Journal of Alloys and Compounds 732 (2018) 385e395
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Journal of Alloys and Compounds
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Phase transitions and related electrochemical performances of Li-Richlayered cathode materials for high-energy lithium ion batteries
Jianqing Zhao a, b, d, Xiaoxiao Kuai a, b, Xinyu Dong a, b, Haibo Wang a, b, e, Wei Zhao f,Lijun Gao a, b, *, Ying Wang d, **, Ruiming Huang c, ***
a Soochow Institute for Energy and Materials InnovationS, College of Physics, Optoelectronics and Energy & Collaborative Innovation Center of Suzhou NanoScience and Technology, Soochow University, Suzhou, 215006, Chinab Jiangsu Provincial Key Laboratory for Advanced Carbon Materials and Wearable Energy Technologies, Suzhou, 215006, Chinac Department of Chemistry, Rutgers-Newark, The State University of New Jersey, Newark, NJ 07103, United Statesd Department of Mechanical & Industrial Engineering, Louisiana State University, Baton Rouge, LA 70803, United Statese Institute of Chemical Power Sources, Soochow University, Zhangjiagang, 215600, Chinaf Shanghai Haiying Machinery Plant, Shanghai, 200436, China
a r t i c l e i n f o
Article history:Received 7 July 2017Received in revised form17 October 2017Accepted 22 October 2017Available online 24 October 2017
Keywords:Chemical activationIon exchangePhase transitionLi-rich layered cathode materialLithium ion battery
* Corresponding author. Soochow Institute for EnerCollege of Physics, Optoelectronics and Energy & CollaSuzhou Nano Science and Technology, Soochow Unive** Corresponding author.*** Corresponding author.
The present work systematically probes and tracks the phase transition of Li-rich layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2 (marked as LMNCO) by using an ex-situ chemical activation that is realizedthrough ion-exchange and post-annealing processes, in order to understand related electrochemicalperformances of Li-rich cathode materials for advanced lithium-ion batteries. Ion exchanges of Hþ-Liþ
and subsequent TBAþ-Hþ (TBA: tetrabutylammonium) in LMNCO are carried out, resulting in its layered-to-spinel phase transition after optimal heat treatments. The resultant compound shows a Li4Mn5O12-type spinel structure. This converted spinel cathode material can deliver discharge capacities higher than300 mAh/g at 0.1 C and 200 mAh/g at 1 C (1 C ¼ 250 mA/g), respectively, and also exhibits better cyclingstability and rate capability in comparison with pristine layered LMNCO and other derivatives. This workoffers a feasible route to study all changes of morphologies, crystal structures, chemical compositions,surface areas and related electrochemical lithium storage behaviors during phase transitions of Li-richlayered cathode materials, and thus provides insights on optimizing electrochemical performances forhigh-energy and high-power lithium ion batteries.
The rechargeable lithium ion batteries have been demonstratedas highly effective power supplies for electric transportation systemand portable electronic devices. Performances of lithium-ion bat-teries crucially rely on energy and power densities of electrodematerials [1]. Recently, tremendous research efforts focus ondeveloping advanced cathode materials, which are expected tooffer high specific capacity and operating voltage together with
gy and Materials InnovationS,borative Innovation Center ofrsity, Suzhou, 215006, China.
outstanding cycling stability and rate capability [2e4]. The Mn-based Li-rich layered oxides have attracted tremendous researchefforts owing to the high lithium storage capacity and workingpotential. These cathode materials marked as Li[LixMnyMz]O2(M¼ Co andNi; xþ yþ z¼ 1 and y > 0.5) can be cycled over a broadvoltage range of 2.0e4.8 V vs. Liþ/Li and deliver specific capacitieshigher than 250mAh/g, along with other merits including low cost,environmental friendliness and safety [5e18].
Li[Li0.2Mn0.54Ni0.13Co0.13]O2 (marked as LMNCO) belongs toaforementioned Li-rich and Mn-rich category, which has thedesirable theoretical capacity (>300 mAh/g) and high workingvoltage (~4.0 V vs. Liþ/Li) [18]. As reported in literature [6e8], Li-rich layered LMNCO is the product with a structural intergrowthof layered lithium-inactive Li2MnO3 (space group C2/m) andlayered lithium-active LiMn1/3Ni1/3Co1/3O2 (space group R-3m) at amolar ratio of 1:1 (0.5Li2MnO3$0.5LiMn1/3Ni1/3Co1/3O2). The highcapacity of LMNCO cathode material can be achieved through the
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electrochemical activation of Li2MnO3 component in the firstcharge reaction above 4.5 V vs. Liþ/Li. However, such a reactionleads to significantly irreversible capacity loss and low Coulombicefficiency in the first charge/discharge cycle, and further can triggera detrimental layered-to-spinel phase transition during next elec-trochemical cycling of activated layered LMNCO. The structuralsimilarity of cubic close-packed oxygen arrays in layered and spinelconfigurations facilitates the layered-to-spinel phase transition inprinciple [7,9,13,19,20]. As a result, transition metal ions migrate tolithium layers in LMNCO and reside on vacant lithium ion sitespermanently. This irreversible phase transformation is continueduntil a hybrid layered-spinel composite material is formed.Accordingly, the voltage plateau of working cathode is reducedfrom ~4.0 V (layered) to ~3.0 V (spinel) vs. Liþ/Li [8,21]. Overall,electrochemical activation of Li2MnO3 component causes structuralinstability and unfavorable phase transition, which accounts for thevoltage fading and decreased energy density of Li-rich layeredcathode materials [9].
On the other hand, it has been reported that the layered-to-spinel phase transformation in Li-rich layered cathode materialsdemonstrates the unexpectedly high-rate capability of hybridlayered and spinel cathode materials [7,8,22], because the spinelphase is lithium-active with enhanced electronic conductivity andlithium ion diffusivity [7]. The formation of the spinel phase in Li-rich layered cathode materials has been demonstrated by highresolution STEM observations [21], high resolution TEM imageswith selected area electron diffraction (SAED) patterns [7], in-situX-ray diffraction patterns [13], X-ray absorption spectroscopic andRaman studies [15,23], and is also reflected on charge/dischargecurves [22], differential capacity plots [14], and cyclic voltammetric(CV) profiles [7,24]. The intergrowth of spinel-layered phases tendsto alleviate the electrochemical inferiority of Li-rich layered cath-ode materials. However, structural details of these phases in thecycled electrodes have not been comprehensively understood yet.The crystal phase of the spinel formed in Li-rich layered oxides isalways reported as either “spinel” or “spinel-like” phase. The effectsof such a phase transformation (whether to improve or deteriorateperformances of spinel-layered composite cathodes) are still underdebate. It is important to explore these effects, in order to under-stand the fundamental electrochemical behavior of high-capacityLi-rich layered cathode materials.
As reported in literature [16,25], Li-rich layered cathode mate-rials can be chemically activated via the protonation (Hþ-Liþ ex-change) in an acidic environment, followed by removing Hþ ions ina post-annealing treatment in air. As a result, the activated cathodematerial show considerably increased capacity and associatedCoulombic efficiency in the initial cycle. Although Hþ-Liþ ion ex-change can induce the formation of the spinel phase, the formationof spinel domains in the structure of Li-rich layered cathode ma-terial is very limited. Consequently, the phase transition succes-sively takes place within the hybrid spinel-layered compositecathode during next electrochemical cycles, leading to structuralinstability and poor cycling stability [25]. As referred to other ion-exchange reports [26e29], alkylammonium hydroxides, such astetrabutylammonium hydroxide (TBAþ$OH�) and tetramethy-lammonium hydroxide (TMAþ$OH�), have been widely employedto exfoliate protonated layered materials into two-dimensional(2D) nanosheets. Due to the organic characteristics and largermolecular size of alkylammonium cations in comparison withprotons and lithium ions, TBAþ$OH� can be utilized for the secondion exchange (TBAþ-Hþ) of protonated Li-rich layered cathodematerials, in order to realize a complete layered-to-spinel phaseconversion.
In our previous work [30], we realized the layered-to-spinelphase transformation of Li-rich Li[Li0.2Mn0.54Ni0.13Co0.13]O2 by
employing ex-situ ion-exchange and post-annealing processes, andfound that the completely-converted material shows a Li4Mn5O12-type spinel structure rather than commonly-reported LiMn2O4-type spinel. The approach we developed not only allows thecomprehensive study of electrochemical effects resulting from thegrowth of a spinel phase within Li-rich layered cathode materials,but also offers a feasible route to precisely identify different crystalstructures during the formation of newly-formed spinel phase.Herein, we report a more comprehensive study during the phasetransition of Li-rich layered materials with aspects to all changes ofmorphologies, crystal structures, chemical compositions, surfaceareas and related electrochemical lithium storage behaviors indetails. Electrochemical performances indicate that introduction ofa spinel phase significantly increases the specific capacity of~100 mAh/g and results in much better high-rate performances ascompared with original Li-rich layered cathode materials, but re-duces the working voltage from 4.0 V to 3.0 V due to the activationof Mn3þ/Mn4þ redox pair. It is also interesting to find that Liþ-TBAþ
can be carried out in the ion-exchanged intermediate material forthe possible recovery of the layered structure from the convertedspinel phase.
2. Experimental
2.1. Synthesis of Li-rich layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2
nanoparticles
Li-rich layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2 nanoparticles weresynthesized by using co-precipitation method. Three precursorsolutions were simultaneously prepared. 0.08 mol transition metalprecursor at a molar ratio of Mn(CH3COO)2$4H2O: Ni(CH3-COO)2$4H2O: Co(CH3COO)2$4H2O ¼ 0.54: 0.13: 0.13 was dissolvedin 50 mL ethanol; the lithium precursor solution was composed of0.12mol LiOH dissolved in 20mL distilled water; and the surfactantsolution was 5.4 mmol F127 (EO106PO70EO106) dissolved in 50 mLethanol. The F127/ethanol solution and transition metal precursorsolution was first mixed together at 40 �C under continuous stir-ring, and then the lithium precursor solution was dropwise addedto precipitate transition metal ions. The resulting suspension washeated at 80 �C to completely remove the solvent and then dried inair at 120 �C for 12 h. The dried powder was annealed in air at300 �C for 3 h at a temperature ramp of 1 �C/min, followed bysintering at 900 �C for 12 h at a temperature ramp of 5 �C/min. Li[Li0.2Mn0.54Ni0.13Co0.13]O2 nanoparticles were obtained after cool-ing to room temperature.
2.2. Chemical activation of Li-rich layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2 via ion exchanges and post-heat treatments
The chemical activation of Li-rich layered oxide was carried asfollows: first, 1 g Li[Li0.2Mn0.54Ni0.13Co0.13]O2 particles weredispersed in 150 mL 2 M HCl aqueous solution for the Hþ-Liþ ionexchange of Li[Li0.2Mn0.54Ni0.13Co0.13]O2 at ambient temperature.The HCl solution was replaced every 2 days for 5 times in order toachieve deep protonation. The protonated intermediates dispersedin 150 mL aqueous HCl was obtained in a brown suspension. Sec-ondly, an aqueous tetrabutylammonium (TBA$OH) solution (SigmaAldrich) with a mass rate of 20 wt.% was employed to performTBAþ-Hþ exchange of protonated particles. The volumetric ratio ofTBA$OH solution over the brown suspension was set to 5: 1, andthese two solution was mixed in a vortex stirrer for 30 min. Allresulting ion-exchanged particles were collected via centrifugationand washed with distilled water for several times. Thirdly, the Liþ-TBAþ exchange was carried out in 1 M LiOH aqueous solution tostudy the reversibility of different ion exchanges. Finally, all ion-
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exchanged derivatives were annealed in air at 500 �C for 3 h at atemperature ramp of 1 �C/min.
2.3. Characterizations
Crystallographic structures and phases of Li-rich layered nano-particles and all derivatives were analyzed by X-ray diffraction(XRD) on a Panalytical X'pert Diffractometer with Cu Ka radiation.Morphology and particle characteristics of different samples wereexamined using a field emission scanning electron microscopy(FESEM, Hitachi S4800). Detailed structures of different sampleswere observed on transmission electron microscopy (TEM, FEITecnai G2 FEG) at an acceleration voltage of 300 kV. Porous struc-ture and specific surface area of powders wasmeasured by nitrogenadsorption/desorption at 77 K on a Quantachrome AS-1 instrumentusing the Brunauer-Emmet-Teller (BET) method. Chemical com-positions of specimens were determined by inductively coupledplasma atomic emission spectrometry (ICP-AES) on a SPCTROCIROS elemental analyzer.
2.4. Electrochemical measurements
The cathodes were consisted of 80 wt.% Li[Li0.2Mn0.54-
Ni0.13Co0.13]O2 particles or its derivatives, 10 wt.% acetylene black asthe conductive carbon (Alfa Aesar, 99.5%), and 10 wt.% poly-vinylidene fluoride (PVDF) as the binder (Alfa Aesar). These elec-trodes were assembled into CR2032-type coin cells forelectrochemical measurements, with themetallic lithium foil as theanode and Celgard 2320 membrane as the separator. The electro-lyte was 1 M LiPF6 dissolved in ethylene carbonate (EC), dimethylcarbonate (DMC) and diethyl carbonate (DEC) at a volumetric ratioof 1:1:1. Galvanostatic charge and discharge were performed atdifferent current densities in a voltage range of 2.0e4.8 V vs. Liþ/Liusing an 8-channel battery analyzer (MTI Corporation). Theoreticalcapacities of different cathode materials are all set to 250 mAh/g,i.e., current density corresponding to 1 C is 250 mA/g. Cyclic vol-tammetric (CV) curves of cathodes were recorded at a scanning rateof 0.1 mV/s between 2.0 and 4.8 V vs. Liþ/Li using an electro-chemical analyzer (CHI 605C).
3. Results and discussion
As shown in Fig. 1, ion-exchange processes and calcinationsresult in dramatically morphological changes of different de-rivatives. The pristine Li-rich Li[Li0.2Mn0.54Ni0.13Co0.13]O2 particlesexhibit a distinct aggregation with an even particle size around~250 nm in Fig. 1a, while Hþ-Liþ ion exchange in acidic environ-ment gives rise to distinct layered cake-shaped blocks of LHMNCO.We speculate that the multilayered morphology and structure ofLHMNCO (Fig. 1b) is probably attributed to the structural intro-duction of layered LMNCO. As reported in literature [26,29,31], theprotonation is required for the subsequent Hþ-TBAþ ion exchange.After shaking the intermediate LHMNCO and HCl mixture (i.e.,protonated LHMNCO particles dispersed in the final HCl solution)via violent vortexes in the aqueous TBA$OH solution at a volumetricratio of 1:5, the collected LHMNCO TBA shows the interestingmorphology of nanoflowers (Fig. 1c). Each particle is composed ofnumerous ultrathin nano-petals. This phenomenal morphologychange is mostly resulted from the cooperative effects of the TBA-assisted exfoliation and the turbulence-induced reaction environ-ment. It is interesting that further Liþ-TBAþ ion exchange in a basicsolution extensively unfolds petals of LHMNCO TBA nanoflowersinto nanosheet stacks of LHMNCO TBA Li as shown in Fig. 1d.Fig. 1eeg reveal morphologies of ion-exchanged derivatives afterpost-annealing treatments, which apparently cause aggregations
and coarse structures of LHMNCO HT, LHMNCO TBA HT andLHMNCO TBA Li HT, respectively, after the removal of Hþ and TBAþ
ions along with other byproducts, such as H3Oþ and OH�.Accordingly, phase transitions and structural reconstructions
accompanying with morphology changes from initial layeredLMNCO to different converted derivatives have been studied fromXRD and TEM characterizations in Fig. 2 and Fig. 3, respectively. Inagreement with reported XRD patterns of Li-rich layered materials[7,8,17,18], pristine LMNCO in Fig. 2a shows typical XRD peaks thatare indicative of the intergrowth of monoclinic Li2MnO3 (spacegroup: C2/m) and rhombohedral LiMn1/3Ni1/3Co1/3O2 (space group:R-3m) in the layered structure [7]. The main layered structure ofLiMn1/3Ni1/3Co1/3O2 can be determined from distinct peak splits oftwo (006)L2-(012)L2 and (108)L2-(110)L2 doublets at 2q ¼ 36 - 38�
and 2q ¼ 64 - 66�, respectively. The weak (020)L1 reflection at2q ¼ 20 - 23� is belong to the layered Li2MnO3 phase, which is thesuperlattice within the parent layered structure [17]. The inter-growth of these two layered components can be further confirmedfrom the other five peaks at 2q¼ 18.7�, 36.9�, 37.9�, 38.5� and 44.5�,resulting in diffraction patterns of (001)L1/(003)L2, (200)L1/(101)L2,(113)L1/(006)L2, (131)L1/(012)L2, and (202)L1/(104)L2. A spinel-likeimpurity is found in the XRD pattern of LHMNCO (marked withred asterisks in Fig. 2b), which is accordant with that reported inthe literature [7,22]. As shown in Fig. 2b, the peakmerger of (113)L1/(006)L2 and (131)L1/(012)L2 pairs reveals the distortion of layeredstructure at a certain degree, while the preserved peak splits of(108)L2-(110)L2 doublets indicate the retained layered structure inthe protonated LHMNCO intermediate [22]. Furthermore, XRDpattern of LHMNCO TBA powder manifests the growth of spinel-like phase during TBAþ-Hþ exchange, since the intensity of onerepresentative XRD peak at 2q ¼ 19.3� in LHMNCO for the spinelphase increases. The cubic close packed oxygen arrays both inlayered and spinel structures is essential to realize the phasetransition in Li-rich layered transition metal oxides by migratingtransitionmetal ions into lithium layers when lithium ion vacanciesexist during ion-exchange processes [13,20]. LHMNCO TBA Li showsidentical XRD pattern to that of LHMNCO TBA, indicating limitedreversibility of phase transition by Liþ-TBAþ ion exchange. How-ever, the peak splitting of (108)L2-(110)L2 doublets in its XRDpattern reveals the prominent layered structure of LHMNCO TBA Liderivative, despite the spinel phase has been detected. Post-annealing treatment has been demonstrated as an effective wayto remove Hþ and TBAþ substituents in air, resulting in the gener-ation of corresponding Liþ vacancies in lithium layers [32]. Suchfacile process can accelerate the diffusion of transition metal ionsinto lithium ion sites, and thus promote the layered-to-spinel phasetransition. The enlarged selected 2q portions in Fig. 2b at 2q ¼ 16-20�, 34-40� and 62-68� illustrate the phase transitions whendifferent ion-exchanged samples are subjected to calcinations inair. The merge of two separate peaks around 2q¼ 19� into one peakoccurs for LHMNCO, LHMNCO TBA and LHMNCO TBA Li after heattreatments, respectively, indicating dramatic phase transitions dueto the removal of foreign Hþ and TBAþ cations. It is surprising tofind that the coupled (108)L2-(110)L2 pair of LHMNCO TBA hasmerged to one broad peak for LHMNCO TBA HT at a lower 2q po-sition, while LHMNCO HT and LHMNCO TBA Li HT still show dis-tinguishing peak splits of (108)L2-(110)L2 between 2q¼ 62� and 64�.As mentioned before, the peak split of (108)L2-(110)L2 doublets ischaracteristic of layered structure, which is distinguished from thespinel phase showing the (440)S reflection at the same position. Asa result, XRD pattern of LHMNCO TBA HT can be indexed to thespinel Li4Mn5O12 phase with a Fd-3m space group, indicating thecomplete layered-to-spinel phase transition from original Li-richlayered LMNCO to Li4Mn5O12-type spinel compound after two-step ion exchanges, followed by a post-annealing process. In
Fig. 1. SEM images of (a) LMNCO, (b) LHMNCO, (c) LHMNCO TBA, (d) LHMNCO TBA Li, (e) LHMNCO HT, (f) LHMNCO TBA HT and (g) LHMNCO TBA Li HT.
Fig. 2. XRD patterns of pristine Li-rich layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2 nanoparticles and corresponding derivatives in (a) full 2q range and (b) enlarged 2q portions between 16and 22� , 34-40� and 62-68� , respectively.
J. Zhao et al. / Journal of Alloys and Compounds 732 (2018) 385e395388
Fig. 3. TEM and HRTEM images of (a and b) LMNCO, (c and d) LHMNCO, (e and f) LHMNCO TBA, and (g and h) LHMNCO TBA HT.
J. Zhao et al. / Journal of Alloys and Compounds 732 (2018) 385e395 389
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contract, XRD patterns of LHMNCO HT and LHMNCO TBA Li HT bothreveal the coexistence of layered and spinel phases (Fig. 2). It issuggested that a second TBAþ-Hþ ion exchange is crucial to realize acomplete phase conversion.We speculate that due to the larger sizeof TBAþ cations than protons, TBAþ substituents can increase c-axisof ion-exchanged layered derivative in comparison with the effectfrom Hþ ions, which will increase the structure instability thatsignificantly facilitate migrations of transition metal ions whenTBAþ are burned away in air. On the contrary, Liþ can be partiallyrestored in the lithium layers through the Liþ-TBAþ ion exchange,resulting in hybrid layered-spinel structure of LHMNCO TBA Li HT.
Fig. 3 shows the structural evolution from the pristine layeredLMNCO to the converted Li4Mn5O12-type LHMNCO TBA HT spinel inTEM and HRTEM observations. In consistence with SEM image inFig. 1a, LMNCO nanoparticles have the solid structure with anaverage particle size of ~250 nm. The lattice fringe as shown inFig. 3b indicates the high crystallinity of pristine Li-rich layeredmaterials, due to the high synthetic heating temperature at 900 �C.The characteristics of overlapped sheets can be observed in Fig. 3cof LHMNCO after Hþ-Liþ ion exchange in an acidic HCl solution. Theformation of spinel phase during this process at the surface ofparticle is identified in HRTEM image (Fig. 3d), which is consistentwith XRD results in Fig. 2. The continuous TBAþ-Hþ ion exchangenot only tailors LHMNCO TBA to the nanoflower-like shape, but alsogenerates porous structure as shown in Fig. 3e. Furthermore, due tothe violent exfoliation effect of TBAþ cations, HRTEM image inFig. 3f shows disordered lattice fringes both at the surface and inthe bulk of LHMNCO TBA; the other reason leading to such adisordered structure possibly results from the partial decomposi-tion of organic TBAþ cations under the attack of high-energy elec-tron beam during HRTEM observations. On the other hand, fringeswith the smaller d-space (d ¼ 0.455 and 0.438 nm in Fig. 3f) at thesurface of the specimen may also be attributed to the partialdecomposition of TBAþ substituents. As shown in Fig. 1c and f,monodispersive LHMNCO TBA nanoflowers convert to LHMNCOTBA HT particles with irregular shapes after post-annealing pro-cesses due to the fold of nanopetals. Accordingly, the removal ofTBAþ cations also contributes to the porous structure of LHMNCOTBA HT as shown in TEM image (Fig. 3g). According to the XRDresults as shown in Fig. 2, HRTEM image in Fig. 3h shows the spinelcrystal structure of LHMNCO TBA HT with a d-space of (111)S equalto 0.471 nm. In general, ex-situ ion-exchange and heat treatmentsresult in the complete phase transition from the layered LMNCO to
Fig. 4. (a) Nitrogen adsorption/desorption isotherms and (b) corresponding p
a Li4Mn5O12-type spinel material, along with intriguing morpho-logical and structural evolutions.
The nitrogen adsorption and desorption isotherms and pore sizedistributions of pristine Li-rich layered LMNCO, ion-exchangedLHMNCO TBA and annealed LHMNCO TBA HT are shown inFig. 4a and b, respectively. The corresponding porous characteris-tics in terms of surface area, pore volume and relative pore size aresummarized in Table 1 together with Li/Mn/Ni/Co ratios of threesamples. It is clear that substitution of Liþ ions within LMNCO by Hþ
protons, followed by continuous replacement with TBAþ cationsresults in significantly increased surface area to 11.109 m2/g ofLHMNCO TBA, almost four times higher than that of originalLMNCO particles (2.327 m2/g), which can be attributed to theexfoliation effect from TBAþ cations as shown in SEM (Fig. 1c) andTEM (Fig. 3e) images [27,31]. Accordingly, LHMNCO TBA also has alarger pore volume of 8.880e-2 cm3/g in comparison with 1.133e-2cm3/g from LMNCO nanoparticles. Both SEM and TEM observationsindicate that the pore volume of LMNCO powder is from specialgaps between numerous agglomerated LMNCO nanoparticles(Figs. 1a and 3a), while the higher pore volume of LHMNCO TBAmostly arises from the porous structure of individual LHMNCO TBAnanoflowers (Figs. 1c and 3e). Therefore, agglomerated LMNCOnanoparticles give rise to a relatively higher pore size distributionof ~6 nm in Fig. 4b as compared with ~4 nm from LHMNCO TBAnanoflowers with monodisperse characteristic. As aforementionedin XRD characterizations in Fig. 2, the post-annealing treatmentplays a crucial role in realizing a complete layered-to-spinel phasetransition, and themorphological and structural changes have beenobserved in SEM (Fig. 1f) and TEM (Fig. 3g) images, respectively. Asa result, heating LHMNCO TBA in air at 500 �C contributes to furtherincreased surface area to 13.725 m2/g of LHMNCO TBA HT and twopore size distributions of ~3 and ~12 nm in Fig. 4d. Its reduced porevolume is probably due to the folded nanopetals to form an internalporous structure (Fig. 3g) and the obvious aggregation (Fig. 1f) afterpost-heat treatment. The chemical compositions of these threesamples are compared in Table 1 in the form of Li/Mn/Ni/Co molarratios. In comparison with the theoretical ratio of Li/Mn/Ni/Co ¼ 1.2/0.54/0.13/0.13, the as-prepared LMNCO shows slightly lessquantity of the lithium component that is probably assigned to thelithium loss during heat treatment at high temperature at 900 �Ctogether with the long duration time for 12 h. Such the harsh heattreatment is required for the synthesis of Li-rich LMNCO material,in order to achieve its nice integrated structure and high
ore size distributions of LMNCO and LHMNCO TBA and LHMNCO TBA HT.
Table 1Porous characteristics and elemental composition of LMNCO, LHMNCO TBA and LHMNCO TBA HT.
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crystallinity, but leads to unfavorable lithium loss as measured byICP results. After two step ion exchanges via Hþ-Liþ and TBAþ-Hþ,LHMNCO TBA preserves 52.7% of the original lithium ions inLMNCO. As reported in the literature [29], layered transition metaloxides can be fully exchanged and teared into transition metaloxide nanosheets. In our case, the partial ion exchange possibly isattributed to the formation of spinel phase at the surface of layeredderivatives as shown in XRD results (Fig. 2) and HRTEM image(Fig. 3d). The detectable Ni loss in LHMNCO TBA after the ion-exchange processes may result from the cationic Liþ-Ni2þ disor-der in LMNCO. Consequently, a few Ni2þ ions occupied in lithiumsites in the lithium layer are replaced with Hþ and TBAþ cations.The post-annealing process has few effect on the chemicalcomposition of LHMNCO TBA HT, resulting in Li/M(M ¼ Mn þ Ni þ Co) ¼ 0.75. Such value is much close to the Li/Mnratio of 0.8 in Li4Mn5O12 spinel rather than 0.5 of LiMn2O4 spinel,which is consistent with the XRD result being indexed toLi4Mn5O12-type spinel for LHMNCO TBA HT. In short summary, ex-situ ion-exchange processes along with post-heat treatments offera feasible approach not only to tailor morphology and structure ofLi-rich layered transition metal oxides, but also to control the phasetransformation between layered and spinel phases. We speculatethat the nanoarchitectured LHMNCO TBA HT spinel material shouldbe favorable to facilely accommodate the electrolyte, maximizeelectrochemical active sites and release reaction strain duringlithiation and delithiation; hence, LHMNCO TBA HT cathode with apure spinel phase and porous structure is expected to showenhanced rate capability and cycleability as compared with thepristine Li-rich layered LMNCO as well as other two hybridLHMNCOHTand LHMNCO TBA Li HTwith an intergrowth of layeredand spinel phases.
The cyclic voltammetric (CV) measurements of LMNCO,LHMNCO HT, LHMNCO TBA HT and LHMNCO TBA Li HT cathodematerials are carried out, in order to study electrochemical prop-erties related to phase transitions from Li-rich layered (LMNCO) toeither layered-spinel (LHMNCO HT and LHMNCO TBA Li) orLi4Mn5O12-type spinel (LHMNCO TBA HT) phase. Fig. 5 showsresulting CV records in the first three cycles of four different sam-ples. LMNCO reveals the typical electrochemical characteristics ofLi-rich layered cathode materials as shown in Fig. 5a. The firstanodic peak at 4.17 V in the initial charge curve is associated withthe oxidation of Ni2þ to Ni4þ, followed by Co3þ to Co4þ, whereasMnstill remains as tetravalent in LiMn1/3Ni1/3Co1/3O2 component [33].The second anodic peak at 4.66 V corresponds to the electro-chemical activation of inert Li2MnO3 component, i.e., the decom-position of Li2MnO3 to Li2O and lithium-active MnO2, along withthe unavoidable decomposition of electrolyte and the formation ofsolid electrolyte interphase (SEI) at such a high potential >4.5 V[10]. Although the electrochemical activation process leads to thelow Coulombic efficiency in the first cycle, but significantly resultsin high capacity of Li-excess layered cathode materials in the suc-cessive cycles. Correspondingly, the reduction of Co4þ/Co3þ andNi4þ/Ni3þ/Ni2þ redox occurs at 3.66 V in the initial discharge curve.In the second cycle, the anodic peak at 4.66 V disappears, while anadditional cathodic peak at 3.26 V appears, which can be attributed
to the reduction of Mn4þ to Mn3þ from the as-activated MnO2component. The third cycle shows the similar profile to the secondcycle in less polarization and higher current density, indicatingimproved electrochemical reversibility after the electrochemicalactivation of Li-rich layered LMNCO. It is clear to see that CV curvesof LHMNCO HT, LHMNCO TBA HT and LHMNCO TBA Li HT are verysimilar to each other, but apparently different from that of LMNCO.Those three materials all show the dominant redox pair around3.0 V in CV curves, together with two minor redox couples locatednear 4.0 and 4.6 V. The CV performance is consistent with thetypical electrochemical characteristics of reported Ni/Co-dopedLi4Mn5O12-type spinel in a wide voltage range [34]. The CV re-sponses of LHMNCO HT, LHMNCO TBA HT and LHMNCO TBA Li HTsupport XRD results in Fig. 2, which reveal the Li4Mn5O12-typespinel structure of newly-formed spinel phase within the originallayered structure of LMNCO after ion-exchange and post-annealingprocesses. The appearance of an anodic peak at ~4.6 V in the initialCV charges of LHMNCO HT and LHMNCO TBA Li HT indicate theexistence of preserved layered structure, in accordance with XRDcharacterizations (Fig. 2). In contract, LHMNCO TBA HT revealsmuch lower current density of such anodic peak in the first CVcharge. Furthermore, a more intensive anodic peak at 2.95 V isgenerated, which can be attributed to the better complete phasetransition of LHMNCO TBA HT in comparisonwith LHMNCO HT andLHMNCO TBA HT. LHMNCO TBA HT shows identical CV curves ofsecond and third cycles, indicating outstanding electrochemicalreversibility of this spinel materials. The anodic peak at 3.07 V andcathodic peak at 2.63 V are attributed to corresponding oxidationand reduction reactions of Mn3þ/Mn4þ redox pair, which areassociated with extracting and inserting lithium ions on 16c sites inthe spinel structure [1]. Two small redox couples at ~4.0 and ~4.6 Vprobably result from the Co3þ/Co4þ and Ni2þ/Ni3þ/Ni4þredox,respectively.
Fig. 6 exhibits charge and discharge curves of four cathodes inthe first five cycles at 0.1 C in a voltage range of 2.0e4.8 V vs. Liþ/Li,which are in well accordance with CV profiles in Fig. 5. It isnoticeable that introduction of a spinel phase with Li-rich layeredcathodematerials can significantly increase the specific capacity. Asshown in Fig. 6aed, charge/discharge curves of the fourth and fifthcycle are almost identical for all cathodes, suggesting that thecathodes are mostly stable after five electrochemical cycles. Thepristine Li-rich layered LMNCO delivers a specific discharge ca-pacity of 211.3 mAh/g at the fifth cycle with a voltage plateauaround 3.7 V, revealing typical electrochemical behavior of Li-richlayered cathode materials. The XRD result in Fig. 2a demonstratesthe coexistence of layered and spinel phases of LHMNCO HT afterex-situ ion-exchange and post-heat treatments, in which the spinelphase is dominant; hence LHMNCO HT shows a higher dischargecapacity of 285.8 mAh/g but along with a predominant voltageplateau at 2.6 V as well as a minor one at 4.4 V. The former voltagestage is due to the active Mn3þ/Mn4þ redox in the spinel compo-nent, and the latter probably results from the Co4þ/Co3þ and Ni4þ/Ni3þ/Ni2þ redox reactions in the reversed layered component.Furthermore, employing TBAþ cations for the second ion exchangeof LHMNCO have contributed to the complete phase conversion of
Fig. 5. Cyclic voltammetric (CV) curves of (a) LMNCO, (b) LHMNCO HT, (c) LHMNCO TBA HT and (d) LHMNCO TBA Li HT in the first three cycles at a scanning rate of 0.1 mV/s in avoltage range of 2.0e4.8 V vs. Liþ/Li.
J. Zhao et al. / Journal of Alloys and Compounds 732 (2018) 385e395392
Li-rich layered LMNCO, resulting in the spinel LHMNCO TBA HTwith the Li4Mn5O12-type spinel characteristics (Fig. 2) and a mes-oporous structure (Figs. 3g and 4b). As a result, an unexpectedlyhigh discharge capacity of 343.2 mAh/g is achieved in the fifth cycleof LHMNCO TBA HT. There are three voltage plateaus located at 4.6,4.0 and 2.8 V, respectively, which is distinctly different from theprofile of LHMNCO HT, again revealing different structural charac-teristics between LHMNCO TBA HT and LHMNCO HT. Those threevoltage stages can be attributed to reductions from Ni4þ/Ni3þ/Ni2þ,Co4þ/Co3þ and Mn4þ/Mn3þ redox pairs in the spinel structure andare favorable to preserve high-voltage performances of LHMNCOTBA HT. Fig. 3g reveals a porous mesoporous structure of nano-architectured spinel cathode, which would be favorable to accom-modate electrolyte and effectively release reaction strains duringlithiation/delithiation. Such a structure may enable to absorblithium ions within the porous structure, resulting in the additionalcapacity contribution [35,36]. On the other hand, the porousstructure may partially collapse when adsorbed lithium ions areextracted during the charge processes, leading to reduced lithiumstorage capacity in corresponding discharge processes. Accordingly,the LHMNCO TBA HT cathode (Fig. 6c) shows relatively low butgradually increased Coulombic efficiencies in initial cycles ascompared with the other three cathode materials. In contrast,LHMNCO TBA Li HT shows similar electrochemical performance to
LHMNCO HT with the intergrowth of layered and spinel structures,delivering a reduced capacity of 259.9 mA/g. This might result fromthe partially recovered layered structure via Liþ-TBAþ ion exchangeof LHMNCO TBA in LiOH aqueous solution.
Fig. 7a and b shows cycling and high-rate performances ofLHMNCO HT, LHMNCO TBA HT and LHMNCO TBA Li HT cathodes incomparison with the pristine Li-rich layered LMNCO, respectively.The effects from the introduced spinel phase within the layeredLMNCO cathode material on improving specific capacity, cyclingstability and rate capability are more phenomenal when cycled athigher current densities. As shown in Fig. 7a, LHMNCO TBA HT canretain a very high discharge capacity of 197.5 mAh/g with a corre-sponding capacity retention of 89.1% after 100 electrochemicalcycles at 1C, much better than 58.1 mAh/g and 65.9% of pristinelayered LMNCO, 116.1 mAh/g and 85.1% of LHMNCO HT, and77.9 mAh/g and 80.0% of LHMNCO TBA Li HT. Moreover, LHMNCOTBA HT delivers initial capacities of 313.6, 267.2, 203.9, 180.7, 126.3,and 89.4mAh/g at 0.1, 0.5, 1, 2, 5, and 10 C, respectively, as exhibitedin Fig. 7b. Such the remarkable cyclability and high-rate capabilityof LHMNCO TBA HT can be attributed to reconstructed spinel phaseand hierarchical mesoporous structure for facile accommodationand diffusion of lithium ions, and effectively releasing reactionstrains in the “buffer” structure in high porous characteristics.Overall, generation of a spinel phase within Li-rich layered cathode
Fig. 6. Charge and discharge curves of (a) LMNCO, (b) LHMNCO HT, (c) LHMNCO TBA HT and (d) LHMNCO TBA Li HT in the first five cycles at a current density of 0.1 C in a voltagerange of 2.0e4.8 V vs. Liþ/Li.
Fig. 7. (a) Cycling performances at 1 C and (d) high-rate performances at different current densities of LHMNCO HT, LHMNCO TBA HT, LHMNCO TBA Li HT in comparison with thepristine layered LMNCO in a voltage range of 2.0e4.8 V vs. Liþ/Li.
J. Zhao et al. / Journal of Alloys and Compounds 732 (2018) 385e395 393
materials can considerably increase the specific capacity and ratecapability, but has to sacrifice the working voltage. Doping transi-tion metal cations, such as Ni2þ, Co3þ and Fe3þ ions, can contribute
to the high-voltage performance. Ion-exchange method offers adesirable way to obtain enhanced electrochemical performance ofLi-rich layered cathode materials.
J. Zhao et al. / Journal of Alloys and Compounds 732 (2018) 385e395394
4. Conclusions
This work sheds light on fundamental understanding of layered-to-spinel phase transition and relevant electrochemical perfor-mances of Li-rich layered cathode materials via ex-situ ion-exchange processes, followed by post-annealing treatments.Employing TBAþ cations for the second ion exchange of pronatedLi-rich layered oxides is critical to realize a complete phase tran-sition, resulting in a Li4Mn5O12-type spinel-structured materialconverted from Li-rich layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2. Ion-exchange procedure also tailor the morphology and structure ofsolid Li[Li0.2Mn0.54Ni0.13Co0.13]O2 nanoparticles into nanostructuredspinel material with high surface area and mesoporous porosity. Incomparison with the pristine Li-rich layered cathode material, thefinal converted spinel cathode material with hierarchical porousstructure reveals significantly increased specific capacity, bettercycling stability and rate capability. This work opens up a route tomaximize electrochemical performance of Li-excess layered cath-ode materials for high-power and high-energy lithium ionbatteries.
Acknowledgements
This work was supported by US National Science Foundation,the Division of Chemical, Bioengineering, Environmental andTransport Systems (NSF CBET) [grant number 1438493]; the USSmall Business Technology Transfer (STTR) [grant number1346496]; the Research Enhancement Award (REA) program, Lou-isiana Space Consortium (LaSPACE) funded via the NASA SpaceGrant College & Fellowship Program Grant 2011-15 Cycle [grantnumber NASA/LEQSF(2010e2015)-LaSPACE]; the National NaturalScience Foundation of China [grant number U1401248]; the NaturalScience Foundation of Jiangsu Province, China [grant numberBK20151227]; the General Financial Grant from the China Post-doctoral Science Foundation [grant number 2016M601876]. Theauthors also acknowledge LSU IAM Shared Instrumentation Facility(SIF) at Louisiana State University and Suzhou Key Laboratory forAdvanced Carbon Materials and Wearable Energy Technologies,Suzhou 215006, China for using facilities, and thank Prof. KerryDooley at Department of Chemical Engineering at Louisiana StateUniversity for BET measurements, and Dr. Jibao He for TEM obser-vations at Tulane University.
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