Structural Stability of LiCoO2 at 400°

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材料力学词汇

材料力学常用英文词汇弹性力学elasticity弹性理论theory of elasticity均匀应力状态homogeneous state of stress应力不变量stress invariant应变不变量strain invariant应变椭球strain ellipsoid均匀应变状态homogeneous state of strain应变协调方程equation of strain compatibility拉梅常量Lame constants各向同性弹性isotropic elasticity旋转圆盘rotating circular disk楔wedge开尔文问题Kelvin problem布西内斯克问题Boussinesq problem艾里应力函数Airy stress function克罗索夫--穆斯赫利什维利法Kolosoff-Muskhelishvili method 基尔霍夫假设Kirchhoff hypothesis板Plate矩形板Rectangular plate圆板Circular plate环板Annular plate波纹板Corrugated plate加劲板Stiffened plate,reinforced Plate中厚板Plate of moderate thickness弯[曲]应力函数Stress function of bending壳Shell扁壳Shallow shell旋转壳Revolutionary shell球壳Spherical shell[圆]柱壳Cylindrical shell锥壳Conical shell环壳Toroidal shell封闭壳Closed shell波纹壳Corrugated shell扭[转]应力函数Stress function of torsion翘曲函数Warping function半逆解法semi-inverse method瑞利--里茨法Rayleigh-Ritz method松弛法Relaxation method莱维法Levy method松弛Relaxation量纲分析Dimensional analysis自相似[性] self-similarity影响面Influence surface接触应力Contact stress赫兹理论Hertz theory协调接触Conforming contact滑动接触Sliding contact滚动接触Rolling contact压入Indentation各向异性弹性Anisotropic elasticity颗粒材料Granular material散体力学Mechanics of granular media 热弹性Thermoelasticity超弹性Hyperelasticity粘弹性Viscoelasticity对应原理Correspondence principle褶皱Wrinkle塑性全量理论Total theory of plasticity 滑动Sliding微滑Microslip粗糙度Roughness非线性弹性Nonlinear elasticity大挠度Large deflection突弹跳变snap-through有限变形Finite deformation格林应变Green strain阿尔曼西应变Almansi strain弹性动力学Dynamic elasticity运动方程Equation of motion准静态的Quasi-static气动弹性Aeroelasticity水弹性Hydroelasticity颤振Flutter弹性波Elastic wave简单波Simple wave柱面波Cylindrical wave水平剪切波Horizontal shear wave竖直剪切波Vertical shear wave体波body wave无旋波Irrotational wave畸变波Distortion wave膨胀波Dilatation wave瑞利波Rayleigh wave等容波Equivoluminal wave勒夫波Love wave界面波Interfacial wave边缘效应edge effect塑性力学Plasticity可成形性Formability金属成形Metal forming耐撞性Crashworthiness结构抗撞毁性Structural crashworthiness拉拔Drawing破坏机构Collapse mechanism回弹Springback挤压Extrusion冲压Stamping穿透Perforation层裂Spalling塑性理论Theory of plasticity安定[性]理论Shake-down theory运动安定定理kinematic shake-down theorem 静力安定定理Static shake-down theorem率相关理论rate dependent theorem载荷因子load factor加载准则Loading criterion加载函数Loading function加载面Loading surface塑性加载Plastic loading塑性加载波Plastic loading wave简单加载Simple loading比例加载Proportional loading卸载Unloading卸载波Unloading wave冲击载荷Impulsive load阶跃载荷step load脉冲载荷pulse load极限载荷limit load中性变载nentral loading拉抻失稳instability in tension加速度波acceleration wave本构方程constitutive equation完全解complete solution名义应力nominal stress过应力over-stress真应力true stress等效应力equivalent stress流动应力flow stress应力间断stress discontinuity应力空间stress space主应力空间principal stress space静水应力状态hydrostatic state of stress对数应变logarithmic strain工程应变engineering strain等效应变equivalent strain应变局部化strain localization应变率strain rate应变率敏感性strain rate sensitivity应变空间strain space有限应变finite strain塑性应变增量plastic strain increment累积塑性应变accumulated plastic strain永久变形permanent deformation内变量internal variable应变软化strain-softening理想刚塑性材料rigid-perfectly plastic Material 刚塑性材料rigid-plastic material理想塑性材料perfectl plastic material材料稳定性stability of material应变偏张量deviatoric tensor of strain应力偏张量deviatori tensor of stress应变球张量spherical tensor of strain应力球张量spherical tensor of stress路径相关性path-dependency线性强化linear strain-hardening应变强化strain-hardening随动强化kinematic hardening各向同性强化isotropic hardening强化模量strain-hardening modulus幂强化power hardening塑性极限弯矩plastic limit bending Moment塑性极限扭矩plastic limit torque弹塑性弯曲elastic-plastic bending弹塑性交界面elastic-plastic interface弹塑性扭转elastic-plastic torsion粘塑性Viscoplasticity非弹性Inelasticity理想弹塑性材料elastic-perfectly plastic Material 极限分析limit analysis极限设计limit design极限面limit surface上限定理upper bound theorem上屈服点upper yield point下限定理lower bound theorem下屈服点lower yield point界限定理bound theorem初始屈服面initial yield surface后继屈服面subsequent yield surface屈服面[的]外凸性convexity of yield surface截面形状因子shape factor of cross-section沙堆比拟sand heap analogy屈服Yield屈服条件yield condition屈服准则yield criterion屈服函数yield function屈服面yield surface塑性势plastic potential能量吸收装置energy absorbing device能量耗散率energy absorbing device塑性动力学dynamic plasticity塑性动力屈曲dynamic plastic buckling塑性动力响应dynamic plastic response塑性波plastic wave运动容许场kinematically admissible Field静力容许场statically admissible Field流动法则flow rule速度间断velocity discontinuity滑移线slip-lines滑移线场slip-lines field移行塑性铰travelling plastic hinge塑性增量理论incremental theory of Plasticity米泽斯屈服准则Mises yield criterion普朗特--罗伊斯关系prandtl- Reuss relation特雷斯卡屈服准则Tresca yield criterion洛德应力参数Lode stress parameter莱维--米泽斯关系Levy-Mises relation亨基应力方程Hencky stress equation赫艾--韦斯特加德应力空间Haigh-Westergaard stress space 洛德应变参数Lode strain parameter德鲁克公设Drucker postulate盖林格速度方程Geiringer velocity Equation。

111比列镍钴锰三元正极材料的开山之作日本Tsutomu Ohzuku发表于Chemistry Letters

––Copyright ©2001 The Chemical Society of Japan6 wt% PVdF. The electrode was dried under vacuum at 150 °C for 12 h. A lithium electrode was prepared by pressing a lithi-um metal sheet onto a stainless steel plate (15 × 20 mm 2). Two sheets of porous polypropylene membrane (Celgard 2500) were used as a separator. The electrolyte used was 1M LiPF 6dis-solved in ethylene carbonate (EC) / dimethyl carbonate (DMC)(3/7 by volume) solution. Open-circuit voltages of the freshly fabricated cells were around 3.2 V. During charge at 0.17 mA cm –2, the cell voltage rapidly increased to 3.7 V and then stayed at 3.7–3.75 V until the charge capacity reaches about 80 mAh g –1. On further charging, the voltage increased monotonously to 4.2 V of charge-end voltage. The first charge capacity was ca. 165 mAh g –1based on LiCo 1/3Ni 1/3Mn 1/3O 2sample weight,while the first discharge capacity was ca. 150 mAh g –1.Irreversible capacity was ca. 15 mAh g –1in this case. For sub-sequent cycles, the curves converged on a single charge or dis-charge curve, as seen in Figure 2. The rechargeable capacity was 150 mAh g –1when the cell was operated in 2.5–4.2 V.Figure 3 shows the charge and discharge curves of a Li /LiCo 1/3Ni 1/3Mn 1/3O 2cell operated at 0.17 mA cm –2in voltages between 2.5 and 5.0 V. The first charge capacity at 5.0 V of charge-end voltage was ca. 250 mAh g –1and the first discharge capacity was ca. 220 mAh g –1. Although a loss in capacity dur-ing charge and discharge was observed mainly due to elec-trolyte oxidation occurring in higher voltages than 4.5 V,rechargeable capacity more than 200 mAh g –1was obtained with little capacity fading. As reported previously,3,10,11LiCoO 2or LiNiO 2was rapidly deteriorated when the sample was charged to higher voltages than 4.5 V against a lithium electrode. As shown in Figure 3, a Li / LiCo 1/3Ni 1/3Mn 1/3O 2cell shows a better capacity retention than those of LiCoO 2or LiNi 1/2Co 1/2O 2in that voltage range. We believe that LiCo 1/3Ni 1/3Mn 1/3O 2with rechargeable capacity of 200 mAh g –1will be used in lithium-ion batteries.In conclusion LiCo 1/3Ni 1/3Mn 1/3O 2 is a possible alternative to LiCoO 2for lithium-ion batteries. In this letter we did not state the crystal structure of this sample. Its structure is close to LiCoO 2or LiNiO 2, but something bothers us about a nature of complex solid solution mechanism. Structural refinements together with the optimization of synthetic processes are now in progress in our laboratory.The authors wish to thank Mr. Hiroyuki Ito of Tanaka Chemical Corp. for his help on the preparation of a series of nickel manganese hydroxides. This work is supported in part by a grant-in-aid for scientific research from Osaka City University Science Foundation.References1K. Mizushima, P. C. Jones, P. J. Wiseman, and J. B.Goodenough, Mater. Res. Bull., 25, 783 (1980).2J. R. Dahn, U. von Sacken, and C. A. Michal, Solid StateIonics , 44, 87 (1990).3T. Ohzuku, A. Ueda, M. Nagayama, Y. Iwakoshi, and H.Komori, Electrochim. Acta , 38, 1159 (1993) and references therein.4J. C. Hunter, J. Solid State Chem., 39, 142 (1981).5M. M. Thackeray, W. I. F. David, P. G. Bruce, and J. B.Goodenough, Mater. Res. Bull., 18, 461 (1983).6T. Ohzuku, K. Sawai, and T. Hirai, J. Electrochem. Soc.,137, 3004 (1990).7T. Ohzuku, S. Kitano, M. Iwanaga, H. Matsuno, and A.Ueda, J. Power Sources , 68, 646 (1997).8T. Ohzuku, A. Ueda, and N. Yamamoto, J. Electrochem.Soc., 142, 1431 (1995).9T. Ohzuku, K. Nakura, and T. Aoki, Electrochim. Acta , 45,151 (1999).10T. Ohzuku, A. Ueda, and M. Nagayama, J. Electrochem.Soc., 140, 1862 (1993).11T. Ohzuku and A. Ueda, J. Electrochem. Soc., 141, 2972(1994).。

正极材料综述

Recent progress in cathode materials research for advanced lithium ion batteries Bo Xu,Danna Qian,Ziying Wang,Ying Shirley Meng*Department of NanoEngineering,University of California San Diego,La Jolla,CA92093,USAContents1.Introduction (51)yered compounds LiMO2 (52)3.Spinel compounds LiM2O4 (56)4.Olivine compounds LiMPO4 (57)5.Silicate compounds Li2MSiO4 (59)6.Tavorite compounds LiMPO4F (62)7.Borate compounds LiMBO3 (63)8.Conclusion (64)Acknowledgements (64)References (64)1.IntroductionWith the worldwide energy shortage being one of the mountingproblems in21st century,efforts have been made to replace thenon-renewable fossil fuels by other green energy sources,such assolar,wind,and hydroelectric power.Different from the conven-tional fossil fuels,most of these green energy sources suffer fromtheir uncontrollable and intermittent nature,therefore thedifﬁculty in energy storage and regulation results in larger cost.This brings in enormous amount of research interests in materialdevelopments for energy storage.The LIB system is regarded as oneof the near-term solutions because of its high energy density andrelatively simple reaction mechanism.Current LIB technology iswell developed for the portable electronic devices and has beenwidely used in the past twenty years.However,to be implementedin the large-scale high-power system such as the plug-in hybridelectric vehicle(PHEV)or plug-in electric vehicle(PEV),perfor-mance requirements are raised especially from the aspects ofenergy/power density,cycling life and safety issues,thereforefurther LIB material and system developments are necessary.The basic working principles of LIB are shown in Fig.1.A lithiumion battery can work as the energy storage device by convertingelectric energy into electrochemical energy.There are three keycomponents in a LIB system:cathode,anode and electrolyte.Fortoday’s commercialized LIB system,both cathode and anodematerials are intercalation materials.The transition metal oxidesin cathode(graphite in anode)consist of a largely unchangeablehost with speciﬁc sites for Li ions to be intercalated in.All Li ionsMaterials Science and Engineering R73(2012)51–65A R T I C L E I N F OArticle history:Available online5June2012A B S T R A C TNew and improved materials for energy storage are urgently required to make more efﬁcient use of ourﬁnite supply of fossil fuels,and to enable the effective use of renewable energy sources.Lithium ionbatteries(LIB)are a key resource for mobile energy,and one of the most promising solutions forenvironment-friendly transportation such as plug-in hybrid electric vehicles(PHEVs).Among the threekey components(cathode,anode and electrolyte)of LIB,cathode material is usually the most expensiveone with highest weight in the battery,which justiﬁes the intense research focus on this electrode.In thisreview,we present an overview of the breakthroughs in the past decade in developing high energy highpower cathode materials for lithium ion batteries.Materials from six structural groups(layered oxides,spinel oxides,olivine compounds,silicate compounds,tavorite compounds,and borate compounds)arecovered.We focus on their electrochemical performances and the related fundamental crystalstructures,solid-state physics and chemistry are covered.The effect of modiﬁcations on both chemistryand morphology are discussed as well.ß2012Elsevier B.V.All rights reserved.*Corresponding author.E-mail address:shirleymeng@(Y.S.Meng).Contents lists available at SciVerse ScienceDirectMaterials Science and Engineering Rj ou r n a l h o m e p a g e:w w w.e l s e v i e r.co m/l oc a t e/m s e r0927-796X/$–see front matterß2012Elsevier B.V.All rights reserved./10.1016/j.mser.2012.05.003are in the cathode sides initially and the battery system is assembled in ‘‘discharged’’status.While charging,Li ions are extracted from the cathode host,solvate into and move through the non-aqueous electrolyte,and intercalate into the anode host.Meanwhile,electrons also move from cathode to anode through the outside current collectors forming an electric circuit.The chemical potential of Li is much higher in the anode than in the cathode,thus the electric energy is stored in the form of (electro)chemical energy.Such process is reversed when the battery is discharging where the electrochemical energy is released in the form of electric energy.The cathode region and anode region are separated by the separator,a micro-porous membrane that allows the electrolyte to penetrate and prevent shorting between the two electrodes.The electrolyte should be ionically conducting and electronically insulating in principle,however the actual properties of the electrolyte is much more complicated.During the ﬁrst cycle,a so-called solid–electrolyte-interphase (SEI)layer will be formed on the surface of electrodes due to the decomposition of organic electrolyte at extreme voltage range (typically <1.2V or >4.6V).In current LIB technology,the cell voltage and capacities are mainly determined by the cathode material that is also the limiting factor for Li transportation rate.The developments of cathode materials therefore become ex-tremely crucial and receive much attention in recent decade.Since 1980when the LiCoO 2was demonstrated ﬁrstly as a possible cathode material for rechargeable lithium battery [1],the transition metal intercalation oxides have caught the major research interests as the LIB cathodes [2–8].Categorized by structure,the conventional cathode materials include layered compounds LiMO 2(M =Co,Ni,Mn,etc.),spinel compounds LiM 2O 4(M =Mn,etc.),and olivine compounds LiMPO 4(M =Fe,Mn,Ni,Co,etc.).Most of the researches are performed on these materials and their derivatives.New structure intercalation materials such as silicates,borates and tavorites are also gaining increasing attentions in recent years.During the materials optimization and development,following designing criterions are often considered:(1)energy density;(2)rate capability;(3)cycling performance;(4)safety;(5)cost.The energy density is determined by the material’s reversible capacity and operating voltage,which are mostly determined by the material intrinsic chemistry such as the effective redox couples and maximum lithium concentration in active materials.For rate capability and cycling performances,electronic and ionic mobilities are key determining factors,though particle morphologies are also important factors due to the anisotropic nature of the structures and are even playing a crucial role in some cases.Therefore materials optimizations are usually made from two important aspects,to change the intrinsic chemistry and to modify the morphology (surface property,particle size,etc.)of the materials.Fig.2compares the gravimetricenergy densities of different cathode materials that are currently under investigations.While some materials such as LiFeBO 3and LiFeSO 4F are already approaching their theoretical energy densi-ties,for other materials including conventional layered and spinel compounds,signiﬁcant gaps are still present between their theoretical and practical energy densities.The materials with promising theoretical properties have high potentials as the candidates of future generation LIB cathode,therefore are under intensive studies.For certain materials such as the LiFePO 4olivine,signiﬁcant property improvements have been achieved during the past decade with assistance of newly developed technologies.To review and summarize those researches could provide inspiring perspectives for further material optimizations.In this review,we will discuss the recent research progress in the past decade of different cathode materials following the structural category,and modiﬁcations on both chemistry and morphology will be discussed.yered compounds LiMO 2The ideal structure of layered compound LiMO 2is demonstrat-ed in Fig.3.The oxygen anions (omitted for clarity in the ﬁgures)form a close-packed fcc lattice with cations located in the 6-coordinated octahedral crystal site.The MO 2slabs and Li layers are stacked alternatively.Although the conventional layered oxide LiCoO 2has been commercialized as the LIB cathode for twenty years,it can only deliver about 140mAh/g capacity which is half of its theoretical capacity.Such limitation can be attributed to the intrinsic structural instability of the material when more thanhalfFig.1.Working principles of LIB(charging).Fig.2.Theoretical and practical gravimetric energy densities of different cathodematerials.Fig.3.Crystal structure of layered LiMO 2(blue:transition metal ions;red:Li ions).(For interpretation of the references to color in this ﬁgure legend,the reader is referred to the web version of the article.)B.Xu et al./Materials Science and Engineering R 73(2012)51–6552of the Li ions are extracted.On the other hand,the presence of toxic and expensive Co ions in LiCoO 2has introduced the environmental problem as well as raised the cost of the LIB.The research focusing on layered compounds,therefore have moved from LiCoO 2to its derivatives in which Co ions are partially/fully substituted by more abundant and environmental friendly transition metal ions,such as Ni and Mn.The approaches include mixing the LiNiO 2and LiMnO 2with 1:1ratio,forming layered LiNi 0.5Mn 0.5O 2,and the formation of Li–Co–Ni–Mn–O layered compound (so-called NMC type materials).Good electrochemical data of LiNi 0.5Mn 0.5O 2was ﬁrstly reported in 2001by Ohzuku et al.[9].Fig.4shows the typical electrochemical performance of LiNi 0.5Mn 0.5O 2[10,11].The charge/discharge voltages of this material are around 3.6–4.3V where Ni 2+/Ni 4+act as the redox couple as conﬁrmed from in situ X-ray absorption spectroscopy (XAS)study [12].Various methods including X-ray and neutron diffraction,nuclear magnetic reso-nance (NMR)spectroscopy,transmission electron microscope (TEM)and ﬁrst-principles calculations [13–18]have been per-formed to investigate the structural change and local cation distribution of this material.The results showed that different from classic layered material composed of pure Li layer and pure MO 2slab,8–10%Ni ions are usually found in the Li layer of LiNi 0.5Mn 0.5O 2synthesized by solid state or sol–gel synthesis methods.It was suggested that such Li–Ni interlayer mixing might be partially reduced under high voltage (>4.6V)[19,20].For the transition metal layer,a ﬂower-like in-plane cation ordering that Li in transition metal layer is surrounded sequentially by Mn rings and Ni rings was suggested by ﬁrst-principles calculations [16]and conﬁrmed by experiments [14,15].With MO 2slabs pined together by the Ni in Li layer,larger reversible capacity ($200mAh/g)can be obtained in LiNi 0.5Mn 0.5O 2at low rate (C /20)with little capacity fading even after 100cycles,therefore the energy density can be signiﬁcantly improved.The material structure is thermally stable until $3008C,above which oxygen release and material decom-position would occur [21].Structural change including migration of transition metal ions to Li layer at high temperature was also reported both experimentally and computationally [21,22].However,with large amount of un-removable Ni in the Li layers blocking Li diffusion pathways,the Li mobility of the materials is negatively affected.The Li diffusion coefﬁcient in LiNi 0.5Mn 0.5O 2is reported to be lower than that in LiCoO 2by one magnitude of order [23],resulting in the low rate capability of LiNi 0.5Mn 0.5O 2.It was also reported by Kang et al.[11]that the Li–Ni exchange is reduced to $4%in LiNi 0.5Mn 0.5O 2synthesized by ion-exchange method,therefore the rate capability can be signiﬁcantly improved (shown in Fig.4(c)and (d)).Considering the Li–Ni disorder being major factor affecting the material rate capability,attempts to create new compounds of LiCo x Ni y Mn 1Àx Ày O 2are motivated.While good electrochemical performance of LiCo 1/3Ni 1/3Mn 1/3O 2was already reported in 2001by Ohzuku et al.[24],the importance of the series of Li–Co–Ni–Mn–O material is more recognized as the presence of Co ions can help to reduce the amount of defect Ni in Li layer.The LiCo 1/3Ni 1/3Mn 1/3O 2layer compound can be regarded as the solid solution of LiCoO 2,LiNiO 2and LiMnO 2.LiCo 1/3Ni 1/3Mn 1/3O 2deliver similar reversible capacity with LiNi 0.5Mn 0.5O 2.Their voltage proﬁle are also similar in shape,but the operation voltage window of LiCo 1/3Ni 1/3Mn 1/3O 2can be extended to 3.6–4.7V.The material’s typical electrochemical performance is shown in Fig.5[25,26].With additional Co ions existing in the structure,the Li–Ni interlayer mixing can be much reduced to 1–6%[27–32].Though certain superlattice in the transition metal layer could be obtained from computations [33,34],diffraction [28]and NMR study [35],it is suggested that only short-range ordering can be found and Li in transition metal layer is surrounded primarily by Mn ions.The changes of transition metal valence state following cycling were investigated experimentally and computationally [29,30,32,34–36].In general,it is believed that,the transition metal ions are oxidized in sequence of Ni 2+to Ni 3+,Ni 3+to Ni 4+,Co 3+to Co4+Fig.4.Performance of layered LiNi 0.5Mn 0.5O 2:(a)compositional phase diagram,(b)cycling performance [10],(c)rate performance of LiNi 0.5Mn 0.5O 2synthesized by ion-exchange method [11],and (d)rate performance of LiNi 0.5Mn 0.5O 2synthesized by solid state method [11].B.Xu et al./Materials Science and Engineering R 73(2012)51–6553during charging.Mn 4+ions keep unchanged.However,due to the overlap of oxygen 2p band M 3+/4+band,at the end of charging,part of the electrons were also removed from the oxygen ions,causing possible oxygen release at high charging voltages (>4.5V).As shown in Fig.5(c)and (d),with certain morphology modiﬁcations,$90%of the capacity can be retained after 200cycles at room temperature and 84%of the capacity can be retained when at a discharge rate as high as 20C [26].The LiCo 1/3Ni 1/3Mn 1/3O 2compound also shows relatively good performance at elevated temperatures.It was reported [25]that more than 80%capacity can be retained at 558C and half of the capacity can still be achieved when the operation temperature is raised to 958C.We also want to point out that the reduced amount Co can be used to achieve the same beneﬁts.It has been shown by Li et al.that as little as 20%Co (LiCo 0.2Ni 0.4Mn 0.4)can lead to excellent electrochemical perfor-mance [37].While the introduction of Co ions into LiNi 0.5Mn 0.5O 2could improve the material stability,the introduction of extra Li ions,on the other hand,could improve the material capacity.A series of Li-rich layered oxides Li[Li 1/3À2x /3Ni x Mn 2/3Àx /3]O 2therefore are created by making a compound between Li 2MnO 3and LiNi 0.5Mn 0.5O 2to achieve higher capacity beyond the limitation of one Li ion per MO 2formula.With excess Li ions introduced,the theoretical capacity of this series of materials can be increased to more than 300mAh/g.The compound Li[Li 1/9Ni 1/3Mn 5/9]O 2is ﬁrstly reported by Lu et al.in 2001[17].Different compositions (x =1/6,1/5,1/4,1/3,5/12,etc.)of this series of materials were studied by the same group later and similar electrochemical performances were shown [38].For this series of materials,the pristine samples are solid solutions phase following the layered structure in general,although it was reported that some of the spinel feature might be observed when x !1/3being heated to 6008C [38].The Li–Ni interlayer mixing is usually much smaller than that in LiNi 0.5Mn 0.5O 2.With x increasing,the Li–Ni interlayer mixing also increases while the c/a ratio decreases.The structuredetails and cation ordering were investigated by diffraction,TEM,NMR spectroscopy and ﬁrst-principles calculations [39–44].The results show that the excess Li ions were located in the transition metal layer,surrounded mostly by 5or 6Mn 4+ions.The Ni/Mn zigzag ordering is regarded as another competent driving force for the in-plane cation ordering,and actual material may reﬂect a combination of these two orderings [39].A typical voltage proﬁle of Li[Li 1/9Ni 1/3Mn 5/9]O 2for the ﬁrst cycle charging is shown in Fig.6(c)[45].It is composed of a slopy region from the open circuit voltage to $4.5V followed by a plateau region between 4.5V and 4.6V.Such plateau,however,does not appear in following cycles causing a large ﬁrst cycle irreversible capacity.While it is generally agreed that the slopy region originates from the oxidization of Ni 2+to Ni 4+,the mechanism of the anomalous high voltage capacity is still under debate.Several studies [46–48]proposed the mecha-nism of oxygen loss accompanied by Li removal,while other studies proposed the mechanisms involving surface reaction with electrode/electrolyte reduction [49]and/or hydrogen exchange [50].This series of materials deliver the highest reversible capacity (>250mAh/g)for current intercalation cathode materials,but only with low rate (C /50).It was reported that the excess Li ions in transition metal layer are electrochemically active and will migrate to Li layer becoming stable tetrahedral ions during cycling [44,51].Recent study [44]also suggested a possible layer to defective-spinel phase transformation happening near the material surface with signiﬁcant migration of transition metal ions to lithium layer.These un-removable ions in lithium layer will block the lithium diffusion pathways,therefore may be one of the reasons that cause the low Li chemical diffusion coefﬁcient in the plateau region [43]and the intrinsic poor rate capability of this series of materials.A recent study also suggested that the material rate and temperature performances may also be affected by the particle size as shown in Fig.6[45].For the Li-rich layered compounds,the surface characteristics can signiﬁcantly affect the material electrochemicalperformanceFig.5.Performance of layered LiCo 1/3Ni 1/3Mn 1/3O 2:(a)compositional phase diagram,(b)cycling performance at elevated temperatures [25],(c)cycling performance at room temperature [26],and (d)rate performance at room temperature [26].B.Xu et al./Materials Science and Engineering R 73(2012)51–6554especially for theﬁrst-cycle irreversible capacity and rate capability.Different surface modiﬁcations,therefore,have been applied to this series of materials as further optimizations. Common coating materials applied include different types of metal oxides(Al2O3,Nb2O5,Ta2O5,ZrO2,and ZnO)[52–54],metal ﬂuorides AlF3[55,56]and other polyanion compounds such as MPO4(M=Al,Co)[57,58].A systematic research on the metal oxides coating was performed by Myung et al.in2007[54].The area-speciﬁc impedance(ASI)results showed that during theﬁrst cycle,the ASI of un-coated samples dramatically increased,while the ASI of all the coated samples hardly changed.The cycling performance and rate capabilities of the materials were improved especially when coated with Al2O3.It was claimed that during the initial cycling,the oxide coating reacted with the electrolyte forming a solid stableﬂuoride layer which protected the active materials from further HF scavenger.Similar effect was also reported for AlF3coating[56].In addition,the coating can suppress the oxygen loss occurred in the active materials,therefore can also improve the material thermal stability[55,59].The coating morphology is also under optimization.Double layer coating combing two or more coating materials has been developed as well [58].Another approach involves the construction of composition gradient from surface to bulk and forming core–shell structured particles.The structure is shown in Fig.7[60].By introducing composition gradient,the performance of‘‘core’’active materials can be maintained,while the less active‘‘shell’’materials can actasFig.7.Demonstration of core–shell structured particles[60].Fig.6.Performance of layered Li[Li1/3À2x/3Ni x Mn2/3Àx/3]O2:(a)compositional phase diagram,(b)cycling performance(x=1/3)[45],and(c)rate and temperature performance (x=1/3)[45].B.Xu et al./Materials Science and Engineering R73(2012)51–6555a buffer layer and help improve the material performance in surface.Although there are general hypotheses of why the coating materials can improve the active materials’performance,the detailed mechanism is still unknown and under intensive investigations.In summary,the layered oxides LiMO2can deliver high capacities after activation at high voltages,therefore leading to promising energy densities.However,their practical reversible capacities are usually limited by the intrinsic structural instability at low lithium concentrations and high voltages,causing reduced efﬁciency of the active material utilization.Besides,for the cobalt-free lithium nickel manganese oxides,the intrinsic low rate capability has become the bottleneck problem impeding the commercialization of these materials.3.Spinel compounds LiM2O4The structure of LiM2O4spinel is shown in Fig.8.The oxygen framework of LiM2O4is the same as that of LiMO2layered structure.M cations still occupy the octahedral site but1/4of them are located in the Li layer,leaving1/4of the sites in transition metal layer vacant.Li ions occupy the tetrahedral sites in Li layer thatshare faces with the empty octahedral sites in the transition metal layer.The structure is based on a three-dimensional MO2host and the vacancies in transition metal layer ensure the three-dimensional Li diffusion pathways.The spinel LiMn2O4was proposed as the cathode of the lithium ion battery by Thackeray et al.in1983[61–63],but the material was found to encounter sever capacity fading problem.Two reasons have been considered as the main sources for the capacity fading:(1)dissolution of Mn2+ into the electrolyte generated by the disproportional reaction2 Mn3+!Mn4++Mn2+[64,65]and(2)generation of new phases during cycling and the related micro-strains[64,66].Substituting Mn with other metal ions has been used as an important approach to improve cycling performance of spinel materials.Multiple dopants including inactive ions such as Mg,Al,and Zn[67–69],ﬁrst row transition metal ions such as Ti,Cr,Fe,Co,Ni,and Cu[70–74] and rare earth metal ions such as Nd and La[75–77]have been investigated and LiNi0.5Mn1.5O4shows the best overall electro-chemical performances among the above.LiNi0.5Mn1.5O4follows the spinel structure of LiMn2O4where Ni ions are located in the sites of Mn ions originally.With different synthesis conditions[78,79],LiNi0.5Mn1.5O4could possess two different structural symmetries,the ordered structure with space group P4332and the disordered structure with space group Fd¯3m. In ordered LiNi0.5Mn1.5O4,Ni ions occupy4b sites and Mn ions occupy12d sites forming an ordered pattern,while in disordered LiNi0.5Mn1.5O4,Ni and Mn ions are randomly distributed in16d sites.In stoichiometric LiNi0.5Mn1.5O4,the valence of Ni ions is2+ pushing all Mn ions to Mn4+.Comparing to LiMn2O4spinel,the redox couple of LiNi0.5Mn1.5O4is switched from Mn3+/Mn4+to Ni2+/ Ni4+and the voltage is lifted from 4.1V to 4.7V.Such high discharge voltage not only enlarges the energy density but also makes the material capable to be coupled with anode materials which have better safety but relatively higher voltage(Li4Ti5O12, etc.).However,phase-pure LiNi0.5Mn1.5O4is difﬁcult to synthesize because impurities such as nickel oxides and/or lithium nickel oxides usually exist[78,80,81].As an alternative approach,the off-stoichiometric material LiNi0.5Mn1.5O4Àx which adopts the disor-dered structure is synthesized and the performances of these two materials are compared in Fig.9[71].InLiNi0.5Mn1.5O4Àx,there are small amount of Mn3+ions exist as the charge compensation of oxygen loss.The small voltage plateau$4V for LiNi0.5Mn1.5O4Àx therefore is attributed to the Mn3+/Mn4+couple.Different from other doped spinel materials the voltage proﬁle of which is usually composed of two distinct plateaus,there is only oneﬂat plateauatFig.9.Performance of high voltage spinel LiNi0.5Mn1.5O4[71].Fig.8.Crystal structure of spinel LiM2O4(blue:transition metal ions;red:Li ions).(For interpretation of the references to color in thisﬁgure legend,the reader isreferred to the web version of the article.)B.Xu et al./Materials Science and Engineering R73(2012)51–6556$4.7V for LiNi0.5Mn1.5O4,although in LiNi0.5Mn1.5O4Àx,a small voltage step appears at half lithium concentration.The theoretical capacity of LiNi0.5Mn1.5O4is calculated as147mAh/g,and more than140mAh/g reversible capacity can be obtained experimen-tally.With most of the Mn ions keeping Mn4+unchanged during cycling,both ordered and disordered LiNi0.5Mn1.5O4exhibit good cycling performance for lower rate capability that there is little capacity fading after50cycles in room temperature(Fig.9(a)and (c))and elevated temperature(Fig.9(e)).However,their rate capabilities still need to be improved.The disordered LiNi0.5Mn1.5O4Àx shows better rate capability than ordered LiNi0.5Mn1.5O4for the material electronic conductivity is enhanced with the small amount of Mn3+ions.The ionic conductivity is regarded as another rate-limiting factor.The Li diffusion coefﬁcient of LiNi0.5Mn1.5O4was reported to in a wide range between 10À10cm2/s and10À16cm2/s depending on different compositions and material morphologies[82–84].As a further optimization approach,doping small amount of metal ions into LiNi0.5Mn1.5O4forming‘‘bi-doped’’spinel has been widely used,and certain properties of LiNi0.5Mn1.5O4can be further improved.A recent review paper by Yi et al.[85]has summarized the effect of different doping ions including both cation substitu-tion and anion substitution.By doping other transition metal ions such as Fe,Cr,and Ti into ordered LiNi0.5Mn1.5O4,the impurity phases may be limited and the cation disorder could be enhanced [86–90].It was also reported that dopant such as Co and Cu may enhance the material electronic conductivity and/or lithium diffusion coefﬁcient[91,92].These enhancements therefore could further improve the material cycling performance and rate capability(Fig.10[92]).The doped bivalence metal ions such as Cumay also shift the4.7V plateau to even higher voltage therefore could further improve the energy density.However,for most of the dopants,the high voltage capacity is shortened and the overall reversible capacity is reduced.It was reported[86,92,93]that some doped ions such as Fe,Cu,Al,and Mg tend to occupy the tetrahedral sites and become inactive ions,not only reducing the capacity but also blocking the lithium diffusion pathways,therefore may impose a negative effect to the material performances.Apart from chemistry modiﬁcation,size minimization is also reported as an effective approach to improve the material rate capability.Highly crystalline LiNi0.5Mn1.5O4nano-sized particles can be successfully synthesized through different methods [79,94,95].It was shown that the bulk properties of nano-sized particles are generally the same compared to micro-sized particles, although their surface areas increase causing increasing surface reactions.However,the small size shortened the Li diffusion length inside the active materials thus largely enhanced the material ionic conductivity.The rate capability of the nano-sized LiNi0.5Mn1.5O4 therefore is highly improved as an overall effect[95,96].It is important to point out the volumetric energy density(Wh/L) suffers from the size minimization greatly as most of the nano-sized materials do not have the optimized packing scheme yet.In summary,the high voltage spinel material is promising due to its high energy density,perfect structural stability and good cycling performance under certain material modiﬁcations.The high voltage,however,is out of the voltage window of the current electrolyte,therefore causes the electrolyte decomposition and the formation of unstable SEI on composite cathode side during cycling.It is important to point out that the reversible capacity of this material is currently limited to0.5Li per MO2formula,which although is similar to the practical capacity of LiCoO2,still is signiﬁcantly lower comparing to the lithium nickel manganese layered compounds.4.Olivine compounds LiMPO4Despite the early works back in1980s[97,98],intensive studies on polyanion materials have not been conducted until recent ﬁfteen years.These materials are receiving growing attentions because of the inherent stability of the polyanion group,which can delay or minimize the oxygen loss happening in traditional layer and spinel oxides.Among all polyanion materials,olivine LiFePO4attracts the most interests due to its excellent electrochemical properties,as well as its low cost,non-toxicity,excellent thermal stability and environment friendliness.It wasﬁrst found by Goodenough and coworkers in1997[99,100].The structure of LiFePO4is shown in Fig.11.It contains slightly distorted hcp anion oxygen arrays with half of the octahedral sites occupied by Fe and one eighth by Li.The LiO6octahedra are edge-shared while the FeO6octahedra are corner-shared.Both of the LiO6and FeO6run parallel to the c axis and they alternate in the b direction.The a–c planes containing the Li atoms are bridged by PO4tetrahedral.Three different paths of Li diffusion were proposed[101,102]and computational studies suggested that the one along b axis is much more favored than other paths across the channels[101,102].In addition,in2008 Yamada et al.further conﬁrmed from experiments that the Li ion diffusion path along the(010)is a curved one dimensionchainFig.11.Crystal structure of olivine LiMPO4(blue:transition metal ions;yellow:P ions;red Li ions).(For interpretation of the references to color in thisﬁgure legend, the reader is referred to the web version of thearticle.)Fig.10.Rate performance of rate capability of LiNi x Cu y Mn2ÀxÀy O4[92].B.Xu et al./Materials Science and Engineering R73(2012)51–6557。

锂离子

锂离子电池摘要：锂离子电池与其他二次电池相比，具有比能量大、自放电小、质量轻、无记忆效应和环境友好等优点，因此决定了其在电动汽车、存储电源等方面极具发展前景[1];介绍了锂离子电池的电化学反应原理、特点及发展历程。

研究了锂离子电池的成型工艺、电解液以及化成进行，综述了锂离子电池的正极材料、负极材料及电解质材料的研究进展, 并展望了其发展前景。

关键词:锂离子；电极材料；制造技术受到能源危机和环境保护等因素影响，纯电动汽车做为一种新兴的交通工具，由于其能源利用率高、无排放、噪声小以及能量来源多样化等优点[2]成为汽车工业一个重要的研究领域。

传统的电动汽车存在续航里程有限、蓄电池使用寿命太短以及蓄电池尺寸和质量的制约等缺点，而锂离子电池具有高能量密度、高工作电压、无记忆效应、循环寿命长、无污染、质量轻、自放电小等特点[3]，已经应用在很多领域，如便携式电子产品、电动交通工具、大型动力电源和二次充电及储能领域等［4］。

可以很好地解决以上问题，所以基于锂离子电池的电动汽车受到越来越多人们的关注，锂离子电池已经成为纯电动汽车的候选能源之一[5-6]。

1. 锂离子电池工作原理及特点锂离子电池是在锂二次电池基础上发展起来的一种新型充电电池, 它的正负极材料都是能发生锂离子嵌入——脱出反应的物质。

在充电态负极处于富锂态, 正极处于贫锂态, 随着放电的进行, 锂离子从负极脱嵌, 经过电解质嵌入正极, 放电时则以相反过程进行。

在充放电过程中, 锂离子在正负极间摇来摇去, 而无金属锂的析出, 因此, 锂离子电池又被称为“摇椅电池”[ 7] 。

显然这种电池的工作电压与构成电极的锂离子嵌入化合物的浓度有关, 用作电极的材料主要是过渡金属、锰的锂离子嵌入化合物和锂离子嵌入碳化合物。

锂离子电池的一般特点:( 1) 体积及质量比能量高; ( 2) 单电池的输出电压高, 约为4. 2 V; ( 3) 自放电率小; ( 4) 能在较高的温度下使用; ( 5) 对环境污染小。

原位XRD与中子衍射研究LiCoPO4磷酸钴锂

• Summary
The ‘black box’ problem
U, I I, U
There is some thing happening. But what?
In-situ probes
Structural probes that can be used in-situ:
• X-rays • Neutrons
charge-/dischargeprocess
particle do not break additional „internal“ interface formation
Domaines with different
Li-content
Neutron Tomography of commercial „18650“ Batteries
In-situ X-ray diffraction
–High-energy radiation –High intensity –Low (vertical) divergence
~ 1/γ
–Measurement of „thick“ samples –Time resolved measurements –High resolution
5.0
LiCoPO4 – In situ X-ray diffraction
200
LixCoPO4
020
<x>~0.4
discharge
<x>~0
charge
<x>=1.0
5.2
2θ/°
5.6
9.0
2θ/°
3 phases with broad coexistence ranges

顶板厚度对组合桥面板U_肋与顶板构造疲劳性能的影响

第 39 卷第 4 期2023 年8 月结构工程师Structural Engineers Vol. 39 ， No. 4Aug. 2023顶板厚度对组合桥面板U肋与顶板构造疲劳性能的影响江臣1苏庆田2，*郭赵元1周青3傅晨曦3（1.江苏省交通工程建设局，南京 210004； 2.同济大学土木工程学院，上海 200092；3.华设设计集团有限公司，南京 210014）摘要为研究钢顶板厚度对于正交异性组合桥面板U肋与顶板构造疲劳性能的影响，通过建立板壳-实体有限元模型，计算分析不同顶板厚度下正交异性钢桥面板中关键部位的受力情况。

计算结果表明，有无刚性铺装对车轮荷载作用下的桥面板受力状态以及U肋与钢桥面板焊缝焊趾处应力大小有显著的不同，采用刚性铺装的正交异性组合桥面板U肋与钢桥面板焊缝焊趾处应力幅值只有不采用刚性铺装桥面板的8%；改变钢顶板的厚度对正交异性组合桥面板U肋与钢桥面板焊缝焊趾处的疲劳强度影响甚微，采用12 mm的钢顶板也能满足U肋与钢桥面板焊缝焊趾处的疲劳强度要求。

关键词正交异性钢桥面板， U肋，顶板，疲劳性能，有限元法Effect of Plate Thickness on Fatigue Performance of Connection between Trough and Deck Plate in Steel-concreteComposite Bridge DeckJIANG　Chen1SU　Qingtian2，*GUO　Zhaoyuan1ZHOU　Qing3FU　Chenxi3（1.Jiangsu Provincial Transportation Engineering Construction Bureau， Nanjing 210004， China；2.College of Civil Engineering， Tongji University， Shanghai 200092， China；3.China Design Group Co.， Ltd.， Nanjing 210014， China）Abstract In order to study the effect of deck plate thickness on fatigue performance of connection between trough and deck plate in steel-concrete composite bridge deck，a shell-solid finite element model was established to calculate and analyze the stress distribution in an orthotropic steel-concrete bridge deck with different thicknesses of deck plate. The results show that the presence or absence of rigid pavement has a significant effect on the stress state of the bridge deck as well as the stress at the weld toe between U-rib and steel plate under vehicle loading. The stress amplitude at the weld toe between rib and plate is only 8% of that without rigid pavement. Changing the thickness of deck plate has little effect on fatigue strength at the weld toe between U-rib and steel plate， and a 12-mm-thick steel plate can also meet the fatigue requirements. Keywords orthotropic steel bridge deck， U-rib， bridge plate， fatigue performance， finite element method收稿日期：2022-05-02作者简介：江臣（1971-），男，研究员级高级工程师，主要从事交通工程建设管理工作。

Electrochemical and structural properties of xLi(2)M'O-3 center dot(1-x)LiMn0.5Ni0.5O2 eIectrodes

Electrochemical and Structural Properties of x Li2M′O3‚(1-x)LiMn0.5Ni0.5O2Electrodes for Lithium Batteries(M′)Ti,Mn,Zr;0e x E0.3)Jeom-Soo Kim,Christopher S.Johnson,John T.Vaughey,andMichael M.Thackeray*Electrochemical Technology Program,Chemical Engineering Division,Argonne National Laboratory,Argonne,Illinois60439Stephen A.HackneyDepartment of Metallurgical and Materials Engineering,Michigan Technological University,Houghton,Michigan49931Wonsub YoonBrookhaven National Laboratory,Upton,New York11973Clare P.GreyDepartment of Chemistry,State University of New York,Stony Brook,New York11794 Received December24,2003.Revised Manuscript Received March4,2004Electrochemical and structural properties of x Li2M′O3‚(1-x)LiMn0.5Ni0.5O2electrodes(M′)Ti,Mn,Zr;0e x e0.3)for lithium batteries are reported.Powder X-ray diffraction,lattice imaging by transmission electron microscopy,and nuclear magnetic resonance spectroscopy provide evidence that,for M′)Ti and Mn,the Li2M′O3component is structurally integrated into the LiMn0.5Ni0.5O2component to yield“composite”structures with domains having short-range order,rather than true solid solutions in which the cations are uniformly distributed within discrete layers.Li2TiO3and Li2ZrO3components are electrochemically inactive, whereas electrochemical activity can be induced into the Li2MnO3component above4.3V vs Li0.When cycled in lithium cells,x Li2MnO3‚(1-x)LiMn0.5Ni0.5O2electrodes with x)0.3 provide capacities in excess of300mA‚h/g over the range4.6-1.45V.IntroductionLayered lithium transition-metal-oxide electrodes Li x-MO2(M)Co,Ni,and Mn)become unstable at low lithium content when lithium cells are charged above 4.3V.At this potential,the oxygen activity at the electrode surface is increased to a level at which side reactions can occur,such as electrolyte oxidation or oxygen loss.For several years,following the discovery that acid treatment of Li2MnO3produces a novel layered Li2-x MnO3-x/2structure,1which on relithiation yields a compound that can be formulated as x Li2M′O3‚(1-x)-LiMO22(ignoring the H+ion exchange that occurs during acid treatment),1,3we have pursued an approach using an electrochemically inactive Li2M′O3component (e.g.,Mn,Ti,or Zr)to stabilize these layered LiMO2 structures.In particular,we have used this approach in attempts to prevent the conversion of layered LiMnO2 to spinel.4,5Our efforts to stabilize layered LiMn1-x Ni x O2electrodes followed reports by Pacific Lithium Ltd.that Li1.2Mn0.4Cr0.4O2electrodes(alternatively,x Li2MnO3‚(1-x)LiCrO2[M′)Mn;M)Cr;x)0.5])provided high capacity and good cycling stability at50°C when cells were charged and discharged between4.4and2.5V.6,7 Our more recent efforts were influenced by reports that LiMn0.5Ni0.5O2and related lithium manganese-nickel oxides can provide capacities reaching200mA‚h/g.8-11 In particular,our approach is similar to that of Dahn et al.10who have studied the Li[Ni x Li(1/3-2x/3)Mn(2/3-x/3)]O2 system(0<x<0.5),which can be written,alternatively, in composite notation as(1-2x)Li2MnO3‚(3x)LiMn0.5-*To whom correspondence should be addressed.Email:thackeray@ .(1)Rossouw,M.H.;Thackeray,M.M.Mater.Res.Bull.1991,26,463.(2)Rossouw,M.H.;Liles,D.C.;Thackeray,M.M.J.Solid State Chem.1993,104,464.(3)Tang,W.;Kanoh,H.;Yang,X.;Ooi,K.Chem.Mater.2000,12, 3271.(4)Johnson,C.S.;Korte,S.D.;Vaughey,J.T.;Thackeray,M.M.; Bofinger,T.E.;Shao-Horn,Y.;Hackney,S.A.J.Power Sources1999, 81-82,491.(5)Johnson,C.S.;Thackeray,M.M.The Electrochem.Soc.Inc., PV2000-36,2001,47.(6)Ammundsen,B.;Desilvestro,J.;Steiner,R.;Pickering,P.10th International Meeting on Lithium Batteries,Como,Italy,28May-2 June2000;Ext.Abstr.No.17.(7)Ammundsen,B.;Paulsen,J.Adv.Mater.2001,13,943.(8)Spahr,M.E.;Novak,P.;Schnyder,B.;Haas,O.;Nesper,R.J. Electrochem.Soc.1998,14,1113.(9)Ohzuku,T.;Makimura,Y.Chem.Lett.2001,744.(10)Lu,Z.;MacNeil,D.D.;Dahn,J.R.Electrochem.Solid State Lett.2001,4,A191.(11)Lu,Z.;Dahn,J.R.J.Electrochem.Soc.2002,149,A815.1996Chem.Mater.2004,16,1996-200610.1021/cm0306461CCC:$27.50©2004American Chemical SocietyPublished on Web04/20/2004Ni0.5O2(also for0<x<0.5).During the course of ourstudies,it was discovered that LiMn0.5Ni0.5O2andcomposite x Li2M′O3‚(1-x)LiMn0.5Ni0.5O2(M)Mn,Ti)electrode structures could accommodate an additionallithium within their structures at approximately1.5Vvs Li0without disturbing the layered transition-metalframework structure.12,13This reaction,which is revers-ible,greatly increases the capacity of the layeredelectrode.The structures of x Li2M′O3‚(1-x)LiMO2compounds(alternatively,in layered notation,Li(2+2x)/(2+x)M′2x/(2+x)-M(2-2x)/(2+x)O2)appear to be complex.For example,magicangle spinning(MAS)nuclear magnetic resonance(NMR)and X-ray absorption spectroscopy(XAS)dataof Li1.2Mn0.4Cr0.4O2electrodes(Li2MnO3‚LiCrO2)haveprovided evidence that the Li2MnO3and LiCrO2com-ponents do not form a random solid solution,but ratherthat the structures tend to have short-range order withlocal clustering of the lithium and transition-metalions.14,15Li2MnO3-like local environments or smalldomains were observed by NMR,which were electro-chemically inert.15Similar Li2MnO3-like local environ-ments were seen for Li[Ni x Li(1/3-2x/3)Mn(2/3-x/3)]O2for0e x e0.5,but the Li ions in these environments participated in the electrochemistry,suggesting that thedomains were smaller in size.16However,Dahn et al.17recently opposed the idea of Li2MnO3-like domains fromanalyses of XRD data and argued that the cations arearranged homogeneously within the layers of Li[Cr x-Li(1-x)/3Mn(2-2x)/3]O2for0e x e1(alternatively,(2-2x)-Li2MnO3‚(3x)LiCrO2)and Li[Ni x Li(1/3-2x/3)Mn(2/3-x/3)]O2for0e x<0.5(alternatively,(1-2x)Li2MnO3‚(3x)LiMn0.5-Ni0.5O2)structures,suggesting“a true solid solution phase.”In this paper,we address the ordering of these materials.For simplicity,we restrict the definition of structures in this paper,in most instances,to the composite notation x Li2M′O3‚(1-x)LiMO2rather than the equivalent layered notations used by Dahn et al.10 We define and discuss x Li2M′O3‚(1-x)LiMn0.5Ni0.5O2 electrode structures(M′)Mn,Ti,or Zr;0e x e0.3)as determined by powder X-ray diffraction(XRD),high-resolution transmission electron microscopy(HRTEM), and MAS NMR.Electrochemical properties of the electrodes at both high potential(4.6-2.0V)and at low potential(2.0-1.0V)vs a metallic lithium reference electrode are presented.The electrochemical properties of electrodes containing a Li2MnO3component are interpreted with reference to reports of the electro-chemical activity that can be induced into Li2MnO3by acid treatment1,18or by charging Li2MnO3electrodes directly in lithium cells.19To facilitate a discussion of the results,we first consider the structures of x Li2M′O3‚(1-x)LiMO2elec-trodes.Because of the structural similarities between Li2M′O3-and LiMO2-type compounds,it was proposed several years ago that it might be possible to use an electrochemically inactive Li2M′O3component,such as Li2MnO3,Li2TiO3,or Li2ZrO3,to stabilize LiMO2inser-tion electrodes(M)Co,Ni,Mn,and combinations thereof),particularly at low lithium loadings,by forming “composite”electrode structures,x Li2M′O3‚(1-x)Li-MO2.2,4,20,21In LiCoO2,LiNiO2,and LiMnO2,the lithium and transition-metal ions are located in octahedral sites in alternating layers(Figure1a),whereas in Li2MnO3 and Li2TiO3(in layered notation,Li[Li0.33Mn0.67]O2and Li[Li0.33Ti0.67]O2,respectively),one-third of the transi-tion-metal ions are replaced by lithium(Figure1b).22 In Li2ZrO3,the structure is not layered because each cation layer contains Li+and Zr4+ions in a2:1ratio (Figure1c).23The composition of x Li2M′O3‚(1-x)LiMO2electrodes for the complete range of x(0e x e1)is defined by the Li2M′O3-LiMO2tie-line in the schematic phase diagram of the Li2M′O3-MO2-LiMO2-Li2MO2system in Figure 2.In ideal layered Li2M′O3and LiMO2structures,such as Li2MnO3and LiCoO2,respectively,the cations in the transition-metal layers are perfectly ordered in layers between close-packed oxygen planes.Cation ordering(12)Johnson, C.S.;Kim,J.-S.;Kropf, A.J.;Kahaian, A.J.; Vaughey,J.T.;Thackeray,mun.2002,4,492.(13)Johnson, C.S.;Kim,J.-S.;Kropf, A.J.;Kahaian, A.J.; Vaughey,J.T.;Fransson,L.M.;Edstrom,K.;Thackeray,M.M.Chem. Mater.2003,15,2313.(14)Ammundsen,B.;Paulsen,J.;Davidson,I.;Liu,R.-S.;Shen,C.-H.;Chen,J.-M.;Jang,L.-Y.;Lee,J.-F.J.Electrochem.Soc.2002, 149,A431.(15)Pan,C.J.;Lee,Y.J.;Ammundsen,B.;Grey,C.P.Chem.Mater.2002,14,2289.(16)Yoon,W.S.;Paik,P.;Yang,X.-Q.;Balasubramanian,M.; McBreen,J.;Grey,C.P.Electrochem.Solid State Lett.2002,5,A263.(17)Lu,Z.;Chen,Z.;Dahn,J.R.Chem.Mater.2003,15,3214.(18)Paik,Y.;Grey,C.P.;Johnson,C.S.;Kim,J.-S.;Thackeray, M.M.Chem.Mater.2002,14,5109.(19)Robertson,A.D.;Bruce,P.G.Chem.Mater.2003,15,1984.(20)Kim,J.-S.;Johnson,C.S.;Thackeray,M.M.Electrochem. Commun.2002,4,205.(21)Johnson,C.S.;Thackeray,M.M.In Proc.198th ECS Meeting, Phoenix,AZ,2000;Abstr.No.67.(22)Strobel,P.;Lambert,A.B.J.Solid State Chem.1988,75,90.(23)Hodeau,J.L.;Marezio,M.;Santoro,A.;Roth,R.S.J.Solid State Chem.1982,45,170.Figure1.Schematic illustrations of the structures of(a) LiMO2(M)Mn,Ni,Co);(b)Li2MnO3and Li2TiO3;(c)Li2-ZrO3;(d)Li2MO2(M)Mn,Ni).positional phase diagram of the Li2MnO3-MO2-LiMO2-Li2MO2system.xLi2M′O3‚(1-x)LiMn0.5Ni0.5O2for Li Batteries Chem.Mater.,Vol.16,No.10,20041997occurs in the Li2Mn layers of Li2MnO3,which is drivenby the large differences in charge of the monovalent andtetravalent Li and Mn ions,respectively.Therefore,when synthesizing compounds that fall on a Li2M′O3-LiMO2tie-line,it seems highly unlikely that perfectlyordered structures or superstructures with a homoge-neous distribution of monovalent,divalent,trivalent,and/or tetravalent cations in the transition-metal layerswill result,except at certain crystallographically allowedvalues of x.Furthermore,because it is extremely dif-ficult to control the precise stoichiometry of lithiatedmetal oxide products at high synthesis temperatures,it is also highly likely that there will be variations inlocal composition in their structures,the extent of whichwill be dependent on(1)localized concentrations oflithium and transition-metal ions at the time of syn-thesis,(2)the oxygen content of the sample,and(3)thefree energy of formation of various thermodynamicallyfavored phases that will compete with one another atthe synthesis temperature.The LiMn0.5Ni0.5O2structure that has been used asthe LiMO2component in the x Li2M′O3‚(1-x)LiMn0.5-Ni0.5O2composite electrodes(M′)Mn,Ti,or Zr;0e xe0.3)of this investigation is of particular interest because the manganese and nickel ions adopt tetrava-lent and divalent oxidation states,respectively,ratherthan a trivalent state.10(Note that both Mn3+(d4)andNi3+(d7)ions are Jahn-Teller active and that relaxationto undistorted states can therefore be achieved simul-taneously by electron transfer from a Mn3+ion to aneighboring Ni3+ion.)Because manganese and nickelions exist in the transition-metal layers of the LiMn0.5-Ni0.5O2structure in a1:1ratio,there is no simple3-fold(point group)symmetry in the transition-metal layersas there is in Li2MnO3.It is therefore impossible forevery nickel ion to be surrounded by six manganese ionsand vice versa.Therefore,if an ordered structure exists,there has to be another arrangement of the nickel andmanganese ions.24Ceder et al.25have predicted by first-principles calculations that a zigzag arrangement ofnickel and manganese has the lowest formation energy.However,no experimental crystallographic data has,toour knowledge,been provided that confirms the exist-ence of the zigzag phase,presumably in part becausematerials synthesized to date contain Li ions in additionto Ni and Mn ions in the predominantly transition-metallayers.Yoon et al.16,26have demonstrated by MAS NMRthat even in a LiMn0.5Ni0.5O2sample evidence could befound of local clustering and cation ordering in thetransition-metal layers(i.e.,Li2MnO3-type character),which was consistent with the presence of Ni ions inthe lithium layers and vice versa,which is common inlayered lithium-nickel oxides.27The spectra were ra-tionalized by using a model based on Li2MnO3-likeordering,the Ni2+substituting in sites that minimizethe number of Li+-Ni2+contents in the transition-metal layers.28Some evidence for Ni2+clustering,presumablyin LiNi0.5Mn0.5O2regions,was also observed for0.3e x <0.5.28Ohzuku has predicted by first-principles calculationsthat a 3× 3supercell should exist in a LiMn0.33-Ni0.33Co0.33O2structure.29Dahn et al.17claim fromanalyses of powder X-ray diffraction data that a similar 3× 3supercell exists for Li[Ni x Li(1/3-2x/3)Mn(2/3-x/3)]O2 samples(0<x e0.5)(alternatively,(1-2x)Li2MnO3‚(3x)LiMn0.5Ni0.5O2)but assert that because the latticeparameters change continuously with x,the cations arehomogeneously distributed in the transition-metal lay-ers over a wide range of x.During our initial studies of LiMn0.5Ni0.5O2and0.03Li2TiO3‚0.97LiMn0.5Ni0.5O2electrodes,it was dis-covered that an additional lithium ion could be insertedinto the LiMn0.5Ni0.5O2component below2V vs Li0toyield a Li2Mn0.5Ni0.5O2structure without destroying thelayered arrangement of the electrode structure(Figure1d).12,13In Li2Mn0.5Ni0.5O2,which is isostructural withLi2MnO230and Li2NiO2,31the manganese and nickelions maintain their octahedral configuration,whereasthe lithium ions fully occupy the tetrahedral sites inadjacent layers.The Li2MO2composition is included inthe compositional phase diagram in Figure2because,in principle,such structures can significantly increasethe capacity of composite x Li2M′O3‚(1-x)LiMO2elec-trodes at potentials below2V vs Li0.Experimental SectionSynthesis of x Li2M′O3‚(1-x)LiMn0.5Ni0.5O2Electrodes. Electrode materials defined by the general composition x Li2M′O3‚(1-x)LiMn0.5Ni0.5O2(M′)Ti,Mn,and Zr)were prepared for x)0.03,0.07,0.14,and0.30.Electrodes in which M′)Ti or Zr were prepared by reacting the appropriate amounts of anatase(TiO2,Aldrich),titanium isopropoxide(Ti-[OCH(CH3)2]4,Aldrich),or zirconium isopropoxide2-propanol complex(Zr[OCH(CH3)2]4‚(CH3)2CHOH,Aldrich)with lithium hydroxide monohydrate(LiOH‚H2O,Aldrich)and Ni0.5Mn0.5-(OH)2.Electrodes in which M′)Mn were prepared directly from Ni1-x Mn x(OH)2and LiOH‚H2O using the required Li:Mn: Ni ratio.The Ni0.5Mn0.5(OH)2and Ni1-x Mn x(OH)2precursors were prepared in-house by precipitation from a basic LiOH solution of Ni(NO3)2and Mn(NO3)2(pH∼11).The reagents were intimately mixed in an acetone slurry,dried in an oven overnight,and subsequently fired at480°C for12h and then at900°C for10h in air.Thereafter,the samples were rapidly quenched(also in air).Powder X-ray diffraction patterns of the products were collected on a Siemens D5000powder diffractometer with Cu K R radiation between10°and80°2θat a scan rate of0.6°2θ/ttice parameters of the phases were determined by Rietveld profile refinements of the X-ray diffraction patterns.High-Resolution Transmission Electron Microscopy. High-resolution images of x Li2TiO3‚(1-x)LiMn0.5Ni0.5O2and x Li2MnO3‚(1-x)LiMn0.5Ni0.5O2structures were obtained for x )0,0.14,and0.30.Samples were prepared for the electron microscope by a standard procedure described elsewhere.32The images were collected on a JEOL-JEM4000FEX-1transmis-(24)Islam,M.S.;Davies,R.A.;Gale,J.D.Chem.Mater.2003,15, 4280.(25)Ceder,G.;Carlier,D.;Grey,C.P.;Gorman,J.P.;Reed,J. Lithium Battery Discussion Workshop,Arcachon,France,14-19 September2003;Ext.Abstr.No.53.(26)Yoon,W.-S.;Grey,C.P.;Balsubramanian,M.;Yang,X.-Q.; McBreen,J.Chem.Mater.2003,15,3161.(27)Delmas,C.In Lithium Batteries,New Materials,Developments, and Perspectives;Pistoia,G.,Ed.;Elsevier:Amsterdam,1994; p457.(28)Yoon,W.-S.;Iannopollo,S.;Grey,C.P.;Carlier,D.;Gorman, J.;Reed,J.;Ceder,G.Electrochem.Solid State Lett.,in press.(29)Koyama,Y.;Tanaka,I.;Adachi,H.;Makimura,Y.;Ohzuku, T.J.Power Sources2003,119-121,644.(30)David,W.I.F.;Goodenough,J.B.;Thackeray M.M.;Thomas, M.G.S.R.Rev.Chem.Miner.1983,20,636.(31)Dahn,J.R.;von Sacken,U.;Michal,C.A.Solid State Ionics 1990,44,87.(32)Shao-Horn,Y.;Hackney,S.A.;Cornilsen,B.C.J.Electrochem. Soc.1997,144,3147.1998Chem.Mater.,Vol.16,No.10,2004Kim et al.sion electron microscope under an accelerating voltage of200 keV.The images were produced without an objective aperture to produce phase contrast.Fourier transform analysis was carried out using Scion Image(Scion Corp.,Release Beta4.0.2) on a64×64pixel section of the digital image file.Magic Angle Spinning Nuclear Magnetic Resonance. 6Li MAS NMR data of x Li2TiO3‚(1-x)LiMn0.5Ni0.5O2samples (x)0,0.14,and0.30)were collected at29.5MHz on a CMX-200spectrometer,with a spinning speed of38kHz and rotor synchronized spin-echoes(π/2-τ-π-τ-acq.;τ)one rotor period;π/2)2µs).The spectra were referenced to1M LiCl at0ppm.Electrochemical Measurements.Electrodes were fabri-cated by intimately mixing85wt%of the appropriate x Li2M′O3‚(1-x)LiMn0.5Ni0.5O2powder,8wt%poly(vinylidenedifluoride)(PVDF)binder(Kynar,Elf-Atochem),and7wt% acetylene black(Cabot)in a1-methyl-2-pyrrolidinone(NMP) solvent(Aldrich,99+%).The mixed slurry was cast onto an aluminum foil current collector and dried at75°C for3-5h. The electrode laminates were dried further at70°C under vacuum overnight.Disk electrodes with an area of1.6cm2 were punched from the laminates.The electrodes were evalu-ated in coin-type cells(size2032,Hohsen Corp.)with lithium foil(FMC Corp.)as the counter electrode,a polypropylene separator(Celgard),and an electrolyte solution consisting of 1M LiPF6in a1:1ethylene carbonate:diethyl carbonate solvent mixture(Merck).Cells were constructed inside a helium-filled glovebox(<5ppm,H2O and O2).Cells were cycled galvanostatically using a Maccor Series2000control unit at a current rate of0.1mA/cm2.ResultsXRD Patterns.The powder XRD patterns of x Li2-M′O3‚(1-x)LiMn0.5Ni0.5O2electrodes for M′)Mn,Ti, or Zr and0e x e0.3are provided in Figure3.Because of the paucity of data,lattice parameters were refined using the high-symmetry trigonal space group,R3h m, rather than the low-symmetry monoclinic space group, C2/m,to monitor the variations in unit cell dimensions as a function of composition,x(Table1).The XRD patterns of the x Li2MnO3‚(1-x)LiMn0.5Ni0.5O2(x)0.00, 0.03,0.14,and0.30)composite structures in Figure3a are similar.Figure3a shows a slight increase in intensity of the peaks at20-23°2θas the Li2MnO3 component in the structure increases,which is consis-tent with an increasing amount of lithium in the transition-metal layers.Our XRD patterns and the changes in lattice parameter are consistent with those reported by Dahn et al.17for Li[Ni x Li(1/3-2x/3)Mn(2/3-x/3)]O2 electrodes;the smooth variation in lattice parameter with composition was interpreted by this group as suggesting that these materials are structurally and chemically homogeneous.Overall,the average unit cell parameters and unit cell volumes decrease marginally with increasing Mn content,which is consistent with the difference in the ionic radii between Mn4+(0.54Å) and Ni2+(0.69Å)ions in octahedral coordination.33By contrast,the average lattice parameters and unit cell volumes of x Li2TiO3‚(1-x)LiMn0.5Ni0.5O2products syn-thesized from a Ti[OCH(CH3)2]4precursor increase in value with increasing x;in this case,the trend is consistent with the larger ionic radius of Ti4+(0.61Å) compared to Mn4+(0.54Å).The patterns in Figure3a,b emphasize the structural compatibility that exists among Li2MnO3,Li2TiO3,and LiMn0.5Ni0.5O2phases and the ability of the Li2M′O3-type component to become inte-grated into a LiMO2-type structure.However,when TiO2(anatase)was used as a precursor,the XRD patterns of the x Li2TiO3‚(1-x)LiMn0.5Ni0.5O2products showed two-phase character,which is clearly noticeable, for x)0.3,at approximately44°2θin Figure3c, suggesting that the Li2TiO3component is not integrated to the same extent into the LiMn0.5Ni0.5O2structure as it is in the products in Figure3b.This observation clearly emphasizes that the structural integration of Li2M′O3and LiMn0.5Ni0.5O2components depends strongly on the type of reagents used and on the reaction conditions during synthesis.The X-ray diffraction patterns of x Li2ZrO3‚(1-x)-LiMn0.5Ni0.5O2products showed that the LiMn0.5Ni0.5O2 structure is less tolerant to accommodating Li2ZrO3. Although the0.03Li2ZrO3‚0.97LiMn0.5Ni0.5O2electrode had an X-ray diffraction pattern similar to those of its M′)Mn and Ti analogues(Figure3d),x Li2ZrO3‚(1-x)LiMn0.5Ni0.5O2products with higher Li2ZrO3content (x>0.03)were multiphase,as were those of x Li2MnO3‚(1-x)LiMn0.5Ni0.5O2and x Li2TiO3‚(1-x)LiMn0.5Ni0.5O2 products for x>0.3(not shown).We are unsure of the structural features of0.03Li2ZrO3‚0.97LiMn0.5Ni0.5O2, and we do not know yet if,at this low concentration, the Li2ZrO3component is embedded within the bulk of the LiMn0.5Ni0.5O2structure or whether it exists as an independent phase.We attribute the difficulty in inte-grating Li2ZrO3into a layered LiMn0.5Ni0.5O2structure to(1)the absence of lithium layering in Li2ZrO3(Figure 1c)and(2)the relatively large second-row transition-metal ion,Zr4+(ionic radius)0.72Å).33Although refinements of the XRD patterns showed that the lattice parameters of both x Li2MnO3‚(1-x)-LiMn0.5Ni0.5O2and x Li2TiO3‚(1-x)LiMn0.5Ni0.5O2sys-tems changed gradually with x,which is consistent with increasing amounts of Li2MnO3and Li2TiO3,respec-tively,X-ray diffraction is not a sufficiently sensitive technique for determining the local variations in com-position and structure in these materials.We therefore relied on more sensitive techniques,such as HRTEM and NMR,to obtain a more detailed understanding of the local structure of these materials at a nanoscopic level.HRTEM Analyses.The HRTEM images of struc-tures in the x Li2MnO3‚(1-x)LiMn0.5Ni0.5O2and x Li2-TiO3‚(1-x)LiMn0.5Ni0.5O2systems were examined at three compositions,x)0.0,0.14,and0.30.To accom-modate the diffraction peaks that occur between20°and 23°2θin the XRD patterns of these compounds,the indexing of the lattice fringes was based on the mono-(33)Shannon,R.D.;Prewitt,C.T.Acta Crystallogr.1969,B25, 925.ttice Parameters ofx Li2M′O3‚(1-x)LiMn0.5Ni0.5O2Structures(M′)Mn,Ti,Zr) x a(Å)c(Å)vol(Å3)x Li2MnO3‚(1-x)LiMn0.5Ni0.5O20.03 2.870(1)14.259(3)101.69(2)0.14 2.864(1)14.246(2)101.21(2)0.30 2.859(1)14.246(2)100.85(2)x Li2TiO3‚(1-x)LiMn0.5Ni0.5O20.03 2.876(1)14.274(2)102.22(1)0.14 2.881(1)14.304(2)102.78(2)0.30 2.887(1)14.342(2)103.52(2)x Li2ZrO3‚(1-x)LiMn0.5Ni0.5O20.03 2.873(1)14.259(2)101.94(2)xLi2M′O3‚(1-x)LiMn0.5Ni0.5O2for Li Batteries Chem.Mater.,Vol.16,No.10,20041999clinic unit cell (C 2/m )defined by Strobel et al.22for Li 2-MnO 3.The examination of the X-ray diffraction patterns of these materials shows the same type of “superlattice”peaks between 20°and 23°2θas described by Dahn et al.17for the Li[Ni x Li (1/3-2x /3)Mn (2/3-x /3)]O 2system (0<x <0.5).These peaks are of particular importance to the local (nearest cation neighbor)environment of the Li ions because the peak intensities are a rationalization of the ordering of the lithium,manganese,and other transition-metal ions present in the transition-metal layers.However,an ideal superlattice can occur only if the atomic distribution is uniform on a length scale approaching that of the superlattice length and if the ions occur in exactly the correct proportion.When such an ideal superlattice occurs,it is expected that the broadening of the superlattice peaks will correspond to the same crystallite size as that determined from “nonsuperlattice”peaks.If,however,the composition is nonideal in the sense that it deviates from that required by an ideal superlattice,or if the composition is not spatially uniform,then the superlattice peaks will be weakened and broadened.34This is of specific concern to this work because of the observations ofrelativelyFigure 3.Powder X-ray diffraction patterns of x Li 2M ′O 3‚(1-x )LiMn 0.5Ni 0.5O 2structures:(a)x Li 2MnO 3‚(1-x )LiMn 0.5Ni 0.5O 2(x )0.0,0.03,0.14,0.30);(b)x Li 2TiO 3‚(1-x )LiMn 0.5Ni 0.5O 2(x )0.03,0.14,0.30)(Ti[OCH(CH 3)2]4precursor);(c)x Li 2TiO 3‚(1-x )LiMn 0.5-Ni 0.5O 2(x )0.03,0.14,0.30)(TiO 2precursor);(d)x Li 2ZrO 3‚(1-x )LiMn 0.5Ni 0.5O 2(x )0.03).2000Chem.Mater.,Vol.16,No.10,2004Kim et al.broad superlattice lines in the nonideal compound compositions,for example,x Li 2M ′O 3‚(1-x )LiMO 2com-positions x )0.14and x )0.30.In such cases,where relatively broad superlattice peaks are observed,two different structural models can be used to rationalize the observation.We may consider either (1)that the broadening is due to isolated regions of composition close to what is ideal for superlattice formation surrounded by disordered regions of nonideal composition 35or (2)that the composition is uniform,but that the ordered structure is interrupted by anti-phase domain bound-aries.36In either case,there will be variations in the local environment of Li compared to the ideal superlat-tice.We therefore resorted to HRTEM to provide further evidence for the existence of such nonuniformities in the superlattice,as suggested by the broadening of the superlattice peaks in Figure 3a,ing this technique,it is possible to directly image the superlattice structure corresponding to the peaks at 20-23°2θ.Figure 4a shows the lattice fringes of a LiMn 0.5Ni 0.5O 2structure associated with the close-packed 001planes of the monoclinic unit cell (d ≈ 4.7Å),which are equivalent to the 003planes of the trigonal (R 3h m )unit cell of LiCoO 2.The image shows well-behaved packing of these planes,as might be expected in a stoichiometric LiMn 0.5Ni 0.5O 2structure.However,if the crystals are at orientations that allow the possibility of coherentdiffraction from the superlattice structure,then fringes of weak contrast and discontinuous morphology corre-sponding to the “superlattice”peaks at 20-23°2θ(i.e.,the (020)and (110)peaks,d ≈4.3Å)are observed within a single grain of LiMn 0.5Ni 0.5O 2(Figure 4b).The nanoscale,spatial variation in the lattice fringe contrast suggests that,even in the parent LiMn 0.5Ni 0.5O 2struc-ture,the distribution of the cations in the transition-metal layers is not random or homogeneous.We also note that the existence of the discontinuous superlattice structure for LiMn 0.5Ni 0.5O 2observed here by HRTEM is indeed consistent with the MAS NMR results,16,26but would be difficult to predict using isolated X-ray dif-fraction spectra.In an attempt to have a more quantitative measure of the structure and chemistry variations suggested in Figure 4b,two adjacent areas are circled in Figure 4b for specific consideration.Qualitatively,it is observed that the lattice fringes corresponding to the superlattice structure show greater definition in region (i)compared to region (ii).This observation is relevant because the amplitude of contrast of the lattice fringes is propor-tional to the superlattice structure factor.37Assuming that the crystallographic orientation and crystal thick-ness are approximately the same for these adjacent regions,the variation in lattice fringe contrast implies a spatial variation in the structure factor for the superlattice.A more quantitative demonstration of the(34)Guinier,A.X-ray Diffraction ;W.H.Freeman and Company:San Francisco,CA,1963;p 254.(35)Rhines,F.M.;Newkirk,J.B.Trans.ASM 1953,45,1029.(36)Jones,F.W.;Sykes,G.Proc.R.Soc.London A 1938,166,376.(37)Williams,D.B.;Carter,C.B.Transmission Electron Micros-copy ;Plenum Press:New York,1996;p441.Figure 4.(a)LiMn 0.5Ni 0.5O 2lattice fringes associated with (001)peak (C 2/m ).(b)HRTEM image of LiMn 0.5Ni 0.5O 2formed using phase contrast from superlattice peaks and local fast Fourier transforms (inset)from adjacent areas (i)and (ii)to emphasize the spatial variation in superlattice order parameter.(c)HRTEM image of 0.30Li 2MnO 3‚0.70LiMn0.5Ni 0.5O 2formed using phase contrast from superlattice peaks and fast Fourier transforms (inset)from adjacent areas (i)and (ii)to emphasize the spatial variation in superlattice order parameter.(d)Lattice fringe images in 0.14Li 2TiO 3‚0.86LiMn 0.5Ni 0.5O 2.(e)Lattice fringe images in 0.30Li 2-TiO 3‚0.70LiMn 0.5Ni 0.5O 2.xLi 2M ′O 3‚(1-x)LiMn 0.5Ni 0.5O 2for Li Batteries Chem.Mater.,Vol.16,No.10,20042001。

LiIonBattery(LIB)-WelcometoCatalysisDatabase：锂离子电池(LIB)-欢迎催化数据库

Li Ion Battery (LIB)
K.Devaki (CH09M001)
Battery
• Battery: Transducer which converts chemical energy into electrical energy and vive versa. • Chemical reactions: Oxidation and reduction • Free energy change of the processes appears as electrical energy. • Primary battery-not rechargeable • Secondary battery- rechargeable
Disadvantages • Low life cycle • Low energy density (30 ~ 40 Wh/Kg)
Recharge-ability : Basically, when the direction of electron
discharge (negative to positive) is reversed, restoring power.
Schematic of Li Ion Battery
The cathode half reaction (with charging being forwards) is: LiCoO2 ↔ Li 1-x CoO2 + xLi+ + xeThe anode half reaction is: xLi + + xe- + 6C ↔ Lix C6
Memory Effect : When a battery is repeatedly recharged

Issues and challenges facing rechargeable lithium batteries

50100150Energy density (W h kg -1)Lighter weightLead–acid200250Ni–CdNi–MHLi ionLi metal ('unsafe')PLiONS m a l l e r s i z eComparison of the different battery technologies in terms of volumetric and gravimetric energy density. The share of worldwide sales for Ni–Cd, Ni–MeH and Li-ion portable batteries is 23, 14 and 63%, respectively. The use of Pb–acid batteries is restricted mainly to SLI (starting, lighting, ignition) in automobiles or standby applications, whereas Ni–Cd batteries remain the most suitable technologies for high-power applications (for example, power tools ).Non-aqueous liquid electrolytePositive (Li x Host 1)Li +Li ++Non-aqueous liquid electrolytePositive (Li x Host 1)(Lithium)Li +Li ++After 100 cyclesa-Metal oxidesCapacity (A h kg -1)1000400600800Nitrides LiM y N 2Positive material:of Li ion of Li metal( limited cycling)Negative material:© 2001 Macmillan Magazines Ltdsolubility of the resulting thiolates in the electrolyte, leading to self discharge.Materials for negative electrodesAs a result of numerous chemical (pyrolitic processing) or physical (mechanical milling) modifications, carbon negative electrodes display electrochemical performances that are improving continu-ously. Reversible capacities of around 450 mA h g reached, compared with a practical value of 350 mA h g (372 mA h g research efforts are focused on searching for carbon alternatives in the hope of finding materials (Fig. 4) with both larger capacities and slightly more positive intercalation voltages compared to Li/Li to minimize any risks of high-surface-area Li plating at the end of fast recharge, which are associated with safety problems.resulted in the emergence of Li transition-metal nitrides as a new potential class of anode materials reversible capacity (600 mA h g NATURE |VOL 414|15 NOVEMBER 2001|363200150100500C + (mA h)C – (mA h)© 2001 Macmillan Magazines Ltd364NATURE |VOL 414|15 NOVEMBER 2001|2.73.12.93.53.334561251β -A l u mi n a10080604020.18P2S5.0.37L i2S .0.45Li l g l as s68conductivity of the electrolyte. In contrast, there are only a few Li-based salts or polymers to choose from, the most commonly used ones being based on polyethylene oxide (PEO). The results from research efforts aimed at counterbalancing this deficit have led to the present level of research and development on electrolytes.Guided by general concepts of viscosity and dielectric constants, optimizing the ionic conductivity of a liquid electrolyte almost becomes a field-trial approach with the hope of finding the key ingredients. For instance, only ethylene carbonate can provide the ad hoc protective layer on the surface of graphite that prevents further reaction (continuous electrolyte reduction and self-discharge). Ethylene carbonate is therefore present in almost all commercial compositions, thinned with other solvents owing to its high melting point. Why the homologous propylene carbonate is unsuitable for this protective layer remains an open question, reminding us that chemistry has its secrets.In contrast, achieving high ionic conductivity in Li-based polymer electrolytes requires a better understanding of the fundamentals of ion dissociation and transport.Both the nature of the polymer–salt interaction and the precise structure of highly concentrated electrolyte solutions have always resisted rationaliza-tion. Nevertheless, a principal goal has been to search for new, highly conductive salts with a large electrochemical window, which form a eutectic composition with PEO that melts at the lowest possible tem-perature57. The concept of non-coordinating anions with extensive charge delocalization was achieved with the perfluorosulphonimide Li+[CF3SO2NSO2CF3]–salt (abbreviated as LiTFSI)58. Figure 8 shows the marked improvement when passing, with simple PEO, from a ‘conventional’ LiCF3SO3 salt (curve 1) to the imide salt (curve 2), where an order of magnitude is gained, not ignoring the larger elas-tomeric domain towards low temperature.The polymer architecture has a role independently of dissociation. Attaching the side chains of the solvating group to the polymer increases the degrees of freedom as a result of dangling chain ends; this improves conductivity (Fig. 8, curve 3), but compromises the mechanical properties.Although efforts aimed at enhancing the ionic conductivity of polymer electrolytes have been insufficient to permit operation at room temperature, they have benefited liquid-based electrolyte systems in terms of cost and safety, so that battery manufacturers of Li-ion cells are eager to see the further development of organic anion-based salts able to operate at voltages greater than 4.5 V. LiTFSI is an example of this cross-fertilization. Although extremely resistant to oxidation itself, the electrochemical use of such a salt is limited to 4 V in presence of an Al collector, because a stable and soluble Al salt can be formed as a consequence of the robustness of the anion bonds. With the less stable coordination anions (LiPF6), decomposition occurs immediately and is accompanied by formation of protective AlF3. However, owing to its high conductivity in any medium, its safety and lack of toxicity, LiTFSI is being used increasingly in Li-ion batteries, the corrosion problem having being solved by simple addition of a passivating coordination-type salt. A wide range of anion-forming systems now exists, especially in the imide family, and these are viewed as candidates for high conductivity and Al passivation.Having exploited most of the possibilities offered by ‘dry’ polymers to improve conductivity (ability, amorphous state and lowest possible glass-transition temperature T g controlling the ion mobility), a remaining option was to use additives, known in polymer science as plasticizers, to act as chain lubricants, so leading to the development of ‘hybrid’ polymer electrolytes59. Indeed, suitable plasticizers are chosen between the same polar solvents as for liquid electrolytes60, such as propylene carbonate, ȍ-butyrolactone or polyethylene glycol ethers, or are formed from short-chain PEO (4–25 monomer units). A lightly plasticized material (10–25% additive) improves conductivity by an order of magnitude (Fig. 8, curve 4). Gels, on the other hand, contain 60–95% liquid electrolyte, and are only 2–5 times less conductive than their liquid counterpart61(Fig. 8, curves 7–10). Interestingly, when the gelling agent is apolyether, most of the solvation still takes place through the polymerchains rather than the carbonate solvents, the latter being less proneto donate electron pairs. Understandably, the lightly plasticizedsystems can be used in a Li-metal configuration, as much of theresilience of the pristine polymer is retained, whereas the much softergels require a Li-ion configuration.It is surprising that in spite of the direct link between their ionicconductivity and their degree of amorphicity, very little is knownabout the structural chemistry of polymer electrolytes. In contrast tothe well established dynamic view of ionic conductivity on thesematerials, Bruce et al.62recently proposed a structural view,highlighting the importance of aligning or organizing the polymerchains in order to enhance the levels of ionic conductivity. Similarly,Wright and co-workers63and Ingram64focused on the liquidcrystalline state to force the solvating polymer into a conformationthat was dictated by the liquid crystal part. The result is a partialdecoupling of the conductivity from the glass-transition tempera-ture of the polymer. The conductivity of such liquid crystalline chainpolymers is low at room temperature, but reaches liquid-like valuesat high temperature or when kept under polarization, and remains soupon cooling to room temperature (Fig. 8, curves 5 and 6), withoutappreciable activation energy64. As these new perspectives generaterenewed interest in the design of polymer electrolytes, it is hoped thatsolutions may eventually be found to the problems of ionicconductivity afflicting this class of materials at ambient orsubambient temperature.The addition of nanoparticle fillers (10% w/w), such as Al2O3orTiO2,to simple PEO compounds increases the conductivity several-fold at 60–80 ᑻC, and prevents crystallization for at least several weeksat room temperature65. Two important advantages of these systemsare an increase in the apparent Li transport number, from a low of≈0.3 (common to polymer, liquid and gels) to ≈0.6, and the forma-tion of a stable, low-resistance interface in contact with Li. Becausethese materials obey different conduction mechanisms, they arepresently the focus of many studies, both practical and theoretical66.Technologies based on either solid polymer or ‘hybrid’ polymerelectrolytes offer great advantages that will be necessary to meet theflexible, shape-effective requirements dictated by today’s electronicminiaturization, while at the same time providing a larger autonomy.Current Li technologies rely on liquid-jellyroll or prismatic-cellconfigurations. Neither fits well in a multiple-cell configuration.This is in marked contrast with the recent thin, plate-like plastic(PLiON) technology that enables excellent packing efficiency, asmultiple plates can be densely packaged in parallel within one cellwhile preserving the flexibility of the overall package. Future technol-ogy improvements should focus on better chemical engineering ofthe bonded laminates, so as to obtain even thinner cells. Similarattributes can be provided by the solid Li-polymer technology, whichin addition exhibits extra capacity and is free of electrolyte leakage.This currently operates at 80 ᑻC. Although warm temperatures maybe an advantage for the large batteries required by the transportationsector, problems of conductivity have to be solved for electronicapplications, as emphasized earlier.The electrode–electrolyte interfaceThe Li-ion cell density can be improved through a selective use ofappropriate existing or new materials for negative and positiveelectrodes. However, optimizing an electrode material is only the firststep in the process leading to its implementation in a practical cell.Indeed, while the capacity of a cell is nested in the structural orelectronic behaviour of its electrode, poor cell lifetimes are rootedmainly in side reactions occurring at the electrode–electrolyte inter-face. Thus, mastering the chemical stability of any new electrodematerial with respect to its operating liquid or polymer electrolytemedium, which requires a control of the electrode–electrolyteinterface through surface chemistry, is as important as designing newNATURE|VOL 414|15 NOVEMBER 2001|365©2001 Macmillan Magazines Ltdmaterials. Tackling interfacial issues is both tedious and complex. We should remember that, despite many years of research devoted to the mechanism by which the solid electrolyte interphase forms on Li or carbonaceous materials, its composition and nature are still the sub-ject of much controversy. In contrast, the positive electrode interface has received little attention over the years, despite its equally crucial role. Its importance is amplified with the Li-ion technology, where high voltages exceed the electrochemical resistance of the electrolyte oxidation, and even favour its catalytically driven decomposition. Thus, it is critical to control the electrode surface so as to modify its catalytic activity towards electrolyte decomposition. The strategy developed to address this issue uses coatings that encapsulate, through chemical or physical means, the electrode grains with either an inorganic or an organic phase. This concept, successfully applied to the spinel LiMn2O4, is based on minimizing the surface area of the active material in direct contact with the electrolyte33. The coating must allow easy diffusion of Li ions and, although insulating in nature, must be thin enough to allow the electrons to tunnel through. Equally relevant is the unexplained role of filler additives in polymer electrolytes65, which markedly reduce the interfacial impedance in contact with Li.Thirty years after its initial observation, the key issue of Li den-drite growth, which was thought to be governed mainly by current densities, remains highly topical, especially in light of recent promis-ing results obtained by Aurbarch’s and Bates’ groups. Revisiting Exxon’s solvent 1-3-dioxolane, Aurbarch and co-workers67showed that the use of LiAsF6salt led to a completely different Li morphology from that obtained from an ethylene carbonate–dimethyl carbonate (EC–DMC) electrolyte. They explained this in terms of the reactivity of dioxolane with lithium, which forms an elastomeric coating endowing the Li surface with plasticity and flexibility, thereby reduc-ing dendrite growth. These findings were implemented in Li/MnO2 commercial Tadiran cells that, under well defined cycling conditions, are claimed to be safe. The bulk polymerization of the cyclic ether, initiated at the positive electrode on overcharge, acts as a thermal shutdown. Even more spectacular are recent reports by Bates et al.68 who succeeded in cycling LiCoO2/Li thin-film batteries for more than 50,000 cycles using a glassy electrolyte in ≈1-Ȗm-thick films obtained by sputtering techniques. By controlling the uniform Li stripping–plating mechanism, the same authors demonstrated the feasibility of a Li-free, rechargeable, thin-film battery — that is, cells constructed in the discharged state with no Li metal initially present69. Such findings, whether resulting from low-current density or the use of solid electrolyte, show that the problem of dendrite growth can be solved, at least with special cell configurations. Visco and co-workers70 recently showed that a glassy nanometric layer deposited on Li metal completely insulates it from its environment, even in the presence of liquids, and that this coating can be applied at a high production rate. With further work devoted to the implemen-tation of these findings to large-size Li batteries, the development of a Li-free rechargeable battery remains a realistic goal for the future.The principal challenge for Li-based rechargeable batteries, or indeed for any battery, lies in gaining better understanding and control of the electrode–electrolyte interface in the hope of designing new solid–solid or solid–liquid interfaces. For example, the nature of the secondary reactions occurring at high temperature, which cause cell failure, remains an unanswered question that must be addressed to ensure the practical success of these technologies. In this case, however, the main difficulty stems from a lack of available techniques to probe the evolution of the electrode–electrolyte interface at a local level. We have so far relied (with the exception of X-ray diffraction) on post-mortem rather than in situ studies to determine how the electrodes or interfaces age with time either under cycling or storage conditions, thereby missing key information. But introduction of the plastic Li-ion-type technology has created new opportunities to perform a wide variety of in situ characterization techniques. These include X-ray absorption near-edge structure, nuclear magnetic resonance and Mössbauer spectroscopies, or even scanning electron microscopy observations that allow real-time visualization of dendrite growth at an interface71. Efforts aimed at developing new characterization tools must be vigorously pursued so as to create a comprehensive database on the electrode–electrolyte interface. ConclusionConsumers are in constant demand for thinner, lighter, space-effective and shape-flexible batteries with larger autonomy. Such demand will continue to generate much research activity towards the development of new cell configurations and new chemistries. In this review we hope to have conveyed the message that the field of energy storage is advancing faster than it perhaps has ever done in the past. The benefits, in terms of weight, size and design flexibility provided by today’s state-of-the-art Li-ion configurations, which owe much to the design engineers’ striving to develop efficient, economical microtech-nologies, are a good illustration. The Li-based battery chemistry is relatively young, and as such is a source of aspirations as well as numerous exciting challenges. The latter are not limited to solid-state chemists. The effort should be highly multidisciplinary with strong roots in the fields of organic and inorganic chemistry, physics, surface science and corrosion. Through materials design we can expect significant improvements in energy density. And although designing new materials can be intuitive or based on chemical concepts, coupling these efforts with those of theorists who are able to perform band-structure calculations on envisioned compounds will prove to be highly beneficial. Of equal importance is a better understanding of the electrode–electrolyte interface to facilitate design of new interfaces. Here the goal is well defined, although we must diverge from the empirical approach used so far, and make full use of the recent progress achieved by in situ characterization. As Li-recharge-able batteries enter their teenage years, scientists and engineers predict an even brighter future lies ahead.1.Takeshita, H. Portable Li-ion,worldwide. Proc. Conf. Power 2000, San Diego, 25 September 2000.2Ikeda, H., Saito, T. & Tamura, H. in Proc. Manganese Dioxide Symp.Vol. 1 (eds Kozawa, A. & Brodd, R.H.) (IC sample Office, Cleveland, OH, 1975).3.Steele, B. C. H. in Fast Ion Transport in Solids(ed. Van Gool, W.) 103–109 (North-HollandAmsterdam, 1973).4.Armand, M. B. in Fast Ion Transport in Solids(ed. Van Gool, W.) 665–673 (North-HollandAmsterdam, 1973).5.Rouxel, J., Danot, M., & Bichon, M. Les composites intercalaires Na x TiS2. Etude gènérale des phasesNa x TiS2et K x TiS2. Bull. Soc. Chim. 11,3930–3936 (1971).6.Di Salvo, F. J., Schwall, R., Geballe, T. H., Gamble, F. R. & Osieki, J. H. Precursor effects ofsuperconductivity up to 35ᑻK in layered compounds. Phys. Rev. Lett.27,310–313 (1971).7.Whittingham, M. S. Electrochemical energy storage and intercalation chemistry. Science192,1226(1976).8.Whittingham, M. S. Chalcogenide battery. US Patent 4009052.9.Rao, B. M. L., Francis, R. W. & Christopher, H. A. Lithium-aluminium electrodes. J. Electrochem. Soc.124,1490–1492 (1977).10.Broahead, J. & Butherus, A.D. Rechargeable non-aqueous battery. 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A.) 212 (Elsevier,London, 1989).60.Feuillade, G. & Perche, P. Ion conductive macromodular gels and membranes for solid lithium cells. J.Appl. Electrochem.5,63–69 (1975).61.Stallworth, P. E. et al.NMR, DSC and high pressure electrical conductivity studies of liquid andhybrid electrolytes. J. Power Sources81,739–747 (1999).62.MacGlashan, G. S., Andreev, Y. G. & Bruce, P. Structure of the polymer electrolyte poly(ethyleneoxide): LiAsF6. Nature398,792–793 (1999).63.Zheng, Y., Chia, F., Ungar, G. & Wright, P. V. Self-tracking in solvent-free, low-dimensional polymerelectrolyte blends with lithium salts giving high ambient DC conductivity. Chem. Commun.16, 1459–1460 (2000).64.Imrie, C. T., Ingram, M. D. & McHattie, G. S. Ion transport in glassy polymer electrolytes. J. Phys.Chem. B103,4132–4138 (1999).65.Croce, F., Appetecchi, G. B., Persi, L. & Scrosati, B. Nanocomposite polymer electrolytes for lithiumbatteries. Nature394,456–458 (1998).66.Sata, N., Eberman, K., Eberl, K. & Maier, J. Mesoscopic fast ion conduction in nanometre-scale planarheterostructures. Nature408,946–949 (2000).67.Aurbach, D. Review of selected electrode solution interactions which determine the performance ofLi and Li-ion batteries. J. Power Sources89,206–218 (2000).68.Bates, J. B., Dudney, N. J., Neudecker, B., Ueda, A. & Evans, C. D. Thin film lithium and lithium-ionbatteries. Solid State Ionics135,33–45 (2000).69.Neudecker, B. J., Dudney, N. J. & Bates J. B Lithium-free thin film battery with in situ plated Li anode.J. Electrochem. Soc.147,517–523 (2000).70.Visco, S. J. The development of reversible lithium metal electrodes for advances Li/S batteries. Conf.Proc. Int. Meeting on Power Sources for Consumer and Industrial Applications, Hawaii, 3-6September 2001.71.Orsini, F. et al.In situ SEM study of the interfaces in plastic lithium cells. J. Power Sources81–82,918–921 (1999).AcknowledgementsThe authors thank their colleagues, both in academic institutions and industry, for sharing the gratifying dedication to this field of progress, and P. Rickman for help drawing the figures.NATURE|VOL 414|15 NOVEMBER 2001|367©2001 Macmillan Magazines Ltd。

2019湖北自然科学奖拟提名项目公示

2019年度湖北省自然科学奖拟提名项目公示1.项目名称：电池材料中离子/电子快速输运构筑及存储机理研究2. 提名单位及提名意见提名单位：武汉理工大学提名意见：电化学储能技术是新能源发展的热点，锂/钠离子电池储能因其独特的性能已成为优先发展方向之一，其关键在于锂/钠离子电池电极材料。

然而材料的离子/电子快速输运困难、效率低下、产品安全是制约储能技术发展的重大瓶颈，构建快速离子/电子传输路径、提高效率、揭示材料存储机理和离子输运规律是攻克这一瓶颈的关键科学问题。

该项目在国家863计划、自然科学基金委的资助下，针对锂/钠电极材料中离子/电子快速输运、材料存储机理和离子输运规律，开展了系统的基础性研究。

发现了三维快速离子通道的新体系通过碳包覆构筑了电子快速传输的路径，离子/电子导电性能差的材料通过纳米化缩短离子传输距离、多种碳复合构筑电子传输通道，提高了功率密度，同时减少了副反应，提高了循环稳定性和安全性；发现了新型电解液体系，充放电过程中在电极表面形成稳定的SEI膜，减少副反应，首次将NaFSI基电解液用于钠离子电池，实现了电极材料的高效稳定循环；通过球差矫正透射电子显微技术，在原子尺度直接观察了锂/钠离子脱嵌过程中的离子占位变化，揭示了材料的存储机理和输运机制。

为新型高性能储能电池材料的研制与开发提供理论指导。

该项目发表的8篇代表性论文具有较大的国际影响力，4篇入选ESI前1%高被引论文，SCI他引1734次，单篇最高SCI他引506次，锂电先驱、国内外院士等许多国际著名学者正面引用。

同意提名该项目为湖北省自然科学一等奖。

3. 项目简介该项目属于材料科学领域。

电化学储能是发展新能源汽车、提高电网对间歇性可再生能源发电接纳能力的关键技术，锂/钠离子电池储能因其独特的性能已成为优先发展方向之一。

然而离子/电子快速输运困难、产品安全是制约储能技术发展的重大瓶颈，构建快速离子/电子传输路径、揭示材料存储机理和离子输运规律是攻克这一技术瓶颈的关键科学难题。

锂离子电池英文

Electrolyte
• The selection of the electrolyte depends on the potential window. • Potentials are beyond water potential and hence non aqueous electrolyte. • It should not reactive with Li+ ions, since it will decide the transport of the Li+ ions for the intercalation reaction. • The electrolyte solution commonly comprises a lithium salt dissolved in a mixture of organic solvents. • Examples include LiPF6 or LiBOB (the BOB is the anion with the boron coordinated by two oxalate groups) in an ethylene carbonate/dimethyl carbonate solvent.
Li Ion Battery (LIB)
Battery
• Battery: Transducer which converts chemical energy into electrical energy. • Chemical reactions: Oxidation and reduction • Free energy change of the processes appears as electrical energy. • Primary battery-not rechargeable • Secondary battery- rechargeable

锂离子电池层状结构三元正极材料的研究进展

锂离子电池层状结构三元正极材料的研究进展（中山大学化学与化学工程学院广州510275）摘要为改进锂离子电池的性能，化学家们一直致力于电极材料的研究。

其中，正极材料的研究更是重中之重，各种正极材料层出不穷，而层状结构三元正极材料LiNi x Co y Mn1-x-y O2因为具有较高的可逆容量、循环性能好、结构稳定性、热稳定性和相对较低的成本等优点，近年来成为研究热点。

本文主要简介其结构特点与电化学特性，并综述其制备方法的改良和改性手段，并分析该材料目前存在的问题和对其未来发展做一个设想。

关键词锂离子电池层状结构LiNi x Co y Mn1-x-y O2 研究进展Research progress in layered structural ternary cathode materials forlithium ion batteriesAbstract To improve the properties of Li-ion Battery, the chemist have been working for suitable electrode materials. Among them, the study of cathode materials is a top priority. There are a variety of cathode material. And in recent years, Layered Structural LiNi x Co y Mn1-x-y O2 as a cathode has been a hot topic, because it has a lot of advantages, such as, it has a high reversible capacity, good cycle performance, structural stability, thermal stability and relatively low cost, etc. This paper is about the introduction of its structural features and electrochemical characteristics, as well as a review of the improvement and modification means of their preparation. Finally, there are analysis of the existing problems of the materials and a vision of its future development.Key words lithium ion batteries; layered structure; LiNi x Co y Mn1-x-y O2; research progress1.引言锂离子电池的具有工作电压高、能量密度高、自放电效率低、循环寿命长、无记忆效应和环保等优点，因此广泛应用于生产生活中。

锂电池高温存放后的电化学容量衰减

Electrochemical Investigations on Capacity Fading of Advanced Lithium-Ion Batteries after Storing at Elevated TemperatureMao-Sung Wu,*,z Pin-Chi Julia Chiang,and Jung-Cheng LinIndustrial Technology Research Institute,Materials Research Laboratories,Hsinchu 310,TaiwanCapacity fading of advanced lithium-ion batteries after elevated temperature storage was investigated by three-electrode measure-ments.Capacity fading of a battery increases by increasing the state-of-charge ͑SOC ͒during storage,especially at elevated temperatures.The reversible capacity of a battery ͑SOC =100%͒at 60°C decreases from 820to 650mAh ͑79.3%capacity retention ͒after 60days.At room temperature,a battery SOC inﬂuences the capacity fading only slightly;after 65days of storage,the reversible capacity decreases from 820to 805mAh ͑98.2%capacity retention ͒.Individual effects by the anode,cathode,and electrolyte on capacity fading are analyzed with three-electrode electrochemical ac impedance.The major contribution,from X-ray photoelectron spectroscopy ͑XPS ͒and energy-dispersive spectroscopy results,comes from cathode degradation as a result of cobalt dissolution at the LiCoO 2surface layer.A minor contribution comes from the continuous reactions between lithiated mesocarbon microbead ͑MCMB ͒electrode and electrolyte components,which in turn thicken the SEI ﬁlm and consume available lithium ions.From X-ray diffraction and XPS results,high-temperature storage inﬂuences only the surface properties of MCMB and LiCoO 2electrodes;bulk properties remain unchanged.©2005The Electrochemical Society.͓DOI:10.1149/1.1896325͔All rights reserved.Manuscript submitted August 17,2004;revised manuscript received December 15,2004.Available electronically April 21,2005.In recent years,a new type of lithium-ion battery,the advanced lithium-ion battery ͑ALB,with laminated aluminum foil exterior ͒,has emerged because of its high energy density,long cycle life,and low self-discharge properties.ALB offers similar energy character-istics as the traditional lithium-ion battery but with a higher ﬂexibil-ity on the wide variety of sizes and shapes in design.1,2In practical application,batteries are operated and stored at vari-ous conditions ͑temperature and humidity ͒.Temperature is a crucial factor in the performance of lithium-ion batteries.Detriments may result from high temperature because it signiﬁcantly affects capacity fading.3-5Amatucci et al.3report that LiMn 2O 4-based lithium-ion rechargeable batteries suffer from poor storage and cycling perfor-mance at elevated temperatures.A LiMn 1.7Al 0.3O 4-hard carbon bat-tery is deteriorated because of anode ﬁlm formation between 50and 75°C.The ﬁlm is generated from the decomposed products of LiPF 6,polymerized ethylene carbonate ͑EC ͒,and Mn ions dissoci-ated from the positive active materials.4Wang et al.5propose a mechanism for irreversible capacity loss of lithium-ion spinel cells ͑coin cell ͒in high-temperature storage.Loss of cyclable lithium ions to the carbonaceous anode because of cathode acid generation is the reason.Another effect of the acid is that spinels form from Mn dissolution,but the formation cannot be accounted for capacity loss,nor does it cause degradation of the SEI layer on the carbonaceous anode.Capacity fading of the commercially available LiCoO 2-based lithium-ion batteries cycled at room temperature has been investi-gated by means of electrochemical impedance spectroscopy.Results show that cycled positive electrode contributes more to the fading because of continuous electrolyte oxidation.6Capacity fading of Sony 18650cells cycled at elevated temperatures has been investi-gated by Ramadass et al.,7concluding that the fading was due to a repeated ﬁlm formation and dissolution over the surface of anode.This repetition increases the rate of lithium loss and increases the anode resistance.In both cases,6,7the external metallic cans are opened for electrode retrieval,and new half-cells are made in glove boxes ﬁlled with ultrapure argon to test for the electrodes’separate properties.Reassembly is inconvenient and may cause damage to the electrodes.As mentioned earlier,capacity fading of lithium-ion batteries may result from the anode,the cathode,and the electrolyte.It is difﬁcult to analyze the phenomena with a two-electrode system.If areference electrode may be added,then more mechanisms may be studied and phenomena understood.Therefore,this paper is to in-vestigate the capacity fading of commercial ALB after high-temperature storage using a three-electrode system.Three-electrode electrochemical impedance is used to analyze the individual effects by the anode,the cathode,and the electrolyte.Structural changes in the electrode materials after storage are also studied.ExperimentalComposition of the lithium-cobalt-oxide electrode was 90wt %LiCoO 2͑10␮m diam,Nippon Chemical ͒,7wt %KS6͑Timcal SA ͒,and 3wt %polyvinylidene ﬂuoride ͑PVDF,Kuraha Chemical ͒binder.Powder was mixed in a solvent of N -methyl-2-pyrrolidone ͑NMP,Mitsubishi Chemical ͒to form slurry.The slurry was coated onto aluminum foil ͑20␮m in thickness ͒and dried at 140°C.The electrode ͑200␮m in thickness ͒was then pressed to a resultant thickness of 150␮m.The mesocarbon microbead ͑MCMB ͒elec-trode,composed of 92wt %MCMB ͑Osaka Gas,25␮m diam ͒with 8wt %PVDF binder and NMP,was subjected to the same processing steps as the lithium-cobalt-oxide electrode,except that it was coated onto copper foil ͑15␮m thick ͒.Resultant thickness of the MCMB electrode was 135␮m ͑before pressing the thickness was 180␮m ͒.Batteries were assembled in a dry room.The manufacturing pro-cess was as follows:Both electrodes were dried at 120°C for 3h in vacuum and then cut into appropriate sizes for winding with sepa-rator ͑Celgard 2320,20␮m in thickness ͒.The roll of electrodes and separator was inserted into an aluminum-plastic laminated ﬁlm case.3.2g of electrolyte was injected and then the case was sealed off at a reduced pressure.Electrolyte was 1M lithium hexaﬂuorophos-phate ͑LiPF 6,Tomiyama Pure Chemical ͒in a mixture of 25%EC ͑Merck ͒,25%propylene carbonate ͑PC,Merck ͒,and 50%diethyl-ene carbonate ͑DEC,Merck ͒by volume.Water content of the elec-trolyte measured via Carl Fischer titration in an argon-ﬁlled glove box was less than 10ppm.The fresh battery had external dimen-sions of 3.8ϫ35ϫ70mm.The capacity was about 820mAh and weighed 17.5g.To monitor changes in voltage and impedance of the anode or cathode,a reference electrode was placed in the center of the battery between the two electrodes.A lithium chip was pressed onto one end of a ﬁne copper wire to make the reference electrode.Before stor-age,the three-electrode batteries were cycled between 4.2and 2.75V for three times with a charge/discharge unit ͑Maccor model series 4000͒.The procedure consisted of constant current at 82mA followed by constant voltage at 4.2V until the current tapered down*Electrochemical Society Active Member.zE-mail:ms គwu@Journal of The Electrochemical Society,152͑6͒A1041-A1046͑2005͒0013-4651/2005/152͑6͒/A1041/6/$7.00©The Electrochemical Society,Inc.A1041to 20mA.Discharge current was 82mA.The batteries were charged to different SOCs ͑40,70,and 100%͒and stored open-circuited at room temperature and at 60°C for 1-65days.During storage,in order to determine the reversible capacity,batteries were charged/discharged occasionally for two cycles at 82mA ͑about 0.1C ͒at room temperature.Then the batteries were charged again to the desired SOC ͑s ͒and the storage process continued.Maccor facilitated simultaneous and independent recordings of the total cell voltage and the half-cell voltage for both positive and negative electrodes vs.the reference electrode.Three-electrode im-pedance measurements were taken by means of a potentiostat/galvanostat ͑Schlumberger SI 1286͒and a frequency response ana-lyzer ͑Schlumberger SI 1255͒.Scanning frequencies ranged from 50kHz to 0.01Hz,perturbation amplitude 10mV.Scanning electron microscopy ͑SEM ͒and energy-dispersive spectrometry ͑EDS ͒were done with a ﬁeld emission SEM ͑FE-SEM,LEO-1530at an accelerating voltage of 15keV and coupled with an EDS ͑LEO-1550͒.Crystal structures of the MCMB and LiCoO 2were identiﬁed by X-ray diffraction ͑XRD,XD-5͒with a Cu K ␣target ͑wavelength 1.54056Å͒.Diffraction data were col-lected for 1s at each 0.04°step width over 2␪,ranging from 10to 90°.Surface properties of the cathode after storage were conﬁrmed by X-ray photoelectron spectroscopy ͑XPS;Perkin Elmer,PHI Quantera SXM ͒with a focused monochromatic Al K ␣radiation ͑1486.6eV ͒.Before any experiment,batteries were fully charged,disassembled in a glove box,washed with DEC,and dried in vacuum at 100°C for 5h.Sample powders of anode and cathode were scraped off the electrodes’current collector.Results and DiscussionCapacity variations of ALB during storage .—Figure 1shows the capacity variation of ALBs during storage at different SOCs and temperatures.Capacity decay at room-temperature storage ͑Fig.1a ͒is negligible as compared with 60°C ͑Fig.1b ͒.At room temperature,a battery’s SOC inﬂuences the fading trend slightly.The original capacity of a battery SOC =100%is 820mAh;after 65days of room-temperature storage it decreased to 805mAh ͑98.2%capacity retention ͒.Lowering a battery’s SOC hinders its capacity decay,as one can see from Fig.1a that the capacity remains unchanged for a battery SOC =40%.Therefore,in addition to the storage tempera-ture,a battery’s SOC is a factor in capacity fading.Batteries stored at 60°C show a steeper capacity fading trend,and the decrease is most signiﬁcant in the ﬁrst few days ͑Fig.1b ͒.The fading depends strongly on a battery’s SOC;the higher the SOC,the more the fading.Capacity of a fully charged battery ͑SOC =100%͒decreases from 820to 650mAh after 60days at60°C ͑79.3%capacity retention ͒.The fade of a battery SOC =40%is relatively less,from 820to 750mAh ͑91.5%capacity re-tention ͒.The inﬂuence of SOC on capacity fading becomes more pronounced with elevating the storage temperature.Three-electrode electrochemical impedance analysis .—Early re-searchers believed that impedance of a battery is contributed by different factors,such as the electrolyte,the passivation ﬁlm,charge transfer,lithium-ion diffusion in electrodes,etc.8-13Three-electrode electrochemical impedance spectroscopy ͑EIS ͒has been developed to analyze the individual effects of each component on capacity paring the Nyquist plot with an equivalent circuit model identiﬁes the sources of impedance.Figure 2a shows both the measured and the simulated impedance spectra of a full battery before and after 15days of storage at 60°C ͑batteries are fully charged,SOC =100%͒.With respect to the ref-erence lithium electrode,impedance of the anode ͑Fig.2b ͒,and cathode ͑Fig.2c ͒are also shown in the ﬁgure.Before any measure-ment,the spectra of individual anode and cathode are summed up to check the method validity by ensuring that the resultant combined spectra are to be equal to the full battery spectra ͑Fig.2a ͒.The nearly overlapping curves prove its applicability and reliability.Cor-responding equivalent circuits of the anode and cathode are pre-sented in Fig.3.R e resembles the ohmic electrolyte resistance.R 1,Figure 1.Capacity variations of ALB with different SOCs during storage at ͑a ͒room temperature and ͑b ͒60°C.Figure 2.Measured and simulated impedance spectra of the ͑a ͒full battery,͑b ͒anode,and ͑c ͒cathode before and after 15days storage at 60°C.͑Bat-teries are fully charged,SOC =100%.͒Figure 3.The corresponding equivalent circuit used for the analysis of the impedance spectra of ͑a ͒anode and ͑b ͒cathode.11A1042Journal of The Electrochemical Society ,152͑6͒A1041-A1046͑2005͒R 2,and R 3are the different-layer SEI-ﬁlm resistances.C 1,C 2,and C 3are the corresponding capacitance to R 1,R 2,and R 3.R CT is the charge-transfer resistance and C DL is the double-layer capacitance.W is the Warburg impedance.The semicircle in the high-frequency range,corresponding to the surface ﬁlm resistance,is composed of smaller semicircles.Each contributes resistance and capacitance from different layers of the SEI.In the low-frequency range,the semicircle resembles the charge-transfer resistance,and the linear section resembles the solid-state lithium-ion diffusion.11In general,the presence of such a linear portion implies that diffusion of lithium ions is the semi-inﬁnite diffusion condition.Semi-inﬁnite diffusion in host materials is slower than in the electrolyte solution;therefore,the linear portion is assumed to be the semi-inﬁnite diffusion in solid materials.Literature has shown many different corresponding cir-cuits to simulate the anode and cathode precisely.10-13In order to have a higher precision in modeling,different circuits have been simulated,shown in Fig.3.There are three R-C combinations in parallel to resemble anode SEI,but only one in the cathode circuit.Simulation results are identical to the experimental measurements.From Fig.2,changes in the cathode spectra are quite different from that of the anode after storage.The two electrodes have different resistance and surface chemistry,and therefore are affected differ-ently by high temperature.Individual contribution from each of the electrolyte resistance,ﬁlm resistance,and charge-transfer resistance in anode and cathode are presented,respectively,in Fig.4.During high-temperature stor-age,changes in the anode resistance are smaller than that of the cathode.In the anode ͑Fig.4a ͒,resistance changes are larger in the surface ﬁlm ͑the sum of R1,R2,and R3͒and charge-transfer than in the electrolyte.Both resistances are increased with time,especially the ﬁlm resistance.Electrolyte is believed to decompose partially and continuously on the MCMB surface to thicken the SEI ﬁlm,and this process is accelerated above room temperature.14As the SEI ﬁlm thickens,lithium-ion migration in the ﬁlm may be delayed and results in increased ﬁlm resistance.The thickened ﬁlm covers the active sites on the MCMB surface and blocks lithium ions from intercalating/deintercalating into the layer structure and charge-transfer resistance increases.Resistance of the electrolyte does not change because the decomposition amount is small ͑compared with the total electrolyte amount in a test battery ͒and therefore has little effect.In the cathode ͑Fig.4b ͒,resistance change patterns are similar;electrolyte resistance remains unchanged,and both the resistance of surface ﬁlm and charge transfer are increased.The increase in ﬁlm resistance comes from the formation of SEI on the surface of theLiCoO 2electrode.However,unlike the anode,the major contribu-tion to cathode impedance is the charge-transfer resistance,which increases most signiﬁcantly with storage at 60°C.Charge-transfer resistance generally depends strongly on the surface properties of electrode materials.Therefore,a possible source for the increased charge-transfer resistance is the structural collapse of the LiCoO 2electrode surface during high-temperature storage.The deteriorated surface may block lithium ions from intercalating/deintercalating into the layer structure and increases the charge-transfer resistance.When the anode and cathode componential resistances are com-pared,the resistance that is most high-temperature-storage affected and controlled is the charge-transfer resistance of the cathode.With the three-electrode system,individual resistances of the ALB com-ponents may be studied separately,so improvements on the electro-chemical performance of each and even of the whole battery are possible.As the cathode resistance contributes primarily to the total cell resistance after high-temperature storage,it is also interesting to investigate the changes in the diffusion resistance of cathode after storage.In general,at low frequencies,the electrochemical interca-lation process is controlled by the semi-inﬁnite diffusion.Ideally Z Јvs.Z Љis a 45°straight line ͑Warburg region ͒.15-17The slope of the straight line in the Warburg region yields the Warburg prefactor ͑␴͒.Apparent diffusion coefﬁcients of lithium intercalation can be cal-culated according to the following equation 16,17␴=RTn 2F 2A ͱ2ͩ1C LiD Li0.5͓ͪ1͔where C Li is the concentration of Li ion incorporated inside a com-posite electrode,D Li is the apparent diffusion coefﬁcient,A the geo-metrical area of the composite electrode,n the number of electrons transferred,F is Faraday’s constant,R the ideal gas constant,T absolute temperature,and ␻is the angular frequency.The real part of the complex impedance ͑Z Ј͒obtained from the cathode before and after 15days storage at 60°C plotted vs.␻−1/2is shown in Fig.5͑batteries are fully charged,SOC =100%͒.When comparing these two plots,lithium-ion concentration is assumed to be the same be-cause the electrodes are charged to the same state;composition of the materials,geometrical area,and the density are the same too.The change in Warburg prefactor value is only attributed to the diffusion coefﬁcient.Warburg prefactors for the fresh and high-temperature-aged cathode obtained from the slopes are 0.00072⍀s −1/2and 0.0011⍀s −1/2,respectively.A small Warburg prefactor value may lead to high utilization of the electrodeunderFigure 4.Individual contribution from each of the electrolyte resistance,ﬁlm resistance,and charge-transfer resistance in the ͑a ͒anode and ͑b ͒cath-ode.͑Batteries are fully charged,SOC =100%.͒Figure 5.Real part of the complex impedance ͑Z Ј͒obtained from the cath-ode ͑a ͒before and ͑b ͒after 15days of storage at 60°C plotted vs.␻−1/2.͑Batteries are fully charged,SOC =100%.͒A1043Journal of The Electrochemical Society ,152͑6͒A1041-A1046͑2005͒high-rate discharge conditions ͑diffusion control ͒.After high-temperature storage,the cathode surface layer structure has changed and the destructed layers may reduce the amount of diffusion path-ways,decreasing the utilization of active materials.In order to show the signiﬁcant changes in charge-transfer resis-tance,impedance spectra of the cathode are measured while passing through with a C/10current.Generally,discharging a battery during impedance scanning,the battery’s charge-transfer resistance de-creases with increasing the current value,because of the increasing driving force in the electrode kinetics.Figure 6shows the EIS after different storage periods at 60°C by passing a C/10current through.The EIS spectrum of a battery after 3day storage ͑Fig.6a ͒shows signiﬁcant changes.Charge-transfer resistance changes from 0.03to 0.0245⍀with the addition of C/10current.The decrease in charge-transfer resistance shows that an imposing current acceler-ates electrochemical reactions on the electrode surface.However,the battery after 30day storage at 60°C ͑Fig.6b ͒shows no signiﬁ-cant changes.Charge-transfer resistance changes only from 0.1365to 0.1339⍀with the addition of C/10current.When a bat-tery with an original higher charge-transfer resistance is applied with a C/10current density,electrochemical reactions are not enhanced signiﬁcantly.Electrochemical reactions are improved by increasing the driving force,i.e.,increased the implied currents,so to decrease the charge-transfer resistance.It may be concluded that the charge-transfer resistance of a LiCoO 2electrode after high-temperature storage for a period of time has increased signiﬁcantly.Surface properties and bulk structures of the MCMB and LiCo O 2electrodes .—It has been reported that LiCoO 2is unstable at an open-circuit potential ͑OCP ͒higher than 4.2V vs.Li/Li +due to a possibility of cobalt dissolution from LiCoO 2.18Dissolved cobalt ions therefore should be deposited onto MCMB surfaces in fully charged batteries because the reduction potential of cobalt is much higher than the potential for lithium ions to intercalate into MCMB during charging.19EDS,XPS,and XRD are used to observe the changes in surface properties and bulk structure of the electrodes after storage.Figure 7shows the EDS patterns of fully charged MCMB electrodes during storage at different temperatures after 25days.The battery stored at 60°C shows a cobalt peak,indicating the presence of cobalt on the MCMB electrode surface.After high-temperature storage,the LiCoO 2surface structure has deteriorated,and cobalt is dissociated and deposited onto the MCMB surface during charging.Cathode surface deterioration is responsible for the increased charge-transfer resistance ͑Fig.4b ͒.High storage tempera-ture and high SOC of a battery may be the reasons for the acceler-ated dissolution rate of cobalt.The XPS technique was chosen to observe the changes on sur-face properties after high-temperature storage.The Co 2p XPS spec-tra of the fully charged cathode electrodes at different storing tem-peratures after 25days is shown in Fig.8.There are two main peaks of binding energies,corresponding to Co 2p 1/2͑around 795eV ͒and Co 2p 3/2͑around 780eV ͒.20The two XPS spectra have a similar shape except a shift in their binding energy.The binding energy of cathode after high-temperature storage shifts to a higher value and has a shoulder on the high-energy side of the Co 2p 1/2component.This difference indicates that the oxidation state of cobalt in the cathode after storage is higher than that of a fresh cathode.Previous publications show that when the amount of Co 4+ions increases in a redox system of lithium-cobalt-oxide ͑Co 4+/Co 3+͒,the XPS peaks of Co shift toward the high-energy side.20Accordingly,as both the lithium and oxygen contents are kept constant in the lithium-cobalt-oxide electrode,an increase of the cobalt oxidation state increases the amount of cobalt dissolution,suitably explaining the XPSdata.Figure 6.EIS of cathode after ͑a ͒3-day and ͑b ͒30-day storage at 60°C.EIS is measured while passing a current ͑C/10͒through.͑Batteries are fully charged,SOC =100%.͒Figure 7.EDS pattern of MCMB electrodes at fully charged states after 25days of storage at ͑a ͒room temperature,and ͑b ͒60°C.Figure 8.Co 2p XPS peaks of cathodes with different storage temperature at fully charged state after 25days of storage at ͑a ͒room temperature,and ͑b ͒60°C.A1044Journal of The Electrochemical Society ,152͑6͒A1041-A1046͑2005͒According to Amatucci et al.,18LiCoO 2starts structural deterio-ration when the voltage charged is higher than 4.20V ͑vs.Li/Li +͒.Lithium ions may not be able to intercalate/deintercalate into the cathode,leading to a decrease in battery capacity.The OCP of com-mercial lithium-ion batteries in fully charged state ͑SOC =100%͒is around 4.2V,referring to an OCP of LiCoO 2electrode higher than 4.2V vs.Li/Li +.Therefore,in practical usages,capacity fading of high-temperature storage lithium-ion batteries is inevitable.Figure 9shows the XRD pattern of MCMB and LiCoO 2elec-trodes ͑SOC =100%͒after storing at different temperatures.Changes in the patterns of MCMB are small ͑Fig.9a ͒,referring to a very little changed MCMB in its bulk structure.The only noticeable change occurs on the surface ﬁlm,according to the ac impedance data ͑Fig.4a ͒.Similar results are found in the LiCoO 2electrode:high-temperature storage has little effect on the bulk structure,with the only difference being in its surface chemistry.From XPS data,cobalt dissolution occurs at the cathode.From EDS,these dissoci-ated cobalt ions are deposited on the anode.Cobalt dissolution af-fects the surface structure of lithium-cobalt-oxide electrodes,and affects the surface ﬁlm of MCMB electrodes.A generality may be concluded that capacity fading of ALBs after high-temperature stor-age is not caused by structural changes of the materials but the surface phenomena on both the MCMB and LiCoO 2electrodes.OCP and charge/discharge curve of ALB after high temperature storage .—During storage,batteries are charged/discharged occa-sionally for two cycles at 82mA ͑about 0.1C ͒to determine their reversible capacity.After two cycles,batteries are charged to their original SOCs and the storage process continues.During the two-cycle capacity-determining step,batteries are charged at room tem-perature and OCP measured.Figure 10shows the OCP variations of the MCMB and LiCoO 2electrodes after 60°C storage at their fully charged states ͑SOC =100%͒.OCP of the MCMB electrode in-creases with the storage time,from 0.01to 0.06V after 25days.A fully charged MCMB electrode self-discharges and loses lithium ions during storage,leading to an increase in its OCP ͑OCP is de-pendent on the lithium-ion concentration in a particular electrode ͒.The loss is mainly attributed by the SEI.As partial dissolution and decomposition of the SEI ﬁlm possibly thins itself,slowly becoming more porous and less protective,the ﬁlm becomes incapable of pre-venting electrons from tunneling through anymore.14Intercalated lithium ions may continuously diffuse out from the interior of the MCMB electrode through the damaged SEI to react with the elec-trolyte;consequently,a decrease in the lithium-ion concentration in the MCMB electrode ͑higher OCP ͒has resulted.In LiCoO 2,OCP increases from 4.20to 4.25V after 25days of storage at 60°C.Increase in OCP indicates a decreased lithium-ionconcentration in LiCoO 2electrode,and the decrease results from the consumption of lithiated lithium ions with electrolyte.In both elec-trodes,lithium-ion concentration decreases because reversible lithium ions from LiCoO 2are decreased after high-temperature stor-age;in which anode SEI and surface structural deterioration of LiCoO 2͑cobalt dissolution ͒are the two major sources.Generally,high temperature accelerates the continuous decomposition-formation process of the SEI and accelerates the surface structure deterioration.When a battery is fully charged ͑SOC =100%͒,MCMB has high reactivity with the electrolyte,and LiCoO 2has a low structural stability which favors the cobalt dissolution.Figure 11shows the charge/discharge curves of a MCMB elec-trode before and after 25days of storage.Below 0.2V ͑vs.Li/Li +͒,there are three signiﬁcant oxidation-reduction plateaus ͑marked 1,2,3͒,each representing the formation and decomposition of lithiated carbons.According to previous studies on lithiation of carbon ﬁber and graphite,these oxidation-reduction plateaus correspond to the potentials of two-phase coexistence.21-23Charge/discharge curves before and after storage are almost identical,different only in poten-tial plateau 3.Due to a shortage in the reversible lithium ions,con-sumed by anode SEI and cobalt dissolution from cathode,the bat-tery after storage can never be fully charged back to its original capacity,and the difference in plateau yer structuresofFigure 9.XRD patterns of ͑a ͒MCMB and ͑b ͒LiCoO 2electrodes after storage at room temperature and 60°C ͑SOC =100%͒.Figure 10.OCP variations of MCMB and LiCoO 2electrodes after 60°C storage at fully charged state ͑SOC =100%͒.Figure 11.Charge/discharge curves of the MCMB electrode before and after 25days of storage.A1045Journal of The Electrochemical Society ,152͑6͒A1041-A1046͑2005͒MCMB remain unchanged after storage͑Fig.11͒,corresponding to the XRD pattern͑Fig.9͒.The only difference is in the surface char-acteristics.ConclusionsCapacity fading of an ALB after storage depends on the battery’s SOC and its storage temperature.The relationships are directly pro-portional.Capacity of the battery SOC=100%decreases from 820to650mAh after60days of storage at60°C.Three-electrode electrochemical ac impedance technique is used to analyze the indi-vidual effects by the anode,cathode,and electrolyte on capacity fading.After storage,changes in the anode resistance are smaller than that of the cathode.In anode,changes in electrolyte resistance are small.Both theﬁlm and charge-transfer resistance increase slightly with storage time.But a different resistance result has been obtained for the cathode.After high-temperature storage,the surface layer structure has changed.The binding energy of cathode after high-temperature storage shifts to a higher value and has a shoulder on the high-energy side of the Co2p1/2component,indicating the cobalt dissolution.The destructed cathode layers therefore reduce the amount of diffusion pathways for lithium ions and decrease the utilization of the active material.A major contribution to capacity fading is the cathode degradation due to cobalt dissolution from the surface layer.Lithium-ion concentration decrease in both the MCMB and LiCoO2electrodes after storage suggests less reversible lithium ions,mainly due to the continual SEI formation/ decomposition on MCMB electrode,and to the surface structural deterioration of LiCoO2electrode͑cobalt dissolution͒.From the XRD results,high-temperature storage affects only the surface prop-erties of electrodes,the original bulk properties remain unchanged. Charge/discharge curves of MCMB electrodes demonstrate a short-age of reversible lithium ions,and most importantly,an undamaged internal structure.AcknowledgmentsThis work was supported by the Ministry of Economic Affairs of Taiwan under Contract no.93-EC-17-A-08-R7-0312.The authors also thank Dr.J.T.Lee for assistance with sample preparation and XPS analysis.The Industrial Technology Research Institute assisted in meeting the pub-lication costs of this article.References1.N.Takami,T.Ohsaki,H.Hasebe,and M.Yamamoto,J.Electrochem.Soc.,149,A9͑2002͒.2.N.Takami,M.Sekino,T.Ohsaki,M.Kanda,and M.Yamamoto,J.Power Sources,97-98,677͑2001͒.3.G.G.Amatucci,C.N.Schmutz,A.Blyr,C.Sigala,A.S.Gozdz,rcher,andJ.M.Tarascon,J.Power Sources,69,11͑1997͒.4.K.Araki and N.Sato,J.Power Sources,124,124͑2003͒.5. E.Wang,D.Ofer,W.Bowden,N.Iltchev,R.Moses,and K.Brandt,J.Electro-chem.Soc.,147,4023͑2000͒.6. D.Zhang,B.S.Haran,A.Durairajan,R.E.White,Y.Podrazhansky,and B.N.Popov,J.Power Sources,91,122͑2000͒.7.P.Ramadass,B.Haran,R.White,and B.N.Popov,J.Power Sources,112,614͑2002͒.8. A.Funabiki,M.Inaba,Z.Ogumi,S.Yuasa,J.Otsuji,and A.Tasaka,J.Electro-chem.Soc.,145,172͑1998͒.9.M.S.Wu,P.C.Chiang,J.C.Lin,and Y.S.Jan,Electrochim.Acta,49,1803͑2004͒.10.M.S.Wu,P.C.Chiang,and J.C.Lin,J.Electrochem.Soc.,152,A47͑2005͒.11. D.Aurbach,J.Power Sources,89,206͑2000͒.12.S.Zhang,M.S.Ding,K.Xu,J.Allen,and T.R.Richard,Electrochem.Solid-StateLett.,4,A206͑2001͒.13. C.R.Yang,J.Y.Song,Y.Y.Yang,and C.C.Wan,J.Appl.Electrochem.,30,29͑2000͒.14.T.Zheng,A.S.Gozdz,and G.G.Amatucci,J.Electrochem.Soc.,146,4014͑1999͒.15. C.Ho,I.D.Raistrick,and R.A.Huggins,J.Electrochem.Soc.,127,343͑1980͒.16.T.S.Ong and H.Yang,Electrochem.Solid-State Lett.,4,A89͑2001͒.17. A.J.Bard and L.R.Faulkner,Electrochemical Methods;Fundamentals and Ap-plications,p.328,John Wiley&Sons,Inc.,New York͑1980͒.18.G.G.Amatucci,J.M.Tarascon,and L.C.Klein,Solid State Ionics,83,167͑1996͒.19.S.Komaba,N.Kumagai,and Y.Kataoka,Electrochim.Acta,47,1229͑2002͒.20.J.C.Dupin,D.Gonbeau,H.Benqlilou-Moudden,P.Vinatier,and A.Levasseur,Thin Solid Films,384,23͑2001͒.21.T.Ohzuku,Y.Iwakoshi,and K.Sawai,J.Electrochem.Soc.,140,2490͑1993͒.22.N.Takami,A.Satoh,M.Hara,and T.Ohsaki,J.Electrochem.Soc.,142,2564͑1995͒.23.N.Takami,A.Satoh,M.Hara,and T.Ohsaki,J.Electrochem.Soc.,142,371͑1995͒.A1046Journal of The Electrochemical Society,152͑6͒A1041-A1046͑2005͒。

Validation Studies for Concrete Constitutive Models with Blast Test Data

Karagozian & Case, 700 N Brand Blvd, Glendale, CA 91203
Many concrete constitutive models are available for use in LS-DYNA®. A thorough validation related to their applicability for the types of problems at hand should be made before application of any of these models. The process for validating a constitutive model includes examining the results produced with the model related to the behaviors it exhibits, gathering a suite of measured data collection pertinent to the problem to be addressed, and comparisons of measured and computed data. This paper addresses issues related to blast response analyses, which include simplification of boundary conditions (such as support condition and contact interfaces), numerical discretization, and material modeling. It was found important that the strain rate effects should be imposed properly since blast loadings usually excite high frequency and high strain rate responses. The impact of boundary conditions was also identified through the numerical studies.

(2014)Xianhua Hou, Xiaoli Zou et al. RSC ADV.

RSC Advances
PAPER Surfactant CTAB-assisted synthesis of Li1.13[Ni0.233Mn0.534Co0.233]0.87O2 with festoon-like hierarchical architectures as cathode materials for Li-ion batteries with outstandinguction
The introduction of non-aqueous rechargeable lithium-ion batteries in the 1990s for electric vehicles (EV), plug-in hybrid vehicles (PHEVs) and power portable electronic devices bring about a revolution in battery technology and a remarkable transformation from the relatively low-voltage, low-capacity and water-based systems such as nickel–cadmium1,2 and nickel– metal hydride batteries3,4 because of their relatively high energy density and design exibility. The electrochemical performances of the Li-ion batteries mainly depend on the cathode materials, the anode materials and the electrolyte. While the anode materials have the characteristics of the specic capacity much higher than the cathode material, such as siliconbased,5–7 tin-based,8,9 lithium transition metal oxides10,11 and metal-oxides.12–14 So it has become very important to improve the capability of cathode material, and further to improve the capacity of lithium ion secondary batteries.

材料力学双语教学学习资料(英汉对照)

材料力学双语教学学习资料第一章绪论Chapter 1 Introduction§1-1 材料力学的任务The Tasks of Mechanics of Materials1＊. 材料力学: Mechanics of Materials2. 构件: Structural Members3. 变形: Deformation4＊. 强度: Strength5＊. 刚度: Rigidity6＊. 稳定性: Stability§1-2 变形固体的基本假设Fundamental Assumptions of SolidDeformation Bodies1. 连续性假设: Continuity2. 均匀性假设: Homogeneity3. 各向同性假设: Isotropy§1.3 外力及其分类External Forces and Classification1. 分布力: Distributed Force2. 集中力: Point Force3. 静载荷: Static Load4. 动载荷: Dynamic Load§1.4 内力、截面法和应力的概念Concepts of Internal Forces，Method ofSection and Stress1＊. 内力: Internal Force2＊. 截面法: Method of Section3. 截面法的三个步骤：截开，代替，平衡Three steps of method of section: cut off, substitute , and equilibrium.4＊. 应力: Stress5. 平均应力：Average stress6. 应力（全应力）：Whole stress（sum stress）7＊. 正应力: Normal Stress8＊. 剪应力(切应力)：Shearing Stress§1.5 变形与应变Deformation and Strain1．线应变: Strain2．剪应变: Shearing Strain§1.6 杆件变形的基本形式Basic Types of Deformations of Rods1＊. 拉伸或压缩: Tension or Compression2＊. 剪切: Shear3＊. 扭转: Torsion4＊. 弯曲: Bending第二章拉伸、压缩与剪切Chapter 2 Tension，Compression andShear§2.1 轴向拉伸与压缩的概念和实例The Concept and Examples of AxialTension and Compression1. 拉杆: Tensile Rod2. 压杆: Compressive Rod3. 受力特点：外力合力的作用线与杆轴线重合Characteristic of the External Forces: The acting line of the resultant of external forces is coincided with the axis of the rod.4. 变形特点：杆沿轴向伸长或缩短Characteristic of Deformation: Rod will elongate or contract along the axis of the rod.§2.2 轴向拉伸或压缩时横截面上的内力和应力Internal Force and Stress of Axial Tension or Compression on the Cross Section1＊. 横截面: Cross Section2＊. 轴力: Normal Force3＊. 轴力图: Diagram of Normal Force§2.3 直杆轴向拉伸或压缩时斜截面上的应力Stress of Axial Tension or Compressionon the Skew Section1. 斜截面: Skew Section2.ασσα2cos = αστα2s i n 2=§2.4 材料在拉伸时的力学性能Mechanical Properties of Materialswith Tensile Load1. 标准试件: Specimen2. 低碳钢(C ≤0.3%): Low Carbon Steel3. 弹性阶段:Elastic Region4. 屈服阶段:Yielding Stage5. 强化阶段:Hardening Stage6. 颈缩阶段: Necking Stage 7＊.σp ----比例极限: Proportional Limit 8＊.σe ----弹性极限: Elastic Limit 9＊.σs ----屈服极限: Yielding Stress 10＊.σb ----强度极限: Ultimate Stress 11. 延伸率: Percent Elongation12. 断面收缩率: Percent Reduction of Area 13. 塑性材料: Ductile Materials 14. 脆性材料: Brittle Materials 15. 铸铁:Cast iron§2.7 失效、安全系数和强度计算 Failure, Safety factor and Strengthcalculation1＊. 许用应力: Allowable Stress 2. 安全系数: Safety Factor 3＊. 强度条件: Strength Condition][max σσ≤=AF N4＊. 强度校核: Check strength][max σσ≤5＊. 截面设计: Section design][σNF A ≥6＊. 确定许可载荷:Determine allowable load][σA F N ≤§2.8 轴向拉伸或压缩时的变形 Deformation in Axial Tension orCompression1. 弹性变形: Elastic Deformation2. 塑性变形: Plastic Deformation3. 纵向应变: Longitudinal Strainll l l l -=∆=1ε 4. 横向应变: Lateral Straindd d d d -=∆=''ε5．线弹性变形:Linear Elastic Deformation6.泊松比:Poisson’s ratioεεμ'=7＊.弹性模量-E ：表示材料抵抗拉压变形的能力 E - modulus of elasticity ：Indicates the capability of materials for resisting tension or compression 8＊.抗拉刚度-EA :表示构件抵抗拉压变形的能力EA -the axial rigidity: Indicates the capability of constructive members for resisting tension or compression 9＊. 胡克定律（Hooke’s Law ）:当应力不超过材料的比例极限时，应力与应变成正比.The stress is proportional to the strain within the elastic region.εσE =§2.12 应力集中的概念The Concept of Stress Concentration 1.由于截面尺寸的突然变化，使截面上的应力分布不再均匀，在某些部位出现远大于平均值的应力，称应力集中。

LixCoO2(0≤x≤1)结构稳定性和力学性质的第一性原理计算

第49卷第6期 2021年6月硅酸盐学报Vol. 49，No. 6 June ，2021JOURNAL OF THE CHINESE CERAMIC SOCIETY DOI ：10.14062/j.issn.0454-5648.20200537Li x CoO 2(0≤x ≤1)结构稳定性和力学性质的第一性原理计算沈丁1,2，王来贵2，唐树伟1，孙闻1，董伟1，杨绍斌1(1. 辽宁工程技术大学材料科学与工程学院, 辽宁阜新 123000；2. 辽宁工程技术大学力学与工程学院, 辽宁阜新 123000)摘要：采用基于密度泛函的第一性原理，从电子层次计算研究了锂离子电池正极材料Li x CoO 2(0≤x ≤1)的晶体结构、电子结构和力学性质。

结果表明：Li x CoO 2在脱锂过程中晶体结构会发生转变，当x =0.5时由六方结构R 3m 晶型转变为单斜结构P 2/m 晶型，当完全脱锂(x =0)时又转变为六方结构P 3m 1晶型。

随着Li 原子的脱出，Li x CoO 2导带被部分填充，价带被完全填充，金属性质和电子导电性增强，并出现自旋极化。

Li x CoO 2中Co —O 键为含有部分离子性特征的共价键，随着Li 原子的脱出，Co —O 键的离子性特征减弱，共价性特征增强。

随着Li 原子的脱出，Li x CoO 2的体积模量(B )、剪切模量(G )和弹性模量(E )均呈现逐渐减小的趋势，而Poisson 比和各向异性指数逐渐增大。

G/B 值变化趋势表明，LiCoO2呈脆性，脱锂过程中向韧性转变。

关键词：钴酸锂；结构稳定性；力学性质；第一性原理；锂离子电池中图分类号：TM 911, O646 文献标志码：A 文章编号：0454–5648(2021)06–1056–09 网络出版时间：2021-04-07Structural Stability and Mechanical Property of Li x CoO 2 Calculated by First-Principles MethodSHEN Ding 1,2, WANG Laigui 2, TANG Shuwei 1, SUN Wen 1, Dong Wei 1, YANG Shaobin 1(1. College of Materials Science and Engineering, Liaoning Technical University, Fuxin 123000, Liaoning, China2. School of Mechanics and Engineering, Liaoning Technical University, Fuxin 123000, Liaoning, China)Abstract: The structure stability, electronic structure properties and mechanical properties of Li x CoO 2(0≤x ≤1) as cathode material for lithium-ion battery were investigated by the first principles method based on density functional theory. The results show that the crystal structure of Li x CoO 2 changes from hexagonal R 3m structure to monoclinic P 2/m structure as x = 0.5, and then changes to hexagonal P 3m 1 structure as x = 0. The conduction band of Li x CoO 2 is partially filled and the valence band is completely filled with the de-intercalation of Li atom. As a result, the metal property and the electron conductivity of Li x CoO 2 both increase, and the spin polarization phenomenon appears. The Co —O bond in Li x CoO 2 is a covalent bond partially with an ionic bond. The ionic properties of Co —O bond weaken and the covalent properties increase with the de-intercalation of Li atom. The bulk modulus (B ), shear modulus (G ) and Young’s modulus of Li x CoO 2 decrease with the de-intercalation of Li atom, while the Poisson ratio and the anisotropy index increase. The change of G/B value shows that LiCoO 2 is brittle and then changes to ductile during de-intercalation process.Keywords: lithium cobaltite; structural stability; mechanical property; first-principles; lithium-ion battery锂离子电池电极材料在充放电过程中，锂离子嵌入或脱出时发生电化学反应，电极材料的晶体结构发生一系列变化，导致此过程产物的力学性质发生变化，从而影响电极材料的结构稳定性和储锂性能。

AN APPARATUS FOR PROTECTING STRUCTURAL SUPPORTS

专利名称：AN APPARATUS FOR PROTECTING STRUCTURAL SUPPORTS发明人：Roller, Joseph A.申请号：EP98939374.9申请日：19980813公开号：EP1003937A1公开日：20000531专利内容由知识产权出版社提供摘要：An apparatus for protecting structural supports from damage when impacted by an object such as a moving vehicle is provided. The apparatus has a shaped component which in the preferred embodiment is a semi- cylindrical component having a body defined by a wall, a top, and a base. The wall has at least one flat wall face and surrounds a hollow interior. An indentation for receiving a structural support is present in at least one flat wall face. A means for securing the shaped component to the structural suppport and for firmly seating one component at its flat wall face against the second component at its flat wall face when two components are present. Preferably, each component has a plurality of impact absorbing indentations, each having an aperture, a base and a wall extending from the base to an aperture mouth. These indentations function to re-distribute the energy of impact when a collusion occurs between the apparatus and a moving object. Preferably, each component is formed by rotational molding from a plastic resin.申请人：Roller, Joseph A.地址：4736 Diane Drive Ashtabula, OH 44004 US国籍：US代理机构：Winter, Brandl & Partner 更多信息请下载全文后查看。

临界屈曲应变能力英文

临界屈曲应变能力英文English:The critical buckling strain capacity refers to the ability of a material to withstand large deformations before reaching the point of structural instability, known as buckling. This property is crucial in determining the overall stability and strength of structural components under compressive loads. The critical buckling strain capacity is influenced by various factors including material type, geometry, and the presence of imperfections or defects. It is commonly evaluated through experimental testing and numerical simulations to ensure that the structure can withstand expected loads without buckling.中文翻译:临界屈曲应变能力是指材料在达到结构不稳定点（即屈曲）之前能够承受大变形的能力。

这种性质对于在受压载荷下结构构件的整体稳定性和强度至关重要。

临界屈曲应变能力受材料类型、几何形状以及缺陷或不完美的影响。

通常通过实验测试和数值模拟来评估这一能力，以确保结构在受预期载荷时不会发生屈曲。

混凝土结构设计中常见的稳定问题

第 40 卷第 1 期2024 年2 月结构工程师Structural Engineers Vol. 40 ， No. 1Feb. 2024混凝土结构设计中常见的稳定问题刘建飞1，*周袁凯1袁媛2（1.中国建筑标准设计研究院有限公司，北京 100048； 2.中移建设有限公司，北京 100038）摘要混凝土结构的稳定问题受重视程度不及钢结构，本文对比中国、欧洲、美国和新西兰的混凝土设计规范中关于结构整体稳定、构件稳定的条文，并结合示例和公式推导，给出相关设计建议。

对于整体稳定，刚重比数值受侧向合力位置、建筑重力大小及位置的影响，不应简单地以刚重比的数值来评判不同建筑的整体稳定性。

对于构件稳定，分析了几个常见而易被忽视的问题。

穿层柱及支撑连桥的框架柱宜考虑二阶效应；对于大跨度梁和长悬臂梁可借鉴外国规范，控制梁的高宽比和侧向无支撑长度；通过数值屈曲分析来反算柱子计算长度时，应正确施加荷载；国内对墙肢的稳定限制较严，对比墙肢的稳定等效荷载与轴压比限值荷载，给出了楼梯间外墙和L形剪力墙的截面由稳定控制的临界高度。

关键词整体稳定，二阶效应，刚重比，墙肢稳定，穿层柱Common Stability Problems for Design of Concrete StructuresLIU　Jianfei1，*ZHOU Yuankai1YUAN Yuan2（1.China Institute of Building Standard Design and Research Co. LTD.， Beijing 100048， China；2.China Mobile Jianshe Co.LTD.， Beijing 100038， China）Abstract The stability of concrete structure is not considered as vital as that of steel structure. This paper compares the provisions on overall stability of structure and members’stability in concrete design codes of China， Europe， USA and New Zealand， some relevant design suggestions are given by combining examples and formula deduction. For the overall stability，the rigid-gravity ratio value is affected by the resultant position of lateral forces，the magnitude and position of building gravity. The overall stability of different buildings should not be simply judged by the stiffness-weight ratio. For the stability of members，several common but easily ignored problems are analyzed：second-order effect should be considered for columns supporting bridges and cross-story columns；by reference foreign codes，the height-width ratio and lateral unbraced length of beam should be controlled for long-span and long cantilever beams；loads should be applied correctly when solving the effective length of columns by buckling analysis. The stability of shear walls is strictly restricted in China. In this paper， the critical height of outer wall in staircase and L-shaped shear wall limited by stability is given by comparing the equivalent load of stability of wall limbs with the limit load of axial load ratio.Keywords overall stability， second-order effect， stiffness-weight ratio， stability of shear walls， cross-story columns收稿日期：2023-02-20*联系作者：刘建飞，男，正高级工程师，主要从事复杂结构的抗震设计与研究。

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Journal of Solid State Chemistry 168,60–68(2002)doi:10.1006/jssc.2002.9679Structural Stability of LiCoO 2at 4001C 1Y.Shao-Horn,n ,2S.A.Hackney,n A.J.Kahaian,w and M.M.Thackeray wnDepartment of Materials Science and Engineering,Michigan Technological University,Houghton,Michigan 49931;and w Electrochemical TechnologyProgram,Chemical Technology Division,Argonne National Laboratory,Argonne,Illinois 60439Received April 1,2002;accepted June 9,2002The relative stability of the lithiated-spinel structure,Li 2[Co 2]O 4,at 4001C to the layered LiCoO 2structure has been investigated.‘‘Low-temperature’’LT-LiCoO 2samples were synthesized at 4001C by the solid-state reaction of Li 2CO 3with CoCO 3(or Co 3O 4)for various times between 10min and 232days.Least-squares reﬁnements of X-ray powder diffraction patterns were used to determine the fractions of lithiated-spinel Li 2[Co 2]O 4and layered LiCoO 2in the samples.X-ray powder diffraction and transmission electron microscope data show that Li 2[Co 2]O 4nucleates from an intermediate Li x Co 1Àx [Co 2]O 4spinel product before transforming very slowly to layered LiCoO 2.The experimental data conﬁrm the theoretical predic-tion that layered LiCoO 2is thermodynamically more stable than the lithiated-spinel structure at 4001C and support the arguments that a non-ideal cation distribution in Li 2[Co 2]O 4,non-stoichiometry and kinetic factors restrict the transformation of the lithiated-spinel structure to layered LiCoO 2at this temperature.#2002Elsevier Science (USA)Key Words:lithium–cobalt–oxide;spinel;structure,phase transition;X-ray diffraction;transmission electron microscopy;electron diffraction;lithium batteries.INTRODUCTIONA ‘‘low-temperature’’form of LiCoO 2(LT-LiCoO 2),synthesized at 4001C,has been studied extensively as a positive electrode material for lithium batteries (1–10).X-ray and neutron powder diffraction data and vibrationspectroscopy studies have determined that LT-LiCoO 2has the lithiated-spinel Li 2[Co 2]O 4structure (space group Fd 3m )(10).LT-LiCoO 2electrodes do not perform well in lithium cells,and there is considerable hysteresis between the lithium de-intercalation and intercalation processes (1–7).This behavior contrasts with the superior electro-chemical performance of electrodes with the layered ‘‘high-temperature’’HT-LiCoO 2structure (space group R %3m )that can be synthesized at 8001C (1,11).We have shown previously with electron diffraction data that the cobalt distribution in LT-LiCoO 2samples is not ideal,and that it can be considered to be intermediate between an ideal layered (Li)3a {Co}3b O 2conﬁguration and an ideal lithiated-spinel (Li 2)16c [Co 2]16d O 4conﬁguration,in which cation layers with 75%Co and 25%Li alternate with layers containing 25%Co and 75%Li.Therefore,it has been proposed that in LT-LiCoO 2structures,the presence of some cobalt ions on the octahedral 16c sites would hinder lithium diffusion in the interstitial space of the [Co 2]O 4spinel framework and deleteriously affect the electrochemical performance of Li/LT-LiCoO 2cells (5,10).Such a structural phenomenon raises questions about the relative thermodynamic stability of the lithiated-spinel versus the layered LiCoO 2structure at 4001C.We have,therefore,investigated the relative stability of these two structure types at this temperature.Three experimental observations of the phase stability of LiCoO 2samples heated between 4001C and 9001C have previously shown that (1)at 4001C,the lithiated-spinel structure is the dominant phase (1–11);(2)the fraction of layered LiCoO 2increases as the samples are heated at intermediate temperatures,e.g.,5001C (9);and (3)a single-phase,layered LiCoO 2structure is obtained between 8001C and 9001C (1,11).These experimental results imply that at 4001C the lithiated-spinel is the thermodynamically pre-ferred phase.However,quantum mechanical calculations have predicted that at all temperatures the lithiated-spinel structure is thermodynamically unstable relative to the layered structure (12),yered LiCoO 2is the ground state structure.Moreover,the lithiated-spinel structure1The submitted manuscript has been created by the University of Chicago as Operator of Argonne National Laboratory (‘‘Argonne’’).Part of this research was performed under the auspices of the US Department of Energy under Contract W-31-109-ENG-38.The US Government retains for itself,and others acting on its behalf,a paid-up,non-exclusive,irrevocable worldwide license in the said article to reproduce,prepare derivative works,distribute copies to the public,and perform publicly and display publicly,by or on behalf of the Government.2To whom correspondence should be addressed.Present address:Department of Mechanical Engineering,Massachusetts Institute of Technology,77Massachusetts Avenue,Cambridge,MA 02139.E-mail:shaohorn@.600022-4596/02$35.00r 2002Elsevier Science (USA)All rights reserved.(cubic,Fd3m)has higher symmetry than the layered structure(trigonal,R%3m).This observation contradicts the general rule of thumb for phase transitions that low-temperature,thermodynamically stable phases tend to have lower symmetry than high-temperature,stable phases(13). We have,therefore,undertaken an investigation of the relative thermodynamic stability of the lithiated-spinel Li2[Co2]O4structure versus the layered LiCoO2structure at4001C.In this paper,we report data concerning the reactions of Li2CO3with CoCO3(or Co3O4)that occur between10min and232days of heat treatment at4001C.In particular,we have determined the relative volume frac-tions of the lithiated-spinel and layered structures as a function of heat-treatment time.Structural analyses of the products by powder X-ray diffraction and transmission electron microscopy have allowed us to propose a reaction pathway for the formation of Li2[Co2]O4at4001C and an explanation for the slow and sluggish transformation of the lithiated-spinel structure to the layered structure.Our interpretation is consistent with the poor electrochemical behavior of LT-LiCoO2electrodes in lithium cells(5,10).EXPERIMENTALLT-LiCoO2samples were prepared by the solid-state reaction of stoichiometric amounts of Li2CO3(Aldrich) and CoCO3(Aldrich)at4001C in air for various times that ranged from10min to232days.These samples are referred to as LT-LiCoO2-400-X,where X is the reaction time (days).A CoCO3reference was heated at4001C in air for 10min in the absence of Li2CO3to demonstrate its rapid decomposition at this temperature.An LT-LiCoO2sample was also prepared at5001C over4days(LT-LiCoO2-500-4).The evolution of lithiated-spinel Li2[Co2]O4and layered LiCoO2structures was also investigated by reacting Li2CO3with submicron-size Co3O4powder(Aldrich)at 4001C for1and4days;these samples are referred to as LT-LiCoO2–Co3O4-400-1and LT-LiCoO2–Co3O4-400-4, respectively.All of the X-ray diffraction data were collected on a Siemens D5000powder diffractometer with Cu K a radiation at room temperature(221C).The LT-LiCoO2patterns used for structural reﬁnements were recorded between101and 11012y using a step size of0.0312y and a step time of15s. The powder X-ray diffraction proﬁles were reﬁned with the GSAS program(14)using a Pseudovoigt function(15)in which the peak-shape parameters W(peak width)and X (peak tail)were varied.With these reﬁnements,the ‘‘goodness ofﬁt’’between the observed and calculated data is given by the weighted residual factor,w R p(15).The lattice parameters and oxygen positional parameters of the lithiated-spinel and layered structures were held constant during the two-phase reﬁnements and quantitative analysis of the LT-LiCoO2patterns.The phase transformation from the lithiated-spinel structure to the layered structure in LT-LiCoO2samples was investigated by imaging and convergent-beam electron diffraction with a transmission electron microscope(JEOL-JEM4000FX-1)under an accelerating voltage of200keV. Sample preparation for these microscopy studies has been described elsewhere(16).RESULTS AND DISCUSSIONX-Ray DiffractionThe X-ray diffraction patterns obtained after heating CoCO3in air in the absence of Li2CO3,and together with Li2CO3,for10min are shown in Figs.1(a)and(b), respectively.The patterns for LT-LiCoO2samples heated at4001C for0.5,1and2days are shown in Figs.1(c–e), and in Figs.2(a–e)for10,39,67,131and232days.In the absence of Li2CO3,CoCO3decomposes very rapidly to CO3O4with a spinel structure at4001C(Fig.1(a)).When Li2Co3is present,the X-ray diffraction patterns show that the decomposition of CoCO3proceeds at a much faster rate than that of Li2CO3.In addition,the subsequent growth of a lithiated-spinel phase,Li2[Co2]O4(LS),is derived predominantly from a Co3O4intermediate pro-duct(Figs.1(b–e))that may contain a small amount of substituted lithium on the tetrahedral sites,namely Li x Co1Àx[Co2]O4(x r0:4)(17).The growth ofthe FIG.1.X-ray diffraction patterns:(a)CoCO3heated at4001C for 10min;(b)–(e)LT-LiCoO2samples heat treated at4001C for10min,12, 24and48h,respectively.The{220},{511}and{440}peaks of the spinel phase(S)and the{511}and{440}peaks of the lithiated-spinel phase(LS) are indicated in(a)and(e),respectively.STRUCTURAL STABILITY OF LiCoO261Li 2[Co 2]O 4phase is evident (particularly when the X-ray patterns are enlarged)from (1)a decrease in the magnitude of the {220}spinel peak (S)at B 31.512y ;and (2)the evolution of independent {511}and {440}peaks at slightly higher 2y values (near 591and 6612y ;respectively)which grow at the expense of the {511}and {440}Li x Co 1Àx [Co 2]O 4spinel peaks.As shown in Fig.1(e),after 2days,the reaction is still far from complete.For prolonged heating (Z 10days),the dominant peaks in the patterns of the LT-LiCoO 2samples could be indexed to the lithiated-spinel Li 2[Co 2]O 4structure (Figs.2(a–e)).Note that there is no evidence of the {220}Li x Co 1Àx [Co 2]O 4spinel peak in Figs.2(b–e).The most noticeable changes in the X-ray diffraction patterns of LT-LiCoO 2samples are observed between 581and 7112y ;as shown in Fig.3,in which the X-ray diffraction pattern of the LT-LiCoO 2-500-4sample is also given for comparison (Fig.3(f)).In Figs.3(a–e),it is clear that the {440}peak of the lithiated-spinel phase (Fd 3m )gradually splits on prolonged heating into the (108)and (110)peaks of thelayered LiCoO 2structure ðR %3m Þ;the {511}peak of the lithiated-spinel structure shifts to the left to generate the overlapping (107)and (009)peaks of layered LiCoO 2.These effects are more pronounced in the LT-LiCoO 2-500-4sample (Fig.3(f)).The data clearly indicate that the transformation of the lithiated-spinel LT-LiCoO 2structure to layered LiCoO 2at 4001C is sluggish,and that it can beachieved more rapidly,but not completely,after reaction for 5days at 5001C.The lattice parameters and the oxygen positional parameter of the lithiated-spinel Li 2[Co 2]O 4and the layered LiCoO 2structures were initially obtained from reﬁnements of the LT-LiCoO 2-400-10and LT-LiCoO 2-500-4patterns,respectively (Table 1);they are consistent with those reported previously (1–7,11).These parameter values were held invariant in the subsequent quantitative analyses of these phases in the LT-LiCoO 2X-ray diffrac-tion patterns,the results of which are shown in Table 2.Although the quality of the X-ray patterns does not allow an accurate determination of the volume fraction of the lithiated-spinel and layered LiCoO 2components,the results show a clear trend.As expected,the lithiated-spinel structure is the major component in the LT-LiCoO 2-400-10FIG.2.X-ray diffraction patterns of LT-LiCoO 2samples heated at 4001C for (a)10d,(B)39d,(c)67d,(d)131d and (e)232d.The major peaks in (e)are indexed to the lithiated-spinel Li 2[Co 2]O 4structure.Peaks from unreacted Li 2CO 3are indicated byasterisks.FIG.3.X-ray diffraction patterns between 581and 7112y ;(a)–(e)LT-LiCoO 2-400samples shown in Figs.2(a–e);(f)LT-LiCoO 2-500-4,showing the growth of layered LiCoO 2over long reaction times and at higher temperature.TABLE 1Structural Parameters of the Lithiated-Spinel Li 2[Co 2]O 4and Layered LiCoO 2Structures Determined from the Reﬁne-ment of the X-Ray Diffraction PatternsPhasesSpace group Lattice parameters (nm)Oxygen position parameterLithiated-spinel Li 2[Co 2]O 4Fd 3ma ¼0:79950.259Layered LiCoO 2R %3m a ¼0:28160.240c ¼1:404462SHAO-HORN ET AL.sample,whereas in the LT-LiCoO2-500-4sample,the layered structure predominates.It is also clear from Table2that on prolonged heating of LT-LiCoO2samples at4001C,the volume fraction of the lithiated-spinel phase in the sample decreases gradually with respect to the amount of layered LiCoO2.Although incomplete,the slow transformation of the[Co2]O4spinel framework to a layered conﬁguration implies that layered LiCoO2is thermodynamically more stable than the lithiated-spinel Li2[Co2]O4structure at4001C.Thisﬁnding provides experimental evidence to support the theoretical prediction of the relative stability of these two structure types(12). In the structural analyses of LT-LiCoO2samples (Table3),the peak proﬁle parameters revealed that(1) after heating for232days at4001C,the peak width parameter(W)of the lithiated-spinel phase(108.6)is signiﬁcantly larger than that of the layered LiCoO2phase (21.1);(2)the value of W for the layered LiCoO2 component at4001C(21.1)is signiﬁcantly larger than it is at5001C(0.1);and(3)at5001C,the layered component (21.1)is substantially more crystalline than the lithiated-spinel component(86.2).These points are consistent with the expectation that low synthesis temperatures should provide small crystals that increase in size and crystallinity on further heating;they also emphasize the difﬁculty in synthesizing a highly ordered,lithiated-spinel structure Li2[Co2]O4by solid-state reactions,even at moderate temperatures.The transformation from the lithiated-spinel to the layered structure is kinetically sluggish because,after131 days at4001C,there is still57%of the lithiated-spinel phase in the sample and further heating appears to have little effect.It is therefore believed that the following structural factors limit this phase transition at4001C:(1) the ideal Li2[Co2]O4and LiCoO2structures have ordered rocksalt-type conﬁgurations in which all the octahedral sites areﬁlled,(2)the transformation of Li2[Co2]O4to layered LiCoO2requires that there must be an exchange of one quarter of the cobalt and lithium ions within the cubic-close packed oxygen array,and(3)there is no readily available and energetically favorable interstitial space in the rocksalt structures for Li and Co diffusion. Transmission Electron Microscopy StudiesA transmission electron microscope image of the LT-LiCoO2-400-39sample,shown in Fig.4,reveals equiaxed and rod-like crystal shapes.Most of the smaller crystals, which range in size from20to80nm,are equiaxed, whereas the larger crystals,typically100–300nm in length, are rod shaped.Single-crystal electron diffraction revealed that the equiaxed crystals have the lithiated-spinel Li2 [Co2]O4structure,whereas the rod-shaped crystals have the layered LiCoO2structure,with the rod axis being contained within the closed-packed(003)planes.These observations are consistent with the smaller peak-width parameter(W)of the layered LiCoO2component obtained from the reﬁnement of the X-ray diffraction data(Table3). The development of these large,rod-shaped crystals is referred to as‘‘abnormal’’growth.This contrasts with the ‘‘normal’’growth of the small Li2[Co2]O4crystals on heating at4001C from10days to39days(Figs.5(a)and (b))during which they remain equiaxed.Note,however, that some small crystals show the onset of rod-like character.The abnormal growth of the LiCoO2crystals implies that the Gibbs volume free energy of layered LiCoO2is lower than that of the lithiated-spinel Li2 [Co2]O4;this provides the driving force to reduce the surface free energy and to promote the growth of the layered LiCoO2crystals.The transmission electron micro-TABLE2Volume Fractions of the Lithiated-Spinel and Layered Components in LT-LiCoO2Samples,Obtained from Two-Phase Reﬁnements of the X-Ray Diffraction PatternsLiCoO2samples Volume fractionof Li2[Co2]O4Volume fractionof LiCoO2Sigma w R p4001C F10days0.930.070.06830.21834001C F39days0.670.330.00470.10974001C F67days0.610.390.00380.10384001C F131days0.570.430.00340.09764001C F232days0.560.440.00520.12545001C F4days0.240.760.00180.1036TABLE3Peak Proﬁle Parameters Determined from the Reﬁnements of the X-Ray Diffraction PatternsLithiated-spinel Lithiated-spinel Layered Layered LiCoO2samples W X W X4001C F10days204.231.721.1(not reﬁned)28.2(not reﬁned) 4001C F39days104.219.921.1(not reﬁned)28.2(not reﬁned) 4001C F67days97.118.221.1(not reﬁned)28.2(not reﬁned) 4001C F131days106.317.321.1(not reﬁned)28.2(not reﬁned) 4001C F232days108.615.721.128.25001C F4days86.218.40.117.8STRUCTURAL STABILITY OF LiCoO263scope studies are consistent with the X-ray diffraction reﬁnements and theoretical calculations;they provide further evidence for a gradual transition of the lithiated-spinel structure to the layered structure at 4001C and for layered LiCoO 2being thermodynamically more stable than Li 2[Co 2]O 4at 4001C.Proposed Reaction for Li 2[Co 2]O 4FormationIf the layered LiCoO 2structure is the thermodynamically preferred structure at 4001C,then the presence of a lithiated-spinel structure after prolonged heating of CoCO 3and Li 2CO 3at this temperature implies that the formation of a spinel conﬁguration is kinetically preferred.Therefore,we propose the following reaction route for the formation of a lithiated-spinel structure,Li 2[Co 2]O 4,from Li 2CO 3and CoCO 3precursors:Reaction I F lithiation and oxidation of CoCO 3:ð6À2x ÞCoCO 3þx Li 2CO 3þð1þx =2ÞO 232Li x Co 1Àx ½Co 2 O 4þð6Àx ÞCO 2ð0r x r 0:4Þ:Reaction II F subsequent formation of Li 2[Co 2]O 4:Li x Co 1Àx ½Co 2 O 4þð32Àx ÞLi 2CO 3þð14Àx =2ÞO 23ð32Àx =2ÞLi 2½Co 2 O 4þð32Àx ÞCO 2:Reaction I is consistent with the experimental X-ray data in Figs.1and 2that show convincing evidence that anintermediate Li x Co 1Àx [Co 2]O 4spinel phase ð0r x r 0:4Þis nucleated prior to the formation of the lithiated-spinel (rocksalt)Li 2[Co 2]O 4structure.Because layered LiCoO 2appears to be the thermodynamically favored structure at 4001C,we conclude that the presence of anintermediateFIG.4.A transmission electron microscope image of an LT-LiCoO 2-400-39sample showing two crystal shapes:equiaxed and rod like(circled).FIG.5.Transmission microscope images showing the crystal growth and change in morphology on heating LT-LiCoO 2-400samples from (a)10d to (b)39d.64SHAO-HORN ET AL.Li x Co1Àx[Co2]O4spinel structure is essential for the formation of a metastable lithiated-spinel Li2[Co2]O4 product;in this respect,the possibility that a spinel intermediate structure is formed has already been pro-posed,but not demonstrated,by Garcia et al.(9). Reaction of Co3O4with Li2CO3The reaction of submicron-size Co3O4particles with Li2CO3was studied in an attempt to gather more information about the reaction route by which LT-LiCoO2 structures are formed from a spinel precursor and,in particular,about the crystallographic orientation relation-ships that develop between the structures of the precursor and the LiCoO2products.No change in the X-ray powder diffraction pattern was found in the LT-LiCoO2–Co3O4-400-1sample,i.e.,after heating Co3O4and Li2CO3for1 day at4001C,whereas substantial changes were observed after heating for4days as shown in Fig.6.In the latter case,in addition to unreacted Li2CO3,the X-ray diffrac-tion pattern showed two cubic phases that are clearly evident from the doublet peaks at approximately451 and6612y;and could be attributed to Li2[Co2]O4 (a¼0:7995nm)and a spinel component Li x Co1Àx[Co2]O4ð0r x r0:4Þ;a¼0:808020:8089nm).The strong singlet spinel peak at approximately31.512y;which corresponds to the{220}reﬂections,provides unequivocal evidence for tetrahedral-site cobalt in the Li x Co1Àx[Co2]O4phase;this peak is forbidden by the Li2[Co2]O4rocksalt structure.The presence of Li2[Co2]O4in LT-LiCoO2–Co3O4-400-4was conﬁrmed by comparing the X-ray diffraction data near 191and6612y with LT-LiCoO2-400-232and LT-LiCoO2-500-4samples(Figs.7and8,respectively).It is evident that LT-LiCoO2–Co3O4-400-4contains an Li x Co1Àx[Co2]O4 FIG.6.X-ray diffraction pattern of LT-LiCoO2–Co3O4-400-4,show-ing a mixture of a spinel phase,a lithiated-spinel phase(both marked S),and unreacted Li2CO3(L).FIG.7.X-ray diffraction peaks near1912y for(a)LT-LiCoO2-400-232,(b)LT-LiCoO2–Co3O4-400-4,and(c)LT-LiCoO2-500-4.FIG.8.X-ray diffraction peaks near6612y for(a)LT-LiCoO2-400-232,(b)LT-LiCoO2–Co3O4-400-4,and(c)LT-LiCoO2-500-4.STRUCTURAL STABILITY OF LiCoO265spinel component and an Li 2[Co 2]O 4lithiated-spinel component.There is also evidence of a small amount of layered LiCoO 2in LT-LiCoO 2–Co 3O 4-400-4(Fig.8).The crystallographic relationships between the spinel and lithiated-spinel (rocksalt)components in the LT-LiCoO 2–Co 3O 4-400-4sample were examined by transmis-sion electron microscopy;a typical image of a lithiated Co 3O 4crystal is shown in Fig.9.Electron diffraction patterns collected from the surface and the center of the crystal are shown in Figs.10(a)and (b),respectively.Although both patterns show six-fold symmetry,there are marked differences in the scattering intensities of the {220}reﬂections from the face-centered-cubic unit cells.The absence of the {220}reﬂections in Fig.10(a)is consistent with the formation of the lithiated-spinel Li 2[Co 2]O 4structure on the crystal surface,whereas Fig.10(b)shows the set of these reﬂections,which is indicative of an Li x Co 1Àx [Co 2]O 4spinel structure within the bulk of the crystal.The absence of distortion of the reciprocal lattice points in Figs.10(a)and (b)implies that during the crystallographic transformation at the phase boundary,the crystallographic planes of the lithiated-spinel Li 2[Co 2]O 4structure remain parallel to those of Li x Co 1Àx [Co 2]O 4.This is perhaps not too surprising because the transition takes place between two cubic phases with only a small change (B 1%)in the lattice parameter.By contrast,such a strong crystallographic relationship does not exist between cubic Li 2[Co 2]O 4andlayered LiCoO 2,which has trigonal symmetry ðR %3m Þ:In this case,the growth of layered LiCoO 2from the Li 2[Co 2]O 4-rich surface of the lithiated Co 3O 4particle is far more dramatic,as shown in Fig.11.The large crystal-lite was conﬁrmed unequivocally by electron diffraction to have trigonal symmetry,as required by layered LiCoO 2(Fig.12(a)),whereas the symmetry of the bulkstructureFIG.9.A transmission electron microscope image of a lithiated Co 3O 4crystal in an LT-LiCoO 2–Co 3O 4-400-4sample.FIG.10.Single-crystal electron diffraction patterns collected from(a)the surface and (b)the center of the lithiated Co 3O 4crystal as shown in Fig.9.66SHAO-HORN ET AL.was clearly still cubic (Fig.12(b)).The dramatic crystal growth and change in crystal morphology suggest that long-range diffusion of cobalt and lithium (requiring a high activation energy)is needed to bring about the signiﬁcant change in distribution of lithium and cobalt during the rocksalt-to-rocksalt transition of Li 2[Co 2]O 4to layered LiCoO 2,whereas only short-range diffusion of cobalt,lithium and oxygen (with lower activation energy)is required to nucleate the growth of Li 2[Co 2]O 4on the surface of Li x Co 1Àx [Co 2]O 4.Therefore,when CoCO 3is reacted with Li 2CO 3at 4001C,a kinetically favored pathway to the lithiated-spinel conﬁguration Li 2[Co 2]O 4from an intermediate Li x Co 1Àx [Co 2]O 4spinel conﬁgura-tion can account for the formation of this structure type over the thermodynamically favored layered LiCoO 2structure.CONCLUSIONStudies of the reaction of CoCO 3with Li 2CO 3at 4001C have shown that the lithiated-spinel Li 2[Co 2]O 4is initially formed from an intermediate Li x Co 1Àx [Co 2]O 4spinel product before transforming gradually to layered LiCoO 2,thereby conﬁrming theoretical calculations that the layered structure is thermodynamically favored at this tempera-ture.The rocksalt-to-rocksalt transition from Li 2[Co 2]O 4to layered LiCoO 2at 4001C is kinetically hindered,predominantly by structural factors.The lack of an energetically favorable interstitial space,the difﬁculty of exchanging 25%of the lithium ions in the 16c sites of the spinel structure with cobalt ions (and vice versa),and the non-ideal distribution of lithium and cobalt in Li 2[Co 2]O4FIG.11.A transmission electron microscope image of a lithiated Co 3O 4crystal in an LT-LiCoO 2–Co 3O 4-400-4sample,showing the growth of layered LiCoO 2on the crystalsurface.FIG.12.Single-crystal electron diffraction patterns of (a)the trigonal (layered)LiCoO 2crystal shown at the particle surface in Fig.11,and (b)a cubic Li x Co 1Àx [Co 2]O 4spinel crystal in the bulk.STRUCTURAL STABILITY OF LiCoO 267are believed to contribute to the slow and sluggish transformation,and therefore to the apparent stability of Li2[Co2]O4at4001C.The difﬁculty in preparing pure,single-phase Li2[Co2]O4can be attributed to a wide range of solid solution and structural conﬁgura-tions that exist among rocksalt compositions in the Li x Co1Àx O system(0:4r x r0:5).These conclusions support previous reports that the structural features of LT-Li2[Co2]O4electrodes are a major reason for the poor electrochemical properties of room-temperature Li/LT-Li2[Co2]O4cells.ACKNOWLEDGMENTSJohn Vaughey is thanked for collecting X-ray diffraction data from various samples.Support from the US Department of Energy,Chemical Sciences Division,Ofﬁce of Basic Energy Sciences under Contract No. W31-109-Eng-38is gratefully acknowledged.REFERENCES1.R.J.Gummow,M.M.Thackeray,W.I.F.David,and S.Hull,Mater.Res.Bull.27,327(1992).2.R.J.Gummow and M.M.Thackeray,Solid 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