Cite this: RSC Advances, 2013, 3,7479

Structural transformation in a Li1.2Co0.1Mn0.55Ni0.15O2

lithium-ion battery cathode during high-voltage hold3

Received 11th November 2012,Accepted 28th February 2013

DOI: 10.1039/c3ra40510a

www.rsc.org/advances

Debasish Mohanty,*a Sergiy Kalnaus,*a Roberta A. Meisner,a Athena S. Safat,a

Jianlin Li,b E. Andrew Payzant,c Kevin Rhodes,d David L. Wood, IIIa and Claus Danielbe

A decrease in the c-lattice parameter was observed in Li1.2Co0.1Mn0.55Ni0.15O2 during constant voltage

holding at 4.5 V by in situ X-ray diffraction. Comparison of magnetic susceptibility data before and after

high-voltage hold reveals the change in average oxidation states of transition metal ions during high-

voltage holding process. Transmission electron microscopy studies show the spinel reflections with

fundamental trigonal spots from the particles after high-voltage hold indicating substantial structural

modification. The structural transformation was believed to occur due to the oxygen release and/or the

migration of transition metal cations to lithium layer during constant voltage holding.

Introduction

Rechargeable lithium-ion battery (LIB) technology is theleading candidate to power electric vehicles (EV) because ofits high stored energy density, light weight, low maintenance,long service-life, and high efficiency as compared to internalcombustion (IC) engines.1–4 However, in order to realize thecomplete electrification in transportation systems, the perfor-mance of LIBs needs to be improved while maintainingmaximum life and safety.1 Recently, Li-rich materials,Li1+yM12yO2 (M = Co, Mn, Ni) (Li-rich NMC hereafter) revealspromising high capacity when operated at high voltage.5–7

Structurally integrated Li2MnO3-stabilized ‘layered-layered’xLi2MnO3?(1 2 x)LiMO2 (M = Mn, Ni, Co), (for example,0.5Li2MnO3?0.5LiNi0.27Mn0.27Co0.27O2 or Li1.2Co0.1Mn0.55Ni0.15O2

investigated in the present study) has shown improved electro-chemical performance compared to stoichiometric LiMO2 (M =Co, Mn, Ni) cathodes.8–11 Most importantly, these structuresdeliver promising high capacity between 200–250 mAhg21

within the voltage window of 2.0–4.8 V (vs. Li/Li+) which makesthem excellent candidates for LIBs in EV applications.6,8 Thecrystal structure of Li-rich NMC compounds is derived from the

layered LiMO2 with a-NaFeO2 structure with trigonal symmetry(rhombohedral or hexagonal unit cell, R3̄m space group) and theexcess lithium ion present in the structure occupies thetransition-metal layer, filling all the octahedral sites of cubic-close-packed (ccp) oxygen arrays.6 The presence of lithium ionsin the transition-metal (TM) layer generates cation-orderingbetween TM ions, may lead to the formation of Li2MnO3 phase(monoclinic unit cell, C2/m space group).3 The presence ofcation-ordering can be detected from X-ray diffraction (XRD)pattern, where the superlattice peaks appear at lower 2h angles(22u–26u with Cu-Ka radiation). However, LiMO2 and Li2MnO3

structures have similar ccp layers with interlayer spacing ofy4.7 Å for (001) of layered monoclinic and (003) of layeredtrigonal (Fig. 1) phases, which allows perfect integration of these

aMaterials Science and Technology Division, Oak Ridge National Laboratory, Oak

Ridge, Tennessee 37931, USA. E-mail: mohantyd@ornl.gov; kalnauss@ornl.govbEnergy and Transportation Science Division, Oak Ridge National Laboratory, Oak

Ridge, TN, 37831, USAcChemical and Engineering Materials Division, Oak Ridge National Laboratory, Oak

Ridge, Tennessee- 37831dFord Research and Innovation Center, Ford Motor Company, Dearborn, Michigan

48121, USAeBredesen Centre for Interdisciplinary Research and Graduate Education, University

of Tennessee, Knoxville, Tennessee 37996, USA

3 Electronic supplementary information (ESI) available: SEM image, EDS data,electrochemical plots. See DOI: 10.1039/c3ra40510a

Fig. 1 Crystal model showing (001)M and (003)R planes having similar ccp layers.R refers to rhombohedral (trigonal) crystal structure and M refers to monocliniccrystal structure, and TM refers to transition metal ions (Co, Mn, and Ni). Mnatom in monoclinic unit cell (MnM) is shown in purple color. The rhombohedraunit cell and monoclinic unit cells were drawn separately and overlaid with eachother.

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structures at an atomic level to form a layered-layeredLi1+yM12yO2 (M = Co, Mn, Ni) integrated structure.5

During electrochemical charging up to 4.4 V, the lithiumions are extracted from the lithium layers of the parentstructure without destroying the cation-ordering in the TMlayers. However, the Li2MnO3 phase would be activated whenelectrochemical delithiation occurs above 4.4 V.8–10 It has beenreported that an irreversible chemical reaction represented byLi2MnO3 A Li2O + MnO2 occurred during first cycle chargingabove 4.4 V, accompanying a pronounced irreversible capacityloss.8,12 Recently, the structural degradation in these cathodesrelating to the capacity and/or voltage fade8,13–14 duringelectrochemical cycling have been major issues in thesehigh-capacity cathodes.8,17 The fundamental understandingon this structural degradation process in Li-rich NMCmaterials is still unclear and hence, there is a tremendouslygrowing interest in investigating the structure–electrochemicalproperty correlations in these cathode oxides. However, thesereports are solely focused on the structural transformationduring constant current charge/discharge steps9,15,11,13–14,16–17

(voltage of the battery changes) where, the lithium ionscontinuously shuttle between the cathode and anode. Thestructural transformation in Li-rich NMC could also occurwhile holding the cell at constant voltage which has not beeninvestigated yet and is the focus of this present work. Themotivation behind this work is: 1) to investigate the structuraltransformation and 2) to obtain fundamental understandingon the stability of this high voltage and high capacity Li-richNMC cathodes when holding the cell at steady voltage.

In this report, we document the structural evolution in Li-and Mn-rich Li1.2Co0.1Mn0.55Ni0.15O2 cathode while holdingfor prolonged duration in a delithiated state (4.5 V), which wasevidenced by in situ XRD. Furthermore, the magnetic suscept-ibility and transmission electron microscopy (TEM) studiesbefore and after high voltage hold provide the insights into themechanism of phase transformation that might occur duringthe high voltage holding step.

Experimental

The electrode containing Li1.2Co0.1Mn0.55Ni0.15O2 (TODAHE5050) and its pristine powder material were purchasedand used for this study. The material was synthesized by co-precipitation using hydroxide precursor by TODA America Inc.Subsequently, electrodes were fabricated at the U.S.Department of Energy’s (DOE) Cell Fabrication Facility,Argonne National Laboratory. The electrode composition wasas follows: 86 wt% TODA HE5050, 8 wt% Solvay 5130 PVDFbinder, 4 wt% SFG-6 graphite (Timcal), and 2 wt% Super P1Carbon Black (Timcal) and drying and calendering at 0.5 mmin21 resulted in 37% porosity cathodes with 35 mm thickcoatings.

In situ XRD cell was fabricated on 2032 coin cell hardwarebased on the modifying the procedure described by Rhodeset al.18 and more information on this modified cell can be

found in our previous report.19 Briefly, the in situ cell for thisexperiment was done by punching a hole (y13 mm diameter)in the regular coin cell case and then sealed with Kapton1film (Ø19 mm disk), and vacuum sealing epoxy (Torr Seal1)was applied along the edge of the window to preventelectrolyte leakage. Comparison of charge-discharge curvesfrom windowed and non-windowed cells showed no signifi-cant differences, so the Kapton windows may be regarded asgas-tight, and the in situ XRD results as representative of thecell behavior. The prepared coin cell case was transferred to anargon filled glove box for coin cell assembly withLi1.2Co0.1Mn0.55Ni0.15O2 (Ø12 mm disk) as the cathode andlithium foil (Ø12 mm disk ,Alfa Aesar, 0.75 mm thick, 99.9%)as the anode, respectively. 1.2 M solution of LiPF6 in EC : EMC(3 : 7 wt) was used as electrolyte (Purolyte). Celgard 2325separator (16 mm diameter) was used between the electrodesto prevent electrical shorts inside the battery.

The electrochemical cycling was performed by using aBioLogic SP-200 potentiostat controlled with the EC-Labsoftware (V. 10.02). X-ray patterns were collected simulta-neously using PANalytical X’Pert Pro system with molybdenumsource (l = 0.71073 Å) and automatic divergence and anti-scatter slits operated at 60 kV and 45 mA current. The charge/discharge process was carried out between 3 and 4.5 V vs. Li/Li+ at 20 mA g21. After finishing the second cycle charging, thecell was held for 40 h so that current dropped to zero andbecame insignificant in order to avail no charging effect in thecell. XRD patterns were collected with 2h angle range of 5–40u.The collected in situ XRD patterns were refined by Rietveldmethod using X-pert HighScore Plus with zero sampledisplacement errors and available background correction.For the refinement, the unit cell of layered LiMO2 phase withR3̄m space group and Li2MnO3 phase with C2/m space groupwere considered and the detailed crystallographic informationof these unit cells can be found in these references.20,21 Thelattice constants were determined by least squares refine-ments.

The scanning electron microscopy (SEM) image and energydispersive X-rays (EDS) were obtained by Hitachi S4800 FEG-SEM at 20 kV. The ex situ powder XRD patterns were collectedon the PANalytical X’Pert Pro system, the same instrumentwhich was used for in situ XRD study. The condition of datacollection was similar to the condition for in situ XRDexperiment. TEM was performed using a Hitachi HF3300TEM at 300 kV. For pristine material, the Li-rich NMC powderwas dispersed in ethanol and a few drops were added to aholey TEM copper grid. For cycled material (before and afterhigh-voltage holding) the cell was disassembled in the Argonfilled glove box and the cathode was obtained and washedseveral times by di methyl carbonate (DMC). The cathode waskept in the glove box overnight to dry. The powder wasremoved from the electrode, ground, dispersed in ethanol, andsonicated for 10–15 min. Few drops of the solution were addedto the holey Cu grid. Selected area electron diffraction for thestarting powder was simulated by WebEMAPS4 using theLiCoO2 structure and Li2MnO3 crystallographic information.

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DC magnetization was measured as a function of temperature,using a Quantum Design Magnetic Property MeasurementSystem. Each sample was first cooled to 5 K in zero field, thena field of 100 Oe was applied and the data were collected from5 K to 320 K (zfc). The sample was also cooled in the appliedfield from 320 K down to 5 K, while measuring magnetization(fc). Magnetization data were collected directly from theelectrodes after cycling, and the active material (TODAHE5050) was not removed from the metal foil. The cycledelectrodes were washed in DMC to remove any remainingelectrolyte and decomposition products. There was nosignificant difference in the weight of the washed cycledelectrodes compared to the as-received electrodes. Theeffective magnetic moments were calculated from the molarmass of the active material (TODA HE5050).

Results and discussion

Characterization of pristine material

Particle morphology and composition of as receivedLi1.2Co0.1Mn0.55Ni0.15O2 pristine material were identified byscanning electron microscopy (SEM) (Fig. S1a in theSupplemental Information3). The SEM image shows theplatelet like morphology which is very typical for NMCmaterials.22 The EDS (Fig. S1b3) shows the Co: Mn: Ni ratioas 0.11 : 0.54 : 0.15 which confirms the material compositionas Li1.2Co0.1Mn0.55Ni0.15O2. The ex situ powder X-ray diffrac-tion (XRD) collected from pristine Li1.2Co0.1Mn0.55Ni0.15O2

(Fig. S23) shows that the parent compound is a combination ofrhombohedral (trigonal) and monoclinic phases. The majorpeaks in Fig. S23 could be assigned to the trigonal phase (R3̄mspace group) and/or monoclinic phase (C2/m space group).Due to similar d-spacing of these two phases, it is difficult todistinguish these two structures from the XRD. However, thepresence of cation-ordering between Li+ and Mn4+ possiblyoriginating from the Li2MnO3 phase, was confirmed by thesuperlattice peaks (020) and (110) (Fig. S23) present at lower 2h

angle. Moreover, the high-resolution transmission TEM(Fig. 2a) shows the lattice fringes spacing of y4.7 Å indicating

the presence of (001) planes and/or (003) planes originatingfrom the C2/m and R3̄m space group respectively. The selectedarea electron diffraction (SAED) obtained from one represen-tative particle is presented in Fig. 2b. Bright fundamental{112̄0} reflections can be clearly seen which originates fromtrigonal phase with R3̄m space group and the superlatticereflection (marked in a dotted circle for an example) denotesthe presence of monoclinic symmetry with C2/m space group.The XRD, TEM and SAED pattern confirms the pristineLi1.2Co0.1Mn0.55Ni0.15O2 cathode is a trigonal phase as a hostmatrix with embedded monoclinic phase.

In situ XRD study

In situ electrochemical XRD patterns of Li1.2Co0.1Mn0.55Ni0.15O2

cathode were collected by a custom designed half-cell (coin cell)as described in the experimental section between 3.0–4.5 V (vs.Li/Li+) at 20 mA g21 (equivalent to C/16 rate). The cell was heldat 4.5 V for 40 h at the end of the second charging cycle. Thereason for performing the high voltage hold experiment aftersecond charging cycle is to avoid the complexity that might arisedue to structural transformation in the first cycle plateau region(4.5 V–4.7 V). The in situ XRD plots collected during charge(denoted as I in Fig. 3), discharge (denoted as II in Fig. 3), andduring holding at 4.5 V are shown in Fig. 3. The major peaks inFig. 3 were assigned to host rhombohedral (trigonal) phase(R3̄m space group). The reflection denoted as * and # are fromstainless steel (cell cap) and carbon black (from electrode),respectively. The monoclinic superlattice reflections (020) and(110) are not visible in the in situ XRD plots because of lowintensity of these reflections and/or due to high background atlow 2-theta angle in the XRD patterns. Since the (003) positionalways determines the trigonal unit cell dimension along thec-axis, and the d-spacing of this plane always varies as a functionof the charge/discharge process, which is related to lithiumextraction/insertion from/into the cathode,23 the displacement

Fig. 2 High-resolution TEM image (a) and selected area electron diffractionpattern (b) collected from [0001] zone axis from Li1.2Co0.1Mn0.55Ni0.15O2 crystal.It shows fundamental reflections that originate from the rhombohedral(trigonal) unit cell. The super lattice spots (marked with a dotted circle) originatefrom the cation ordering (Li+ and Mn4+) in the TM layers.

Fig. 3 In situ XRD pattern collected from Li1.2Co0.1Mn0.55Ni0.15O2 cathodeduring charge (I), discharge (II) and hold at 4.5 V for 40 h. The reflection fromstainless steel (cell cap) and carbon (from electrode) are designated as * and #,respectively. The rectangular marked region was presented in the expandedview in Fig. 4.

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of (003) peak position during charge, discharge and highvoltage hold was considered for further analysis. The enlargedview of (003) peak position during charge/discharge andholding at 4.5 V is presented in Fig. 4. The experimental latticeparameters calculated from in situ XRD patterns were plottedalong with the electrochemical profile in Fig. 5. During chargingup to 4.5 V, lithium ions were being extracted from the lithiumlayer, generating lithium deficiency in the lithium layer andinducing strong electrostatic repulsion among ccp oxygenlayers. Consequently, the unit cell expanded along the c-axisand the c-lattice parameter increased (Fig. 5). In contrast, thea-lattice parameter decreased during charging due to theformation of transition metal ions with smaller ionic radii,which resulted from oxidation of Ni2+ (0.69 Å) and Co3+ (0.61 Å)to Ni3+ (0.56 Å)/Ni4+ (0.48 Å) and Co4+ (0.53 Å), respectively in anMO6 environment24 for charge compensation due to Li+

extraction. During discharge, the opposite trend in c-anda-lattice parameter values indicate the successful lithiuminsertion into the host matrix and consequently, reduction ofNi4+/Co4+ to Ni2+/Co3+ ions happens. The increase (decrease) ofc-lattice parameter was in agreement with the displacement of(003) peak position to lower (higher) 2h angle during lithiumion extraction (insertion) as shown in Fig. 4 and 5.

Contrary to the shifting to lower 2h during the charge step,the (003) peak position shifted towards higher 2h angle (Fig. 4and Fig. S33 for isoplots) during the constant voltage hold at4.5 V indicating the contraction of the unit cell along c-axis.This was further confirmed from the calculated c-latticeparameter values (Fig. 5) which were found to decrease duringprolonged holding time. The c-lattice parameter value wasreduced from 14.421 ¡ 0.002 Å at the beginning of holdingstep to 14.372 ¡ 0.003 Å after 40 h holding at 4.5 V. The latticeparameter of 14.370 Å corresponds to the value obtained at 4.2V during the preceding charging/discharging process. Anincrease in a-lattice parameter was also observed during thehigh voltage hold. However, the change was insignificant

compared to that of the c-parameter. One can argue thedecrease in c-lattice parameter during holding the cell atsteady voltage in the context of relaxation and/or equilibriumprocess that might occur in the cathode structure. However,the relaxation process may happen when the cell is kept in rest(no external current applied) instead of keeping at constantvoltage. In the former case, when external current is zero, thevoltage of the cell decreases/increases until the structureattains the lowest energy state and becomes stable. If therelaxation and/or equilibrium process happens, one would notexpect the new phase and/or change in electronic states of TMions in the host cathode structure however, which wasobserved in this study (see the TEM and magnetization studybelow). In this present study the voltage of the cell was keptconstant for 40 h after charging the cell to 4.5 V (second cycle).The voltage of the cell is directly related to the difference in thechemical potential between cathode and anode and when it isconstant, there is a reduced likelihood that lithium ionsshuttle between the electrodes after 40 h. Therefore, thechange in lattice parameter during high-voltage hold may notbe a simple relaxation process instead; these are believed to beassociated with the structural rearrangement inside theLi1.22yCo0.1Mn0.55Ni0.15O2 (y , 1) cathode, as confirmed byTEM and magnetic susceptibility studies (see below). Duringthe course of charging to 4.5 V, 45% of the lithium ions havebeen extracted (calculated based on the capacity from Fig. S43)mostly from the lithium layer creating many vacancies. It isreasonable to assume that most of the lithium ions in the

Fig. 4 Displacement of (003) peak of Li1.2Co0.1Mn0.55Ni0.15O2 during chargingto 4.5 V (I), discharge to 3 V (II), and hold at 4.5 V.

Fig. 5 Change in c- and a-lattice parameters (uncertainty of values are plottedas error bars)) of Li1.2Co0.1Mn0.55Ni0.15O2 during electrochemical cycling andhold at 4.5 V. The point A and B represents the region where ex situ TEM andMagnetization susceptibility experiments were performed.

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transition metal layer remain during charging since extractionof lithium ions from the transition metal layers requireshigher energy (¢4.4 V) compared to that from the lithiumlayer. Hence, it is more likely that the activation of Li2MnO3

phase did not occur to a significant extent during the first orsecond cycle charging step. However, the shifting of (003)peaks towards higher 2h angle (change in lattice parameters)provides an indication that the lithium layer vacancy is filledwith foreign migrated cations, which eventually reduces theelectrostatic repulsion between ccp oxygen layers and/orintroduction of new phases in the host trigonal unit cell. Toclarify this speculation, the high-resolution TEM along withFast Fourier Transformation (FFT) and magnetic susceptibilitymeasurement before (marked as A in Fig. 5) and after highvoltage hold (marked as B in Fig. 5) at 4.5 V were obtained andpresented in the following section.

TEM and magnetic susceptibility study

In Fig. 6a the high-resolution TEM and corresponding FFT(inset) from a representative particle from the material beforeholding at 4.5 V was presented. The SAED shows thefundamental trigonal reflections {112̄0} reflections originatesfrom trigonal phase with R3̄m space group. Due to the low

intensity, the superlattice reflections from monoclinic phasewith C2/m space group were not visible in the FFT pattern.However, based on previous reports it is reasonable tospeculate that the cation ordering may have been decreasedup to certain extent, when charging up to 4.5 V since activationof Li2MnO3 starts while charging beyond 4.4 V.7 No strongspinel phase in the material is observed from examination ofseveral particles before high voltage hold in the second cycle.The high-resolution TEM image collected from a representa-tive particle from the material after holding the cell at 4.5 V for40 h is given in Fig. 6b. The FFT generated from this imageshows the appearance of new spots where the presence of twophases were observed. These different diffraction spots can beattributed to different phases as further explained. Thefundamental {112̄0} reflections were from the host trigonalphase (rhombohedral/trigonal, often designated as O3 phasein R3̄m space group). The presence of a spinel-like phase(Fd3̄m space group) in the crystal was confirmed by the strongsuperlattice reflections that appeared half-way between twofundamental trigonal reflections. In case of spinel phase, thetransition metal ions and lithium ions occupy in the layers ofthe interstitial sites with an alternating ratio of 1 : 3 that canbe designated a 16d and 8a lattice sites in a Fd3̄m space groupand the superlattice reflection appears in the middle of twostrong fundamental O3 spots.25 The other set of reflectionsthat appeared in the middle of trigonal symmetry of funda-mental reflections was denoted as (101̄0) planes. Thesediffraction spots are normally forbidden in trigonal symmetryand are an indication of the presence of transition metal ionsin the lithium layer.26,27 Therefore, the TEM results confirmthat major structural modification occur during the highvoltage hold process. The absent spinel phase in the materialat the end of second cycle charging (before hold) is in contrastto the results reported by Ito et al.,14 where the authors findfaint spinel reflections (two-times-periodicity according to thenotation made by those authors) even during charging in the4.5 V plateau region. This may be due to the differentstructures during charging at plateau region (4.5 V) in thefirst cycle compared to the structure during charge at 4.5 V inthe second cycle. Detailed statistical analysis of structure(SAED pattern) from particles at these two points (delithiatedmaterial at 4.5 V in the first cycle and second cycle) should bemade in order to understand this behaviour, which is beyondthe scope of this present work. However, the magneticsusceptibility data presented before and after high voltagehold at 4.5 V clearly describes the structural rearrangementthat happens during holding the cell at steady high-voltagewhich is explained as follows. We would like to mention that,the discussion on magnetic susceptibility on starting materialis out of the scope of this article and hence, we only confineour discussion to the magnetic susceptibility before and afterhigh-voltage hold at 4.5 V.

The variation of inverse molar magnetic susceptibility (1/x;where x = M/H, M is magnetic moment and H is appliedmagnetic field) vs. temperature taken before (point A in Fig. 5)and after high voltage hold (point B in Fig. 5)) is presented inFig. 7 and in the inset we present the change magneticsusceptibility (x) with the temperature (T). From Fig. 7, it isclear that both curves show paramagnetic behaviour obeying

Fig. 6 HRTEM image from Li1.2Co0.1Mn0.55Ni0.15O2 taken before (a) and afterholding (b) at 4.5 V for 40 h. The FFT is shown as insets of a and b represents thesample at point A and B marked in Fig. 5.

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Curie–Weiss law at higher temperature (T > 100 K), differingthe slopes of the curves. The experimental effective magneticmoments were deduced from the plot of inverse molarmagnetic susceptibility vs. temperature between the regionsof 100 K and 320 K using the equation xm = Cm/(T 2 h) as alinear fit. Here, xm is the molar magnetic susceptibility, Cm isthe Curie constant, and h is the Weiss temperature. The valuewas found to be 2.66 ¡ 0.04 mB and 2.75 ¡ 0.02 mB for thematerials before and after high voltage hold respectively whichshows an increase in effective magnetic moment after high-voltage holding process. In order to examine the validity ofincrease in effective magnetic moment values, magnetizationdata were collected from multiple samples (three samplesfrom each point A and B in Fig. 5) and effective magneticmoments were calculated from each sample. The averageeffective magnetic moment values were 2.65 ¡ 0.02 mB and2.76 ¡ 0.03 mB for samples before (point A in Fig. 5) and after(point B in Fig. 5) high-voltage hold respectively; whichconfirms the increase in effective magnetic moment in thesamples after high-voltage hold. The increase in experimentaleffective magnetic moment confirms the change in averageoxidation states of TM ions during constant voltage holdindicating the continued redox reaction during constantvoltage hold. This further indicates the introduction ofmagnetic ions with higher effective magnetic moment suchas Ni2+ (LS/HS, S = 1), Ni3+(HS, S = 3/2 or LS, S = 1/2), Co3+ (HS,S = 2) due to oxidation/reduction of Ni4+ (LS, S = 0), Co4+ (LS, S= 3/2) present in the charged material (delithiated) beforeprolonged hold. When voltage is constant, the difference inthe chemical potential between cathode and anode isconstant; hence, one would not expect the extraction andinsertion of lithium ion between the electrodes until 40 h (inthis study). In a LIB, charge/discharge process involves Li+

extraction/insertion from/into the cathode structure and theaverage oxidation states of TM ions changes to maintain the

charge neutrality in the system. However, when no net lithiumions travelling between the electrodes after a long period ofconstant voltage holding, the charge compensation for thechange in average oxidation states of transition metal ionsmust have been preceded by oxygen release from the cathodestructure. In the following section, we will explain the detailedmechanism which might be responsible for this structuraldegradation during high-voltage holding process.

The pristine Li1.2Co0.1Mn0.55Ni0.15O2 material is an integra-tion of layered trigonal and monoclinic phases where lithiumions are present in the lithium layer (LiLi) as well as in thetransition metal layer (LiTM). Upon charging to 4.5 V, the LiLi

ions are being extracted expanding the unit cell along c-axis,hence increasing the c-lattice parameter. However, duringhigh-voltage hold, the lattice contracts along c-axis. This mightbe due to following reasons. At high the voltage (delithiated)state, lithium deficient Li1.22yCo0.1Mn0.55Ni0.15O2 is not stablesince the Fermi level is situated within the oxygen valenceband by dropping the Fermi energy into the top of the O22:2pband.28–30 This may lead to release of oxygen from LiTMO6 and/or TMTMO6. Hence 1) the release of oxygen from LiTMO6 maydrive the LiTM to migrate from TM layers to the lithium layerfilling the octahedra vacancy. and/or 2) release of oxygen fromTMTMO6 may also drive some of the TM ions to lithium layer tofill the vacancy in the lithium layer. The oxygen loss and thefilling of vacancy in the lithium layer by foreign ions couldreduce the electrostatic repulsion between ccp oxygen layerscausing the contraction of unit cell along c-axis.28 Thehypothesis of migration of TM ions in to the lithium layer isconfirmed by the forbidden reflections and spinel phase fromthe FFT (Fig. 6). For example, Ni2+, which might have beenformed during oxygen release during high voltage hold, hassimilar ionic radius as Li+ and could migrate into Li vacancysites. However, 0.15 mol of Ni2+ in the starting material couldhave been oxidized to Ni4+ when charged to 4.5 V. Hence, thereduction of Ni from +4 to +2 state, suggested by the increasein magnetic susceptibility, must be accompanied by oxygenrelease for charge compensation. This is also in agreementwith the increase in a-lattice parameter (which is related toaverage metal-metal distance) in this study (Fig. 4) andsupports the argument of decrease in average oxidation statesof TM ions that might occur due to reduction of Ni4+ (S = 0) toNi2+ (S = 1). Eventually, the presence of Ni2+ in the lithiumlayer can also create a ferromagnetic (FM) interaction due to180 degree FM exchange along Ni2+ (in the lithium layer) -O-Mn4+ (in the TM layer).24,27,31 The Curie–Weiss temperaturevalues (h) obtained from the Curie–Weiss plot was found to be273 K and 265 K for the material before and after high voltagehold respectively. The decrease in the negative character in theh value in the material after high voltage hold also supportsthe fact that some degree of FM exchange has been introduceddue to magnetic exchange coupling of Ni2+ in the lithium layerwith the Mn4+ in the TM layer via oxygen ion.27

The phase transformation in Li and Mn rich cathodematerials is not a unique behaviour. However, most of thereports show that the ‘‘layer to spinel phase’’ transformationin the cathode during repeated charge/discharge cycles whereLi ions are continuously extracted from and inserted into thecathode structure. The present study shows that the structural

Fig. 7 Temperature dependence of inverse magnetic susceptibility (zfc), inapplied field of 100 Oe, for the material before and after high voltage hold for40 h that is depicted in Fig. 5.The black solid lines shows the Curie–Weiss fit (R2 =0.999). The inset shows the temperature dependence of magnetic susceptibilityin fc and zfc mode.

7484 | RSC Adv., 2013, 3, 7479–7485 This journal is � The Royal Society of Chemistry 2013

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transformation also occurs during constant voltage hold. Thestructural transformation at the constant voltage (4.5 V) maybe preceded by the oxygen release followed by the reduction ofNi4+ to Ni2+ for charge compensation. The filling of lithiumlayer vacancy by Ni2+ and/or spinel phase formation are alsoenvisioned which might contract the c-unit axis. However, theextent of Ni2+ ions formation and their migration to lithiumlayer is not quantified. The verification of this structuralevolution mechanism and quantitative analysis of Ni2+

formations is under further investigation by neutron scatter-ing experiments which are very sensitive to the site occupancyfactor (SOF) of oxygen and lithium ion as well as TM ions andwill be presented in future publications.

Conclusions

Evidence of a phase transformation in a Li-richLi1.2Co0.1Mn0.55Ni0.15O2 cathode during constant voltage hold-ing at 4.5 V is reported. Contrary to the constant currentcharge step, the c-lattice parameter was found to decreaseduring constant voltage holding, which demonstrates thestructural instability of this cathode in a highly delithiatedstate. The results from high-resolution TEM show the spineland forbidden reflections from the particle that has under-gone long-term, high-voltage holding process indicates themigration of transition metal ions from the transition metallayer to the lithium layer modifying the crystal structure. Themagnetic susceptibility results show the change in electronicstates of the transition metal ions resulting from the chargecompensation due to oxygen release. These results shed lighton the important structural transformations that occur in Li-rich cathode materials such as Toda HE5050 during high-voltage hold and can be an instrument for designingstructures with superior structural stability.

Acknowledgements

This research at Oak Ridge National Laboratory, managed byUT Battelle, LLC, for the U.S. Department of Energy (DOE)under contract DE-AC05-00OR22725, was sponsored by theOffice of Energy Efficiency and Renewable Energy for theVehicle Technologies Office’s Applied Battery ResearchProgram (Program Managers: Peter Faguy and David Howell).Part of this research was supported by the DOE, Basic EnergySciences, Materials Sciences and Engineering Division and byORNL’s ShaRE User Facility, which is sponsored by theScientific User Facilities Division, Office of Basic EnergySciences, DOE. The electrodes were produced at the DOE’sCell Fabrication Facility, Argonne National Laboratory, byAndrew Jansen and Bryant Polzin. The Cell Fabrication Facilityis fully supported by the DOE Vehicle Technologies Office(VTO) within the core funding of the Applied Battery Research(ABR). The authors thank Dr. Daniel Abraham at ArgonneNational Laboratory for useful discussion.

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