] Z.H. Lu, J.R. Dahn, Understanding the anomalous capacity of Li/Li[NixLi(1/3-2x/3) Mn(2/3-x/3)]O2 cell using in situ X-ray diffraction and electrochemical studies, J. Electrochem. Soc. 149 (2002) A815eA822. [8] M.M. Thackeray, C.S. Johnson, J.T. Vaughey, N. Li, S.A. Hackney, Advances in manganese-oxide “composite” electrodes for lithium-ion batteries, J. Mater. Chem. 15 (2005) 2257e2267. [9] Y. Li, M. Bettge, B. Polzin, Y. Zhu, M. Balasubramanian, D.P. Abraham, Understanding long-term cycling performance of Li1.2Ni0.15Mn0.55Co0.1O2-graphite lithium-ion cells, J. Electrochem. Soc. 160 (2013) A3006eA3019. [10] T. Liu, A. Garsuch, F. Chesneau, B.L. Lucht, Surface phenomena of high energy Li(Ni1/3Co1/3Mn1/3)O2/graphite cells at high temperature and high temperature and high cutoff voltages, J. Power Sources 269 (2014) 920e926. [11] C. Villevieille, J.L. Gomez-Camer, M. Hess, P. Novak, Reducing mass transfer effects on the kinetics of 5V HE-NCM electrode materials for Li-ion batteries, J. Electrochem. Soc. 161 (2014) A871eA874. [12] D. Mohanty, J. Li, D.P. Abraham, A. Huq, E.A. Payzant, D.L. Wood III, C. Daniel, Unraveling the voltage-fade mechanism in high-energy-density lithium-ion batteries: origin of the tetrahedral cations for spinel conversion, Chem. Mater 26 (2014) 6272e6280. [13] I.A. Shkrob, A.J. Kropf, T.W. Marin, Y. Li, O.G. Poluektov, J. Niklas, D.P. Abraham, Manganese on graphite anode and capacity fade in Li ion batteries, J. Phys. Chem. C 118 (2014) 24335e24348. [14] F. Yang, Y. Liu, S.K. Martha, Z. Wu, J.C. Andrews, G.E. Ice, P. Pianetta, J. Nanda, Nanoscale morphological and chemical changes of high voltage lithium manganese rich NMC composite cathode with cycling, Nano Lett. 14 (2014) 4334e4341. [15] E.-J. Lee, Z. Chen, H.-J. Noh, S.C. Nam, S. Kang, D.H. Kim, K. Amine, Y.-K. Sun, Development of microstrain in aged lithium transition metal oxides, Nano Lett. 14 (2014) 4873e4880. [16] L. Boulet-Roblin, M. El Kazzi, P. Novak, C. Villevieille, Surface/interface study on full xLi2MnO3$(1-x)LiMO2 (M¼ Ni, Mn, Co)/Graphite cells, J. Electrochem. Soc. 162 (2015) A1297eA1300. [17] C. Villevieille, P. Lanz, C. Bunzli, P. Novak, Bulk and surface analyses of ageing of a 5V-NCM positive electrode material for lithium-ion batteries, J. Mater. Chem. A 2 (2014) 6488e6493. [18] S.T. Myung, K. Izumi, S. Komaba, Y.K. Sun, H. Yashiro, N. Kumagai, Role of alumina on Li-Ni-Co-Mn-O particles as positive electrode material for lithiumion batteries, Chem. Mater 17 (2005) 3695e3704. [19] F. Amalraj, M. Talianker, B. Markovsky, L. Burlaka, N. Leifer, G. Goobes, E.M. Erickson, O. Haik, J. Grinblat, E. Zinigrad, D. Aurbach, J.K. Lampert, J.Y. Shin, M. Schulz-Dobrick, A. Garsuch, Studies of Li and Mn-rich Lix[MnNiCo] O2 electrodes: electrochemical performance, structure, and the effect of the aluminum fluoride coating, J. Electrochem. Soc. 160 (2013) A2220eA2233. [20] Y. Yao, H. Liu, G. Li, H. Peng, K. Chen, Synthesis and electrochemical performance of phosphate-coated porous LiNi1/3Co1/3Mn1/3O2 cathode material for lithium ion batteries, Electrochim. Acta 113 (2013) 340e345. [21] A. Mauger, C. Julien, Surface modifications of electrode materials for lithiumion batteries: status and trends, Ionics 20 (2014) 751e787. [22] Q.R. Xue, J.L. Li, G.F. Xu, H.W. Zhou, X.D. Wang, F.Y. Kang, In situ polyaniline modified cathode material Li [Li0.2Mn0.54Ni0.13Co0.13]O2 with high rate capacity for lithium ion batteries, J. Mater. Chem. A 2 (2014) 18613e18623. [23] M. Bettge, Y. Li, B. Sankaran, N.D. Rago, T. Spila, R.T. Haasch, I. Petrov, D.P. Abraham, Improving high-capacity Li1.2Ni0.15Mn0.55Co0.1O2-based lithium-ion cells by modifying the positive electrode with alumina, J. Power Sources 233 (2013) 346e357. [24] S.J. Shi, J.P. Tu, Y.Y. Tang, Y.Q. Zhang, X.Y. Liu, X.L. Wang, C.D. Gu, Enhanced electrochemical performance of LiF-modified LiNi1/3Co1/3Mn1/3O2 cathode materials for Li-ion batteries, J. Power Sources 225 (2013) 338e346. [25] L. Yang, T. Markmaitree, B.L. Lucht, Inorganic additives for passivation of high voltage cathode materials, Inorganic additive for passivation of high voltage cathode materials, J. Power Sources 196 (2011) 2251e2254. [26] Y. Zhu, Y. Li, D.P. Abraham, Mitigation performance degradation of highcapacity lithium-ion cells with boronate-based electrolyte additives, J. Electrochem. Soc. 161 (2014) A1580eA1585. [27] D.V. Chernyshov, S.A. Krachkovskiy, A.V. Kapylou, I.A. Bolshakov, W.C. Shin, M. Ue, Substituted dioxaphosphinane as an electrolyte additive for high voltage lithium-ion cells with overlithiated layered oxide, J. Electrochem. Soc. 161 (2014) A633eA642. [28] K.S. Lee, Y.K. Sun, J. Noh, K.S. Song, D.W. Kim, Improvement of high voltage cycling performance and thermal stability of lithium-ion cells by use of a thiophene additive, Electrochem. Commun. 11 (2009) 1900e1903. [29] W. Weng, Z.C. Zhang, J.A. Schlueter, P.C. Redfern, L.A. Curtiss, K. Amine, Improved synthesis of a highly fluorinated boronic ester as dual functional additive for lithium-ion batteries, J. Power Sources 196 (2011) 2171e2178. [30] Z.Q. Deng, A. Manthiram, Influence of cationic substitutions on the oxygen loss and reversible capacity of lithium-rich layered oxide cathodes, J. Phys. Chem. C 115 (2011) 7097e7103. [31] X. Jin, Q. Xu, H. Liu, X. Yuan, Y. Xia, Excellent rate capability of Mg doped Li [Li0.2Ni0.13Co0.13Mn0.54]O2 cathode material for lithium-ion battery, Electrochim. Acta 136 (2014) 19e26. [32] B. Song, M.O. Lai, L. Lu, Influence of Ru substitution on Li-rich 0.55Li2MnO3$0.45LiNi1/3Co1/3Mn1/3O2 cathode for Li-ion batteries, Electrochim. Acta 80 (2012) 187e195. [33] H. Park, J. Lim, J. Yoon, K.S. Park, J. Gim, J. Song, H. Park, D. Im, M. Park, D. Ahn, Y. Paik, J. Kim, The effects of the Mo doping on 0.3Li[Li0.33Mn0.67]O2 cathode

36

S. Sallard et al. / Journal of Power Sources 359 (2017) 27e36

material, Dalton T 41 (2012) 3053e3059. [34] S.H. Kang, K. Amine, Layered Li(Li0.2Ni0.15þ0.5zCo0.10Mn0.55-0.5z)O2-zFz cathode materials for Li-ion secondary batteries, J. Power Sources 146 (2005) 654e657. [35] M. Menetrier, J. Bains, L. Croguennec, A. Flambard, E. Bekaert, C. Jordy, P. Biensan, C. Delmas, NMR evidence of LiF coating rather than fluorine substitution in Li(Ni0.425Mn0.425Co0.15)O2, J. Solid State Chem. 181 (2008) 3303e3307. [36] L. Croguennec, J. Bains, M. Menetrier, A. Flambard, E. Bekaert, C. Jordy, P. Biensan, C. Delmas, Synthesis of “Li1.1(Ni0.425Mn0.425Co0.15)0.9O1.8F0.2” materials by different routes: is there fluorin substitution by oxygen? J. Electrochem. Soc. 156 (2009) A349eA355. [37] W. He, D. Yuan, J. Qian, X. Ai, H. Yang, Y. Cao, Enhanced high-rate capability and cycling stability of Na-stabilized layered Li1.2[Co0.12Ni01Mn0.54]O2 cathode material, J. Mater. Chem. A 1 (2013) 11397e11403. [38] M.N. Ates, Q.Y. Jia, A. Shah, A. Busnaina, S. Mukerjee, K.M. Abraham, Mitigation of layered spinel conversion of a Li-rich layered metal oxide cathode material for li-ion batteries, J. Electrochem. Soc. 161 (2014) A290eA301. [39] Q. Li, G.S. Li, C.C. Fu, D. Luo, J.M. Fan, L.P. Li, Kþ-doped Li1.2Mn0.54Co0.13Ni0.13O2: a novel cathode material with an enhanced cycling stability for lithium-ion batteries, ACS Appl. Mater. Interfaces 6 (2014) 10330e10341. [40] H.Z. Zhang, Q.Q. Qiao, G.R. Li, X.P. Gao, PO34 polyanion-doping for stabilizing Li-rich layered oxides as cathode materials for advanced lithium-ion batteries, J. Mater. Chem. A 2 (2014) 7454e7460. [41] H.Z. Zhang, F. Li, G. Pan, G.R. Li, X.P. Gao, The effect of polyanion-doping on the structure and electrochemical performance of Li-rich layered oxides as cathode for lithium-ion batteries, J. Electrochem. Soc. 162 (2015) A1899eA1904. [42] R. Yu, X. Wang, Y. Fu, L. Wang, S. Cai, M. Liu, B. Lu, G. Wang, D. Wang, Q. Ren, X. Yang, Effect of magnesium doping on properties of lithium-rich layered oxide cathodes based on a one-step co-precipitation strategy, J. Mater. Chem. A 4 (2016) 4941e4951. [43] Y.X. Wang, K.H. Shang, W. He, X.P. Ai, Y.L. Cao, H.X. Yang, Magnesium-doped Li1.2[Co0.13Ni0.13Mn0.54]O2 for lithium-ion battery cathode with enhanced cycling stability and rate capability, ACS Appl. Mater. Interfaces 7 (2015),

13104e13021. [44] R.D. Shannon, Revised effective ionic radii and systematic studies of interatomic distances on halides and chalcogenides, Acta Cryst. 32 (1976) 751e767. [45] E. Levi, M.D. Levi, O. Chasid, D. Aurbach, A review on the problems of the solid state ions diffusion in cathodes for rechargeable Mg batteries, J. Electroceram. 22 (2009) 13e19. [46] W. Liu, P. Oh, X. Liu, S. Myeong, W. Cho, J. Cho, Countering voltage decay and capacity fading of lithium-rich cathode material at 60 C by hybrid surface protection layers, Adv. Energy Mater 5 (2015) 1500274. [47] J. Lim, J. Moon, J. Gim, S. Kim, K. Kim, J. Song, J. Kang, W.B. Im, J. Kim, Fully activated Li2MnO3 nanoparticles by oxidation reaction, J. Mater. Chem. 22 (2012) 11772e11777. [48] H. Koga, L. Croguennec, M. Menetrier, P. Mannessiez, F. Weill, C. Delmas, Different oxygen redox participation for bulk and surface: a possible global explanation for the cycling mechanism of Li1.20Mn0.54Co0.13O2, J. Power Sources 236 (2013) 250e258. [49] C. Genevois, H. Koga, L. Croguennec, M. Menetrier, C. Delmas, F. Weill, Insight unto the atomic structure of cycled lithium-rich layered oxide Li1.20Mn0.54Co0.13O2 using HAADF and electron nanodiffraction, J. Phys. Chem. C 119 (2015) 75e83. [50] F. Amalraj, M. Talianker, B. Markowsky, D. Sharon, L. Burlaka, G. Shafir, E. Zinigrad, O. Haik, D. Aurbach, J. Lampert, M. Schulz-Dobrick, A. Garsuch, Study of the lithium-rich integrated compound xLi2MnO3$(1-x) LiMO2 (x around 0.5; M ¼ Mn, Ni, Co; 2:2:1) and its electrochemical activity as positive electrode lithium cells, J. Electrochem. Soc. 160 (2013) A324eA337. [51] A.D. Robertson, P.G. Bruce, The origin of electrochemical activity in Li2MnO3, Chem. Commun. (2002) 2790e2791. [52] X. Dong, Y. Xu, S. Yan, S. Mao, L. Xiong, X. Sun, Toward low-cost, high energy density Li2MnO3 cathode materials, J. Mater. Chem. A 3 (2015) 670e679. [53] M. Tabuchi, Y. Nabeshima, T. Takeuchi, K. Tatsumi, J. Imaizumi, Y. Nitta, Fe content effects on electrochemical properties of Fe-substituted Li2MnO3 positive electrode material, J. Power Sources 195 (2010) 834e844.