Nanostructured high-energy cathode materials for advanced lithium batteries

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液氯中微量水分几种测量方式的对比

山东化工SHANDONGCHEMICALINDUSTRY -136 -2021年第50卷液氯中微量水分几种测量方式的对比林君1,范长春2,陈娟2，王鹏2,陈为洪2(1.扬州联博药业有限公司，江苏扬州225009；2.宁夏瑞泰科技股份有限公司，宁夏中卫755550)摘要:液氯是化工生产的重要原料，由于公司没有电解盐水制，所氯外购，槽车运输。

检测方法是用钢瓶氯，到实验室 GB/T 5138—2006中重(简法)的方法。

从期检测，液氯含水一般都大于053%( 要W 0N4%),检测高。

通过查询资料，国外量法检测液氯中的水含量，测定法测定 1/10。

鉴于情况制0定了一些试验方案，来进行几测量方比，为同行进行测量方。

关键词：液氯含水；电解法;激光法中图分类号：0657 文献标识码：A 文章编号：1008-021X ( 2021) 05-0136-03氯气是由氯化钠电解产生，刚解氯气高,伴有大量水气及夹质。

湿氯气对钢铁及大金属有腐蚀作用，决定了如压缩机、管、仪表及容器等腐蚀，关产设备及。

统计一般气化氯气中微量水的控制 <0固1%,工业氯国家水量：优级；0.01%，一；0.03%，合；0.04%[1]o 氯气含水对于日产而言是极大隐患，公司量水作计控点来控制。

测定氯气含水的方法有很多有传统法、激光法解法。

主要通过方法的比氯含水量的测定提法。

1 收法m 1.1测定原理气化通过已量的五氧化二管，中水分。

用已量的氢氧化钠瓶，氯气，并分别称量管瓶，与测定质量差，计：水分含量。

化学：P 2O 5 + 3H 2O = 2H 3PO 41-2试剂及溶剂五氧化二磷:分析纯;氢氧化钠溶液：300 g/L o 干燥的氮气。

1.3仪器五氧化二磷法测定水分装置图如图1所示。

1.氢气缓冲瓶;2・废气处理瓶;3二型干燥管内填装五氧化二磷；4.锐量计;5.氯气瓶；6,7,8.旋塞;9.胶管图1分析装置图1.4 测定(1) 两 U 型干燥管洗净，烘干后装入五氧化二磷，再将磨口塞涂，塞好，放干燥器中备用。

锂离子-Materials for high-energy density batteries

Chapter14Materials for High-energy Density Batteries Arumugam ManthiramAbstract Lithium-ion batteries have emerged as the choice of rechargeable power source as they offer much higher energy density than other systems.However, their performance factors such as energy density,power density,and cycle life depend on the electrode materials employed.This chapter provides an overview of the cathode and anode materials systems for lithium-ion batteries.After pro-viding a brief introduction to the basic principles involved in lithium-ion cells, the structure-property-performance relationships of cathode materials like layered LiMO2(M=Mn,Co,and Ni)and their soiled solutions,spinel LiMn2O4,and olivine LiFePO4are presented.Then,a brief account of the carbon,alloy,oxide, and nanocomposite anode materials is presented.14.1IntroductionBatteries are the main energy storage devices used in modern society.They power invariably the portable devices we use in our daily life.They are also being pursued and developed intensively for widespread automobile applications.With respect to the topic of this book,batteries are also critical to store the energy harvested from various sources like solar,wind,thermal,strain,and inertia and to use the harvested energy efﬁciently when needed.In all of these applications,the energy density of the battery,which is the amount of energy stored per unit volume(W h/L)or per unit weight(W h/kg),is a critical parameter.The amount of energy stored depends on the capacity(amount of charge stored)per unit volume(A h/L)or per unit weight(A h/kg)and the voltage(V)each cell can deliver.In addition,rechargeability,charge–discharge cycle life,and the rate at which the cell can be charged and discharged are also important parameters.Moreover,cost and environmental impact considerations need to be taken into account.All of these parameters and properties are related to the battery chemistry and materials involved.A.Manthiram(B)Electrochemical Energy Laboratory,Materials Science and Engineering Program,The University of Texas at Austin,Austin,TX78712,USAe-mail:rmanth@365 S.Priya,D.J.Inman(eds.),Energy Harvesting Technologies,DOI10.1007/978-0-387-76464-114C Springer Science+Business Media,LLC2009366A.Manthiram50100150200250G r a v i m e t r i c E n e r g y D e n s i t y (W h /k g )Volumetric Energy Density (Wh/L)Fig.14.1Comparison of the gravimetric and volumetric energy densities of lithium–ion batteries with those of other rechargeable systemsAmong the various rechargeable battery chemistries known to date,lithium–ion batteries offer the highest energy density when compared with the other recharge-able battery systems such as lead–acid,nickel–cadmium,and nickel–metal hydride batteries as shown in Fig.14.1.The higher volumetric and gravimetric energy den-sities of the lithium–ion cells are due to the higher cell voltages (∼4V)achievable by the use of non-aqueous electrolytes in contrast to <2V achievable with most of the aqueous electrolyte-based cells.This chapter,after brieﬂy providing the basic principles involved in lithium–ion cells,focuses on the various materials systems employed or being currently pursued to maximize the energy density,power den-sity,or both,while keeping in mind the other parameters like cycle life,cost,and environmental concerns.14.2Principles of Lithium–Ion BatteriesLithium–ion batteries involve a reversible insertion/extraction of lithium ions into/from a host matrix during the discharge/charge process as shown in Fig.14.2.The host matrix is called as a lithium insertion compound and serves as the elec-trode material in the cell.The present generation of lithium–ion cells mostly uses graphite and layered LiCoO 2as the lithium insertion compounds,which serve,respectively,as the anode and cathode materials.A lithium-containing salt such as LiPF 6dissolved in a mixture of aprotic solvents like ethylene carbonate (EC)and diethyl carbonate (DEC)is used as the electrolyte.During the charging process,the lithium ions are extracted from the layered LiCoO 2cathode,ﬂow through the electrolyte,and get inserted into the layers of the graphite anode,while the electrons ﬂow through the external circuit from the LiCoO 2cathode to the graphite anode in order to maintain charge balance.Thus,the charging process is accompanied by14Materials for High-energy Density Batteries367 Fig.14.2The charge/discharge process involved in a lithium–ion cell consisting of graphite as ananode and layered LiCoO2as a cathodean oxidation reaction(Co3+to Co4+)at the cathode and a reduction reaction at the anode.During discharge,exactly the reverse reactions occur at the anode and cathode with theﬂow of lithium ions(through the electrolyte)and electrons from the anode to the cathode.The free energy change involved in the relevant chemical reaction14.1is taken out as electrical energy during the discharge process:LiCoO2+C6→Li1−x CoO2+Li x C6(14.1)The open-circuit voltage V oc of such a lithium–ion cell shown in Fig.14.2is given by the difference in the lithium chemical potential between the cathode(μLi(c))and the anode(μLi(a))as follows:V oc=μLi(c)−μLi(a)F(14.2)where F is the Faraday constant.Figure14.3shows a schematic energy diagram of a cell at open circuit.The cell voltage V oc is determined by the energies involved in both the electron transfer and the Li+transfer.While the energy involved in electron transfer is related to the work functions of the cathode and the anode,the energy involved in Li+transfer is determined by the crystal structure and the coordination geometry of the site into/from which Li+ions are inserted/extracted(Aydinol and Ceder,1997).Although the cell voltage is a function of the separation between the redox energies of the cathode(E c)and anode(E a),thermodynamic stability consid-erations require the redox energies E c and E a to lie within the bandgap E g of the electrolyte,as shown in Fig.14.3,so that no unwanted reduction or oxidation of the electrolyte occurs during the charge–discharge process.Thus,the electrochemical stability requirement imposes a limitation on the cell voltage as follows:FV oc=μLi(c)−μLi(a)<E g(14.3)368A.ManthiramocElectrolyte/separator HOMOE E c E Fig.14.3Schematic energy diagram of a lithium cell at open circuit.HOMO and LUMO refer,respectively,to the highest occupied molecular orbital and lowest unoccupied molecular orbital in the electrolyteSeveral criteria should be satisﬁed in order for a lithium insertion compound to be successful as a cathode or an anode material in a rechargeable lithium cell.Some of the most important criteria are listed below.rThe cathode should have a low-lithium chemical potential (μLi(c))and the anode should have a high-lithium chemical potential (μLi(a))to maximize the cell volt-age (V ).This implies that the transition metal ion M n +in the lithium insertion compound Li x M y X z should have a high oxidation state to serve as a cathode and a low oxidation state to serve as an anode.rThe lithium insertion compound Li x M y X z should allow an insertion/extraction of a large amount of lithium per unit weight or per unit volume to maximize the cell capacity (A h/L or A h/kg).This depends on the number of lithium sites available in the lithium insertion/extraction host for reversible lithium inser-tion/extraction and the accessibility of multiple valences for M in the lithium insertion/extraction host.A combination of high capacity and cell voltage can maximize the energy density (W h/L or W h/kg),which is given by the product of the cell capacity and cell voltage.rThe lithium insertion compound Li x M y X z should support a reversible inser-tion/extraction of lithium with no or minimal changes in the host structure over the entire range of lithium insertion/extraction in order to provide good cycle life for the cell.This implies that the insertion compound Li x M y X z should have good structural stability without breaking any M–X bonds.r The lithium insertion compound should support both high-electronic conductiv-ity (σe )and high-lithium–ion conductivity (σLi )to facilitate fast charge/discharge (rate capability)and offer high-power capability,i.e.,it should support mixed ionic–electronic conduction.This depends on the crystal structure,arrangement of the MX n polyhedra,geometry,and interconnection of the lithium sites,nature,and electronic conﬁguration of the M n +ion,and the relative positions of the M n +and X n −energies.14Materials for High-energy Density Batteries369 r The insertion compound should be chemically and thermally stable without undergoing any reaction with the electrolyte over the entire range of lithium insertion/extraction.r The redox energies of the cathode and anode in the entire range of lithium insertion/extraction process should lie within the bandgap of the electrolyte as shown in Fig.14.3to prevent any unwanted oxidation or reduction of theelectrolyte.r The lithium insertion compound should be inexpensive,environmentally benign, and lightweight from a commercial point of view.This implies that the M n+ion should preferably be from the3d transition series.In addition to the criteria outlined earlier for the cathode and the anode materials, the electrolyte should also satisfy several criteria.The electrolyte should have high lithium–ion conductivity,but should be an electronic insulator in order to avoid internal short-circuiting.A high-ionic conductivity in the electrolyte is essential to minimize IR drop or ohmic polarization and realize a fast charge–discharge process (high rate or power capability).With a given electrolyte,the IR drop due to elec-trolyte resistance can be reduced and the rate capability can be improved by having a higher electrode interfacial area and thin separators.The electrolyte should also have good chemical and thermal stabilities without undergoing any direct reaction with the electrodes.It should act only as a medium to transport efﬁciently the Li+ ions between the two electrodes(anode and cathode).Additionally,the engineer-ing involved in cell design and fabrication plays a critical role in the overall cell performance including electrochemical utilization,energy density,power density, and cycle life(Linden,1995).For example,although ideally high-electronic con-ductivity and lithium–ion diffusion rate in the electrodes are preferred to minimize cell polarizations,the electronic conductivity of the electrodes can be improved by adding electronically conducting additives such as carbon.However,the amount of additive should be minimized to avoid any undue sacriﬁce in cell capacity and energy density.Finally,cell safety,environmental factors,and raw materials and processing/manufacturing costs are also important considerations in materials selec-tion,cell design,and cell fabrication.14.3Cathode MaterialsIntensive materials research during the past three decades has led to the identi-ﬁcation and development of several lithium insertion compounds as cathodes for lithium–ion batteries(Whittingham and Jacobson,1982;Gabano,1983;Venkatasetty, 1984;Pistoia,1994;Julien and Nazri,1994;Wakihara and Yamamoto,1998;Nazri and Pistoia,2003).Although initial efforts were focused on transition metal chalco-genides(sulﬁdes and selenides),their practical use was limited as it is difﬁcult to achieve high-cell voltages with the chalcogenide cathodes.The cell voltage is limited to<2.5V versus metallic lithium anode with the chalcogenides due to an overlap of the higher valent M n+:d band with the top of the nonmetal:p band370 A.Manthiram and the inability to stabilize higher oxidation states of the transition metal ions in chalcogenides.For example,an overlap of the M n+:3d band with the top of the S2−:3p band results in an introduction of holes into the S2−:3p band to form S22−ions rather than oxidizing the transition metal ions and accessing higher valent M n+.Recognizing this difﬁculty with chalcogenides,Goodenough’s group at the University of Oxford focused on oxide cathodes during the1980s(Mizushima et al., 1980;Goodenough et al.,1980;Thackeray et al.,1983).The location of the top of the O2−:2p band much below the top of the S2−:3p band and a larger raising of the M n+:d energies in an oxide when compared with that in a sulﬁde due to a larger Madelung energy make the higher valent states accessible in oxides.For example, while Co3+can be readily stabilized in an oxide,it is difﬁcult to stabilize Co3+in a sulﬁde since the Co2+/3+redox couple lies within the S2−:3p band.Accordingly,several transition metal oxide hosts crystallizing in different struc-tures have been identiﬁed as cathode materials during the past25years.Among them,oxides with a general formula LiMO2(M=Mn,Co,and Ni)having a layered structure,LiMn2O4having the spinel structure,and LiFePO4having the olivine structure have become appealing cathodes for lithium–ion batteries.These three systems of cathodes are discussed in the sections below.14.3.1Layered Oxide CathodesLiMO2with M=Co and Ni crystallize in a layered structure in which the Li+ and M3+ions occupy the alternate(111)planes of the rock salt structure so that the MO2sheets alternate with Li+layers along the c-axis as shown in Fig.14.4.The structure consists of a cubic close-packed oxygen array with the Li+ions occupying the octahedral interstitial sites and three MO2sheets per unit cell.Accordingly,this structure is designated as the O3layer structure in which the letter“O”refers to the presence of the alkali metal ions in the octahedral sites and the number“3”refers to the number of MO2sheets per unit cell.While the interconnected lithium–ion sites through the edge-shared LiO6octahedral arrangement between the stronglyMO2LiMO2LiMO2LiMO2Fig.14.4Crystal structure of layered LiMO2(M=Co or Ni)showing the arrangement of Li+ ions between the strongly bonded MO2sheets14Materials for High-energy Density Batteries 371bonded MO 2layers provide high lithium–ion conductivity σLi ,the edge-shared MO 6octahedral arrangement with a direct M–M interaction provides good electronic conductivity σe ,which are critical to realize fast charge–discharge rates.A large charge and ionic radii differences between the Li +and Co 3+ions offer a good ordering without any mixing between the cations in the lithium and transition metallayers.Also,a strong preference of the low spin Co 3+:3d 6(t 62g e 0g )ions for theoctahedral sites prevents the migration of the Co 3+ions from the transition metal layer to the lithium layer via the neighboring tetrahedral sites during the charge–discharge process unlike in the case of other layered oxides like LiMnO 2.On the other hand,the highly oxidized Co 3+/4+redox couple with a large work function offers a high-discharge voltage of around 4V versus Li /Li +(Mizushima et al.,1980).Moreover,with a direct Co–Co interaction and a partially ﬁlled t 2g orbital forCo 4+(t 52g e 0g )ions,Li 1−x CoO 2becomes a metallic conductor on partially charging it.As a result LiCoO 2has become an attractive cathode,and most of the lithium–ion cells presently use LiCoO 2as the cathode.However,only 50%of the theoretical capacity of LiCoO 2,which corresponds to a reversible extraction of 0.5lithium ions per LiCoO 2formula and around 140A h/kg,could be utilized in practical cells.Although early studies attributed this limi-tation in practical capacity to an ordering of lithium ions and consequent structural distortions (hexagonal to monoclinic transformation)around x =0.5in Li 1−x CoO 2(Reimers and Dahn,1992),more recent characterizations of chemically delithi-ated samples suggest that the limitation is primarily due to the chemical insta-bility at deep charge for (1−x )<0.5in Li 1−x CoO 2(Venkatraman et al.,2003;Choi et al.,2006).This is supported by the fact that the reversible capacity of LiCoO 2has been increased signiﬁcantly close to 200A h/kg with a reversible extrac-tion of 0.7lithium ions per formula unit on modifying the surface of LiCoO 2by other oxides like Al 2O 3,ZrO 2,and TiO 2(Cho et al.,2001a,b).The surface modiﬁ-cation prevents the direct contact of the highly oxidized Co 3+/4+with the electrolyte and thereby improves the chemical stability.The chemical instability of Li 1−x CoO 2at deep charge also leads to safety concerns.The chemical instability of Li 1−x CoO 2is due to a signiﬁcant overlap of the redox active Co 3+/4+:t 2g band with the top of the O 2−:2p band as shown in Fig.14.5and the consequent introduction of signiﬁcant amount of holes into the O 2−:2p band at deep charge (i.e.(1−x )<0.5).Although the analogous layered Li 1−x MnO 2and Li 1−x NiO 2offer better chemical stability than Li 1−x CoO 2as the redox active Mn 3+/4+:e g and Ni 3+/4+:e g bands lie well above the O 2−:2p band or barely touches the top of O 2−:2p band (Fig.14.5),Li 1−x MnO 2suffers from a migration of the Mn 3+ions from the transition metal plane to the lithium plane to form a spinel-like phase while Li 1−x NiO 2suffers from phase changes during the charge–discharge process.As a result,both LiMnO 2and LiNiO 2are not promising cathodes.Although both LiMnO 2and LiNiO 2by themselves could not be used as cath-odes,solid solutions among LiMnO 2,LiCoO 2,and LiNiO 2have become attractive cathodes.For example,layered LiNi 1/3Mn 1/3Co 1/3O 2has emerged as a replace-ment for LiCoO 2as it offers a lower cost and better safety when compared with372A.ManthiramLiNiO 2O 2–: 2p Ni 3+/4+: e g Ni 3+/4+: t 2gLiMnO 2O 2–: 2pMn 3+/4+: e g Mn 3+/4+–: t2g LiCoO 2Co 3+/4+: t 2gO 2–: 2p Co 3+/4+: egFig.14.5Comparison of the energy diagrams of LiCoO 2,LiNiO 2,and LiMnO 2LiCoO 2.While the lower cost arises from a lower raw material cost of Mn and Ni when compared with that of Co,the better safety is due to the lying of the Mn 3+/4+:e g and Ni 3+/4+:e g bands above the O 2−:2p band when compared with the Co 3+/4+:t 2g band as seen in Fig.14.5.The better chemical stability also allows Li 1−x Ni 1/3Mn 1/3Co 1/3O 2to be charged to higher cutoff charge voltages with a higher reversible capacity of around 180A h/kg when compared with Li 1−x CoO 2(Choi and Manthiram,2004).More recently,solid solutions between layered Li[Li 1/3Mn 2/3]O 2,which is com-monly designated as Li 2MnO 3,and layered Li[Ni 1−y −z Mn y Co z ]O 2have also become interesting as they exhibit much higher capacities of around 250A h/kg,which is nearly two times higher than that found with LiCoO 2(Lu et al.,2002;Armstrong et al.,2006;Thackeray et al.,2007;Arunkumar et al.,2007a,b).These layered solid solutions between Li[Li 1/3Mn 2/3]O 2and Li[Ni 1−y −z Mn y Co z ]O 2exhibit an initial sloping region A,which corresponds to the oxidation of the transition metal ions to 4+state,followed by a plateau region B,which corresponds to an oxidation of the O 2−ions and an irreversible loss of oxygen from the lattice,during the ﬁrst charge as seen in Fig.14.6.After the ﬁrst charge,the material cycles with a sloping discharge–charge proﬁle involving a reversible reduction–oxidation of the transition metal ions.010020030012345V o l t a g e (V )Capacity (mAh/g)B A Fig.14.6First charge–discharge proﬁles of solid solutions between layered Li[Li 1/3Mn 2/3]O 2and Li[Ni 1−y −z Mn y Co z ]O 214Materials for High-energy Density Batteries 37305010015020025030035023452345234523452345V o l t a g e (V )Specific capacity (mAh/g)x = 0.3Unmodified Al 2O 3 modified x = 0.5x = 0.4x = 0.6x = 0.7Fig.14.7First charge–discharge proﬁles of the layered (1−x )Li[Li 1/3Mn 2/3]O 2−x Li[Ni 1/3Mn 1/3Co 1/3]O 2solid solutions before and after surface modiﬁcation with Al 2O 3followed by heating at 400◦CHowever,these layered solid solution cathodes tend to exhibit a large irreversible capacity loss in the ﬁrst cycle,i.e.,a large difference between the ﬁrst charge capac-ity and the ﬁrst discharge capacity as seen in Fig.14.6.This large irreversible capac-ity loss is generally believed to be due to the extraction of lithium as “Li 2O”in the plateau region B in Fig.14.6and an elimination of the oxide ion vacancies formed to give an ideal composition “MO 2”without any oxide ion vacancies existing in the lattice at the end of ﬁrst charge,resulting in a less number of lithium sites avail-able for lithium insertion/extraction during the subsequent discharge–charge cycles (Armstrong et al.,2006;Thackeray et al.,2007).However,a careful analysis of the ﬁrst charge and discharge capacity values recently with a number of compositions suggests that part of the oxide ion vacancies should be retained in the layered lattice to account for the high-discharge capacity values observed in the ﬁrst discharge (Wu et al.,2008).The irreversible capacity loss observed during the ﬁrst cycle could also be reduced by a modiﬁcation of the cathode surface with other oxides like Al 2O 3(Wu et al.,2008;Park et al.,2001).For example,Fig.14.7shows the ﬁrst charge–discharge proﬁles of a series of solid solutions between layered Li[Li 1/3Mn 2/3]O 2and Li[Ni 1/3Mn 1/3Co 1/3]O 2before and after surface modiﬁcation with Al 2O 3,while374 A.Manthiram200250300D i s c h a r g e C a p a c i t y (m A h /g )Cycle number 200250300200250300200250300200250300Fig.14.8Cyclability of the layered (1−x )Li[Li 1/3Mn 2/3]O 2−x Li[Ni 1/3Mn 1/3Co 1/3]O 2solid solutions before and after surface modiﬁcation with Al 2O 3followed by heating at 400◦CFig.14.8shows the corresponding cyclability data.Clearly,the surface modiﬁed samples exhibit lower irreversible capacity loss and higher dis-charge capacity values than the pristine,unmodiﬁed samples.This improvement in the surface modiﬁed samples has been explained on the basis of the retention of more number of oxide ion vacancies in the layered lattice after the ﬁrst charge when compared with that in the unmodiﬁed samples.It is remarkable that the surface modiﬁed (1−x )Li[Li 1/3Mn 2/3]O 2−x Li[Ni 1/3Mn 1/3Co 1/3]O 2composition with x =0.4exhibit a high-discharge capacity of 280mA h/g,which is two times higher than that of LiCoO 2.These oxides have high potential to increase the energy density values signiﬁcantly.However,these layered oxides require charging up to about4.8V and more stable,compatible electrolyte compositions need to be developed.Moreover,oxygen is lost irreversibly from the lattice during the ﬁrst charge,and it may have to be vented appropriately during cell manufacturing before sealing the cells.Furthermore,the long-term cyclability of these cathodes under more aggres-sive environments such as elevated temperatures need to be fully assessed.14.3.2Spinel Oxide CathodesLiMn 2O 4crystallizes in the normal spinel structure in which the Li +and the Mn 3+/4+ions occupy,respectively,the 8a tetrahedral and 16d octahedral sites ofthe cubic close-packed oxygen array(Fig.14.9)to give(Li)8a[M2]16d O4.While the interconnection of the8a tetrahedral sites via the neighboring empty16c octahedral sites offers a fast3D lithium–ion diffusion(highσLi)within the covalently bonded [Mn2]O4framework,the edge shared MnO6octahedra with a direct Mn–Mn interac-tion provide good electronic conductivityσe needed for high rate capability.Unlike the layered LiMO2cathodes that could suffer from the migration of the transition metal ions from the transition metal layer to the lithium layer,the3D[Mn2]O4 spinel framework provides excellent structural stability,supporting high rate capa-bility.Additionally,the lying of the Mn3+/4+:e g band well above the O2−:2p band as seen in Fig.14.5offers excellent chemical stability unlike the Co3+/4+couple. Moreover,Mn is inexpensive and environmentally benign.As a result,LiMn2O4 with a discharge voltage of about4V(Thackeray et al.,1983)has become appeal-ing as a cathode,particularly for high-power applications such as hybrid electric vehicles(HEV).However,only about0.8lithium ions per LiMn2O4formula unit could be reversibly extracted,which limit the practical capacity to<120A h/kg.Although an additional lithium could be inserted into the empty16c sites of(Li)8a[M2]16d O4 around3V versus Li/Li+to give the lithiated spinel{Li2}16c[M2]16d O4,it is accom-panied by a transformation of the cubic(Li)8a[M2]16d O4into tetragonal{Li2}16c[M2]16d O4due to the Jahn–Teller distortion associated with the high spin Mn3+:3d4(t32g e1g)ion.This results in poor capacity retention during cycling in the3Vregion due to huge volume changes,so the3V region cannot be utilized in practical cells.In addition to the lower capacity(<120A h/kg)when compared with that of the layered oxides,the LiMn2O4spinel cathode encounters severe capacity fade even in the4V region particularly at elevated temperatures.This has been attributedFig.14.9Crystal structure of LiMn2O4spinel with LiO4tetrahedra and edge-shared MnO6 octahedrato a disproportionation of Mn 3+ions into Mn 4+and Mn 2+ions in the presence of trace amounts of hydroﬂuoric acid generated by a reaction of the LiPF 6salt with the parts per million levels of water present in the electrolyte solution;while the Mn 4+ion remains in the solid,Mn 2+leaches out into the electrolyte and poisons the anode,resulting in capacity fade.Several efforts including cationic and anionic substitutions and surface mod-iﬁcations (Park et al.,2001;Kannan and Manthiram,2002)have been pursued to overcome the capacity fade problem.Among them,optimized substitutions like Li[Mn 1.8Ni 0.1Li 0.1]O 4suppresses Mn dissolution drastically and improves the capacity retention signiﬁcantly as seen in Fig.14.10(Shin and Manthiram,2003;Choi and Manthiram,2006,2007).However,such substitutions increase the man-ganese valence and decrease the capacity to around 80A h/kg.Nevertheless,the capacity values could be increased to around 100A h/kg without increasing the manganese dissolution by a partial substitution of ﬂuorine for oxygen.For example,oxyﬂuoride compositions like Li[Mn 1.8Ni 0.1Li 0.1]O 3.8F 0.2exhibit excellent capacity retention at ambient and elevated temperatures (Fig.14.10)along with the high-rate capability (Choi and Manthiram,2006,2007).Although the capacity values708090100110120C a p a c i t y (m A h /g )010********6080100120Cycle NumberFig.14.10Comparison of the capacity retention of LiMn 2O 4(solid circles ),Li[Mn 1.8Ni 0.1Li 0.1]O 4(solid triangles ),and Li[Mn 1.8Ni 0.1Li 0.1]O 3.8F 0.2(open squares )spinel cathodes at 25and 60◦Cof the spinel cathodes are lower when compared with that of the layered oxides discussed in the previous section,the low cost of Mn and the high charge–discharge rate capabilities make the optimized oxyﬂuoride spinel cathodes appealing for high-power applications such as power tools and hybrid electric vehicles.Interestingly, the capacity values of the spinel cathodes could also be increased by blending with appropriate amounts of a layered oxide like LiNi1/3Mn1/3Co1/3O2without sacriﬁc-ing the cost and rate capability beneﬁts.One interesting way to increase the energy density of LiMn2O4spinel is to increase the operating voltage.For example,substitution of other transition metal ions like Ni2+for Mn3+/4+to give LiMn1.5Ni0.5O4has been found to increase the operating voltage to4.7V with a capacity of about130A h/kg(Zhong et al., 1997).In LiMn1.5Ni0.5O4,Mn is present as Mn4+and Ni is present as Ni2+,and the oxidation of Ni2+to Ni3+and Ni4+during the charging process and the oper-ation of the Ni2+/3+and Ni3+/4+couples along with an extraction/insertion of lithium from/into the tetrahedral sites offers a voltage of4.7V when compared with4V for the Mn3+/4+couple.It is interesting to note that while the Ni3+/4+ couple with an octahedral site lithium operates around4V in the layered oxides like LiNiO2,the same couple with a tetrahedral site lithium operates around4.7V in the spinel LiMn1.5Ni0.5O4.This is a direct reﬂection of the contributions of the energies involved in both the electron transfer(work function)and the Li+ion transfer(site energy)to the cell voltage as discussed in Section14.2.Although the synthesis of LiMn1.5Ni0.5O4encounters the formation of a small amount of NiO impurity,appro-priate cationic substitutions to give compositions like LiMn1.42Ni0.42Co0.16O4and LiMn1.5Ni0.42Zn0.08O4eliminate the impurity phase with a signiﬁcant improvement in capacity retention(Arunkumar and Manthiram,2005).However,the major issue with this class of spinel cathodes is the electrolyte stability at4.7V.Development of more stable,compatible electrolytes could make these cathodes attractive for high power applications with a signiﬁcantly increased energy density when compared with the LiMn2O4cathode.14.3.3Olivine Oxide CathodesOne disadvantage with layered oxide cathodes containing highly oxidized redox couples like Co3+/4+and Ni3+/4+is the chemical instability at deep charge and the associated safety problems.Recognizing this,oxides like Fe2(XO4)3that con-tain the polyanion(XO4)2−(X=S,Mo,and W)were initiated as lithium inser-tion/extraction hosts in the late1980s.Although the lower valent Fe2+/3+couple in a simple oxide like LiFeO2would be expected to offer a lower discharge voltage of<3V,the covalently bonded groups like(SO4)2−lower the redox energies of Fe2+/3+through inductive effect and increase the cell voltage to>3V(Manthiram and Goodenough,1987,1989).Following this,LiFePO4crystallizing in the olivine structure(Fig.14.11)and offering aﬂat discharge proﬁle around3.4V with a theo-retical capacity of around170A h/kg was identiﬁed as a cathode(Padhi et al.,1997) in late1990s.Despite a higher capacity,the lower discharge voltage when compared。

纳米材料在新能源领域的应用研究

纳米材料在新能源领域的应用研究引言：随着全球能源需求的不断增加和传统能源资源的逐渐枯竭，寻找替代能源成为了当下科技界和工业界关注的焦点。

新能源领域的发展对材料科学提出了更高的要求，纳米材料的出现为新能源的开发和应用提供了巨大的潜力。

本文将重点讨论纳米材料在新能源领域应用的研究进展，涵盖太阳能、储能和催化等几个方面。

一、太阳能领域：太阳能是最为可持续的能源之一，但效率和成本一直是太阳能发展的两大瓶颈。

纳米材料的引入为太阳能电池的效率提升和成本降低提供了新的解决方案。

近年来，利用纳米结构、纳米颗粒和纳米薄膜等纳米材料制备高效太阳能电池的研究取得了令人瞩目的成果。

例如，利用纳米材料构造的光捕获结构可以显著增加太阳能的吸收能力，提高电池的光电转换效率。

同时，纳米材料的可调控性使得制备柔性太阳能电池成为可能，进一步增加了太阳能的应用范围和有效利用率。

二、储能领域：储能技术的发展对新能源的大规模应用至关重要，其中纳米材料在电池、超级电容器和储氢材料等方面的应用展现了巨大潜力。

纳米颗粒和纳米薄膜等纳米材料能够显著增加电极的表面积，提高电极材料与电解质的接触面积，从而提高了储能设备的能量密度和功率密度。

举例来说，纳米硅材料的引入使得锂离子电池的容量得到显著提升，有望在电动汽车等领域实现更高的续航里程。

此外，纳米材料还可以构建多孔结构，增加材料的承载能力和电荷迁移速率，进一步提高储能设备的性能。

三、催化领域：催化是许多能源转化过程的关键步骤，而纳米材料的催化性能往往优于传统材料，其巨大的表面积和丰富的表面活性位点使得催化剂的性能得到极大提升。

纳米催化剂的特殊形貌和组成结构可以调控反应的活性、选择性和稳定性，从而改善能源转化过程的效率和经济性。

例如，利用具有高比表面积且可调控孔径的纳米金属催化剂，可以实现高效的氧气还原反应，为燃料电池和金属空气电池等能源设备提供持续的电力输出。

结论：纳米材料在新能源领域的研究成果取得了显著进展，为开发高效、经济、环保的新能源技术提供了重要的基础。

219432416_一种复配型电解液添加剂对锌离子电池性能的影响＊

科技与创新┃Science and Technology&Innovation ·76·2023年第12期文章编号：2095-6835（2023）12-0076-03一种复配型电解液添加剂对锌离子电池性能的影响＊黄美红，潘梦鹞，王锋，甘俊旗，黄维（广东工贸职业技术学院，广东广州510510）摘要：水系锌离子电池由于它独特的优点，如较高的理论容量（820mAh/g）、低氧化还原电位和锌金属阳极的高电位，引起了研究人员的广泛关注。

然而，轻度酸性环境中的锌枝晶生长、腐蚀及氢的释放，导致库仑效率低，电池循环性能差。

针对电池在循环过程中存在副反应和枝晶生长的问题，将一种复配型添加剂加入电解液中，探究它对锌离子电池性能的影响。

由于各种成分添加剂的协同作用，锌负极表面具有良好的平整性和结晶细致性。

含有添加剂的Zn//Zn对称电池可在2000h内提供良好的循环稳定性，且过电位较低；随着电流密度的增大，电压并没有出现大的波动，即使在10mA/cm2的大电流密度下，电解液仍然显示出相对较小的电压滞后；在1.0A/g的电流密度下，添加剂能够有效提高电池的放电比容量和库仑效率；全电池在各种倍率下稳定性良好，并且即使在5.0A/g的大电流密度下，仍然具有125mAh/g的比容量。

提出的复配型添加剂方法具有很多优点，为锌电池和其他金属电池电解液的制备提供了一条新的途径。

关键词：电解液；添加剂；锌离子电池；电化学性能中图分类号：TM912文献标志码：A DOI：10.15913/ki.kjycx.2023.12.022锂离子电池在商业领域占据主导地位，然而全球锂资源短缺，以及人们对安全问题的日益关注，限制了它们的进一步大规模发展应用。

因此，发展替代能源迫切需要成本更低、安全性更高的存储系统。

可持续水系可充电金属离子电池和非金属载体电池已迅速建立。

其中，水系锌离子电池由于它独特的优点，如较高的理论容量（820mAh/g）、低氧化还原电位和锌金属阳极的高电位，引起了研究人员的广泛关注[1]。

耦合光热发电储热-有机朗肯循环的先进绝热压缩空气储能系统热力学分析

第 12 卷第 12 期2023 年 12 月Vol.12 No.12Dec. 2023储能科学与技术Energy Storage Science and Technology耦合光热发电储热-有机朗肯循环的先进绝热压缩空气储能系统热力学分析尹航1，王强1，朱佳华2，廖志荣2，张子楠1，徐二树2，徐超2（1中国广核新能源控股有限公司，北京100160；2华北电力大学能源动力与机械工程学院，北京102206）摘要：先进绝热压缩空气储能是一种储能规模大、对环境无污染的储能方式。

为了提高储能系统效率，本工作提出了一种耦合光热发电储热-有机朗肯循环的先进绝热压缩空气储能系统(AA-CAES+CSP+ORC)。

该系统中光热发电储热用来解决先进绝热压缩空气储能系统压缩热有限的问题，而有机朗肯循环发电系统中的中低温余热发电来进一步提升储能效率。

本工作首先在Aspen Plus软件上搭建了该耦合系统的热力学仿真模型，随后本工作研究并对比两种聚光太阳能储热介质对系统性能的影响，研究结果表明，导热油和太阳盐相比，使用太阳盐为聚光太阳能储热介质的系统性能更好，储能效率达到了115.9%，往返效率达到了68.2%，㶲效率达到了76.8%，储电折合转化系数达到了92.8%，储能密度达到了5.53 kWh/m3。

此外，本研究还发现低环境温度、高空气汽轮机入口温度及高空气汽轮机入口压力有利于系统储能性能的提高。

关键词：先进绝热压缩空气储能；聚光太阳能辅热；有机朗肯循环；热力学模型；㶲分析doi： 10.19799/ki.2095-4239.2023.0548中图分类号：TK 02 文献标志码：A 文章编号：2095-4239（2023）12-3749-12 Thermodynamic analysis of an advanced adiabatic compressed-air energy storage system coupled with molten salt heat and storage-organic Rankine cycleYIN Hang1, WANG Qiang1, ZHU Jiahua2, LIAO Zhirong2, ZHANG Zinan1, XU Ershu2, XU Chao2(1CGN New Energy Holding Co., Ltd., Beijing 100160, China; 2School of Energy Power and Mechanical Engineering,North China Electric Power University, Beijing 102206, China)Abstract:Advanced adiabatic compressed-air energy storage is a method for storing energy at a large scale and with no environmental pollution. To improve its efficiency, an advanced adiabatic compressed-air energy storage system (AA-CAES+CSP+ORC) coupled with the thermal storage-organic Rankine cycle for photothermal power generation is proposed in this report. In this system, the storage of heat from photothermal power generation is used to solve the problem of limited compression heat in the AA-CAES+CSP+ORC, while the medium- and low-temperature waste heat generation in the organic Rankine cycle power收稿日期：2023-08-18；修改稿日期：2023-09-18。

纳米二氧化锰的制备及其应用研究进展

第35卷第12期无机材料学报Vol. 35No. 122020年12月Journal of Inorganic Materials Dec., 2020收稿日期: 2020-03-02; 收到修改稿日期: 2020-05-07基金项目: 国家自然科学基金(61775131, 61376009); 上海高校特聘教授(东方学者)岗位计划(2013-70)National Natural Science Foundation of China (61775131, 61376009); The Program for Professor of Special Ap-pointment (Eastern Scholar) at Shanghai Institutions of Higher Learning (2013-70)作者简介: 王金敏(1975–), 男, 教授.E-mail:*******************.cn文章编号: 1000-324X(2020)12-1307-08 DOI: 10.15541/jim20200105纳米二氧化锰的制备及其应用研究进展王金敏, 于红玉, 马董云(上海第二工业大学工学部, 环境与材料工程学院, 上海 201209)摘要: 二氧化锰作为一种重要的过渡金属氧化物, 因其储量丰富、晶型多样、性能优异而备受关注。

将二氧化锰纳米化后, 其颗粒尺寸变小、比表面积变大、材料性能优化、应用领域得以拓宽。

本文在引言部分从介绍二氧化锰的应用着手, 指出纳米化和晶型多变对二氧化锰的结构和性能有着重要的影响。

正文部分主要从纳米二氧化锰的制备方法和纳米二氧化锰的应用两个方面对近年来的研究进展进行了总结和评述。

(1)介绍了水热法、溶胶-凝胶法、化学沉淀法、固相合成法等纳米二氧化锰的制备方法, 对各种制备方法的优点与缺点以及所制备纳米二氧化锰的形貌与性能进行了总结。

锂硫电池正极材料的研究进展

锂硫电池正极材料的研究进展摘要：锂硫电池具有高达2600 Wh▪kg-1的理论比能量以及1672 mAh▪g-1的理论比容量，远大于现阶段使用的商业化二次电池，被越来越多人所关注。

本文主要介绍了锂硫电池正极材料的研究进展，从结构调控型碳/硫复合正极材料、非金属元素表面修饰碳/硫复合正极材料、非碳添加剂/硫复合正极材料这三个方面进行说明，以此来突破锂硫电池目前所存在的问题。

关键词：锂硫电池；正极材料锂硫电池由单质硫正极、电解液、隔膜和金属锂负极构成。

反应机理为电化学机理，以硫为正极反应物质，以锂为负极。

在构成锂硫电池的四个部分中，正极具有极为重要的作用。

因此，大量的研究者都希望通过对正极材料的设计来攻克锂硫电池目前主要存在的单质硫导电性差、充放电产物绝缘、中间产物具有穿梭效应与活性物质的体积膨胀等本征问题，从而实现锂硫电池的商业化生产。

依照不同的设计角度，正极材料大体分为以下几类：1.结构调控型碳/硫复合正极材料碳材料广泛存在于自然界中，具有稳定的理化性质。

碳材料的引入可以显著提高正极材料的导电性，并有效缓冲活性物质的体积膨胀，避免了充放电过程中正极电极结构的粉化与脱落。

（1）多孔碳多孔碳是以碳质材料为结构基元组成的具有多孔结构的功能材料。

根据碳材料孔径分布，可以将其分为微孔（孔径小于2 nm）、介孔（孔径在2~50 nm 之间）、大孔（孔径大于50 nm）、和分级孔（具有多种孔道结构）碳材料。

其高孔隙率和高比表面积有利于硫的储存和均匀分布，并且多孔结构对多硫化物溶解和扩散具有抑制作用，有效减缓了“穿梭效应”，提高了电池的电化学性能和稳定性[1]。

（2）分级多孔碳微孔碳具有较高的比表面积，能确保单质硫在导电骨架中的分散和接触，同时其强物理吸附能力可以有效抑制“穿梭效应”。

但是微孔难以负载大量的活性物质，限制了电池整体的能量密度。

介孔碳较微孔碳具有更高的硫负载量，并能有效地缓解充放电过程中的体积膨胀。

Review-2012-NanoEnergy-Nanostructured electrodes for high-power lithium ion batteries

journal homepage: /locate/nanoenergyAvailable online at REVIEWNanostructured electrodes for high-power lithium ion batteriesRahul Mukherjee a,Rahul Krishnan b,T oh-Ming Lu c,Nikhil Koratkar a,b,na Department of Mechanical,Aerospace and Nuclear Engineering,Rensselaer Polytechnic Institute,Troy,NY12180,USAb Department of Materials Science and Engineering,Rensselaer Polytechnic Institute,Troy,NY12180,USAc Department of Physics,Applied Physics,and Astronomy,Rensselaer Polytechnic Institute,Troy,NY12180,USAReceived21March2012;received in revised form10April2012;accepted10April2012Available online26April2012KEYWORDSLithium-ion battery;Nano;Anode;Cathode;Power density;Energy densityAbstractLithium ion batteries are popular for use in portable applications owing to their high energy density.However,with an increasing interest in plug-in hybrid electric vehicles over the past few years,stemming from an urgent need to migrate to green technologies,the focus has shifted to enhancingpower densities in Lithium ion batteries.In this review article we focus on some of the recentachievements of the academic and industrial community in boosting the power densities of Lithiumion batteries through the development of novel nanostructured anode and cathode architectures.&2012Elsevier Ltd.All rights reserved.IntroductionLithium ion batteries offer very high energy densities anddesignﬂexibilities[1],thereby making them integral inmodern day consumer devices such as cellular phones,camcorders and laptop computers.However,unlike electro-chemical capacitors,lithium ion batteries are restricted tolow achievable power densities,as is evident from theRagone plot in Fig.1[2].With the advent of electricvehicles(EV)and plug-in hybrid electric vehicles(PHEV)there has therefore been a growing need to build lithium ionbatteries that can not only provide high energy densities butalso deliver high power densities in order to be consideredas a potential replacement for conventional gasolineengines.Lithium based rechargeable batteries wereﬁrst proposedby Whittingham in1976whereby he demonstrated rapid andhighly reversible intercalation reaction between lithium(anode)and layered titanium disulﬁde(cathode)[3].Lithium anodes,however,have a tendency to form dendriticgrowths that pose serious issues regarding safety[4]andaffect the cycle ability of the batteries[5].The concept oflithium ion batteries wasﬁrst introduced by Sony in1990asthey demonstrated the capability of non-graphitisable car-bon to insert lithium[6].Since then there has been a lot ofadvancement globally in theﬁeld of lithium ion batteries.2211-2855/$-see front matter&2012Elsevier Ltd.All rights reserved./10.1016/j.nanoen.2012.04.001n Corresponding author at:Department of Mechanical,Aerospace and Nuclear Engineering,Rensselaer Polytechnic Institute,Troy,NY12180,USA.Tel.:+15182762630;fax:+15182762623.E-mail address:koratn@(N.Koratkar).Nano Energy(2012)1,518–533Most of the research has primarily concentrated on improv-ing the performance of the cathode [7,8]and anode [9,10]materials and building a stable and more efﬁcient electro-lyte [11,12].A lithium ion battery essentially comprises of three compo-nents —cathode,anode and electrolyte.Cathodes are gen-erally categorized into three types,namely (1)lithium based metal oxides [13],such as LiCoO 2,(2)transition metal phos-phates [14,15],such as Li 3V 2(PO 4)3and LiFePO 4and (3)spinels [16]such as LiMn 2O 4.Among anodes,carbon is the typical material used in lithium ion batteries [17].However ,over the past few years the interest has shifted to other host materials capable of reversible lithium insertion such as silicon,tin,aluminum and germanium.More on this shall be described in the next section.Finally ,the electrodes are immersed in electrolytes that offer high ionic conductivity .The electrolyte most commonly used is based on lithium salts in aprotic solvents [18],for example LiPF 6in ethylene carbonate and diethyl carbonate (EC:DEC).The use of aqueous electrolytes [19]and solid [20]or gel type [21]polymer electrolytes in lithium ion batteries has also been reported in the literature.In addition,lithium ion batteries contain a separator to physically isolate the anode and cathode in order to ensure safe operation,while permitting ionic transport and preventing electronic ﬂow [22,23].In a lithium ion battery,lithium ions are extracted from the cathode during charging,transported through theelectrolyte and ﬁnally inserted into the anode.During discharge,lithium ions are extracted from the anode,transported through the electrolyte and inserted back into the cathode.Performance of the anode and cathode are generally quantiﬁed in terms of capacity per unit mass or per unit area of the electrode material,where capacity is the total number of ampere-hours that can be withdrawn from a fully charged cell under speciﬁed conditions of discharge [24].The nature of insertion and extraction,the operating voltage,capacity and cycle life are generally governed by the types of cathode and anode materials,the electrolyte and the rate of charge/discharge (C-rate).Apart from these,temperature of operation and storage also plays a crucial role in lithium ion batteries and signiﬁcant deterioration in performance has been reported at elevated temperatures attributed to factors such as irreversible electrolyte break down [25]and large self discharge [26].Performance deterioration with lower temperatures has also been observed in lithium ion batteries and is mostly attributed to the electrolyte characteristics [27].Although conventional electrolytes such as LiPF 6in EC:DMC:EMC or LiPF 6in EC:DMC have freezing points in the range of À301C to À501C,there is a drop in electrolyte conductivities with a decrease in the temperature by almost one order of magnitude [28].For instance,LiPF 6in EC:DMC:EMC has a conductivity of $1.1Â104s cm À1at room temperature which drops to $1Â103s cm À1at À401C.Irreversible losses in capacities and electrolytic break down thus pose a serious threat to the commercialization of advanced lithium ion batteries,especially in the automotive industry,where the battery would ideally be expected to operate over a wide temperature range.In this review we shall however focus on advancements in developing electrodes with good rate capabilities and high capacities as a stepping stone towards successful expansion of lithium ion batteries for high power and automotive applications.Lithium ion battery anodesCarbon has typically been favoured as anode materials in Lithium ion batteries,owing to its excellent cycling ability and long cycle life [29].Graphite compounds exhibit a theoretical capacity of 372mA h g À1of carbon correspond-ing to the formation of LiC 6as shown in Eq.(1)[30]Li þþ6C þe À"LiC 6ð1ÞIn addition,the diffusivity of lithium ions in graphite is exceptionally high,thereby facilitating rapid charging and discharging.The diffusion coefﬁcient of lithium ions can be effectively estimated incorporating cyclic voltammetry technique and Randles–Sevcik equation [31–33].The magni-tude of the peak current in a cyclic voltammetry plot is related to the scan rate through the Randles–Sevcik equa-tion as shown in Eq.(2).ip ¼0:4463nFAC nFVD RT 0:5ð2ÞIn the above equation,n represents the number of charges involved in the electrochemical half-reaction,F is Faraday’s constant (96,485C mol À1),A is the area of the electrode,C is the molar concentration of Li +in the electrolyte,V istheFigure 1Ragone plot of various types of batteries and capacitors.Lithium ion batteries exhibit high energy density values but suffer from poor power densities.In the current scenario when researchers and industries are actively seeking alternate energy sources for automotive and a variety of such other applications,it becomes increasingly necessary to develop a battery that will be capable of delivering both high energy and high power densities.Reprinted with permission from Ref.[2].Copyright 2008Nature Publishing Group.Nanostructured electrodes for high-power lithium ion batteries 519voltage sweep rate (V s À1),D is the diffusion coefﬁcient of lithium ions (cm 2s À1),R is the universal gas constant (8.314J mol À1K À1)and T is the absolute temperature (K).Since lithium ion batteries are primarily operated at room temperatures,Eq.(2)can be further reduced as follows.ip ¼kn 3=2v 1=2D 1=2ACð3ÞHere,k is a constant (2.69Â105C mol À1V À1/2).From Eq.(2),it can be seen that the peak current is a linear function of square-root of the sweep rate.By plotting the peak current for different scan rates,one can obtain the slope of the linear graph.Knowing this slope,n ,A and C ,the diffusion coefﬁcient of Li +can be calculated as follows.D ¼slope kn AC 2ð4ÞIn highly oriented pyrolytic graphite,the diffusion coefﬁ-cient of Li +in the direction parallel to the graphene plane has been reported to be as high as 4.4Â10À6cm 2s À1[34].Due to the inherent high diffusivity of lithium ions in graphite anodes,rapid charge and discharge is possible,thereby providing an opportunity to improve the available power signiﬁcantly.Extensive research on carbon based anode materials lead to the development of functionalized multiwalled carbon nano-tubes deposited using layer-by-layer technique that were able to provide power density as high as 100kW kg À1and demon-strated excellent cycle ability over a thousand cycles [35].However ,such carbon nanotube anode materials could deliver a reversible capacity of only $200mA h g À1.Carbon nanoﬁ-bers have also demonstrated capabilities of high rate cycling [36].In one such study ,carbon nanoﬁbers with diameters of 30nm and 230nm,heat treated at 10001C have demon-strated impressive rate capabilities.At current densities as high as 2A g À1,the capacities obtained were still above 100mA h g À1over 30charge/discharge cycles [37].More recently ,graphene (Fig.2)has attracted attention as potential anode materials in lithium ion batteries.Y oo et al.demon-strated [38]a reversible capacity as high as 784mA h g À1,thereby suggesting a possible adsorption of lithium on both sides of the graphene sheet leading to the formation of Li 2C 6[39,40].Additionally ,graphene has an inherently high con-ductivity which improves the charge transfer signiﬁcantly ,as shown by Su et al.in their recent study [41].However ,the same group demonstrated that graphene sheets blocked fast lithium ion transport (at rates of 3C and above),thereby limiting its application to high energy lithium ion batteries.Moreover ,the mass loading with graphene is extremely low and scalability of the system becomes a challenge.Wei and co-workers proposed a sputter deposited 1.1m m thick amorphous carbon ﬁlm capable of delivering high capacities with excel-lent coulombic efﬁciencies [42].Such sputter deposited amorphous carbon has inherent advantages when it comes to applications in lithium ion batteries.First,owing to high density of the deposited ﬁlm and the subsequent reduction in porosity ,the lithium ions can intercalate only through diffusion into the carbon material.This makes it easier for removal of lithium ions during the delithiation cycle as compared to delithiation of lithium ions from pore sites,thereby improving the reversible capacities.Given the high diffusion coefﬁcient of lithium ions in carbon,at low to moderate C-rates there will be an appreciable amount of lithiation in a 1.1m m thick ﬁlm.Moreover ,randomly oriented structure in the carbon ﬁlm facilitates more storage space for the lithium ions which further improves the obtainable capacities.Second,the deposited carbon ﬁlm has a very low surface area which reduces the formation of solid electrolyte interphase (SEI)which in turn limits the irreversible capacity loss.Finally ,such physical vapor deposition processes are carried out at high vacuum conditions which diminish the hydrogen concentration in the deposited ﬁlm.The presence of hydrogen in carbon has previously shown to have had a direct effect on increasing irreversible capacity loss [43].As a result of the aforementioned characteristics of sputter deposited carbon ﬁlms,initial discharge capacities as high as 987mA h g À1were obtained with a coulombic efﬁciency of 82%,which further improved over successive cycling.However ,the cycling behavior was demonstrated at low rates of 0.2C.Higher rates would result in partial diffusion of lithium ions in the ﬁlm which would drop the obtainable capacities.Therefore in order to achieve high power densities along with high capacities,alternate materials capable of insertinglithiumFigure 2(a)SEM and (b)cross-sectional TEM images of graphene nanosheets.Graphene has shown considerable promise as potential anode materials in lithium ion batteries due to their high capacities.Further research would be necessary however ,to improve the capacities during high rate applications of graphene-based anodes.One such approach could be by building an effective graphene based composite anode,employing secondary materials capable of better lithium intercalation kinetics.Reprinted with permission from Ref.[38].Copyright 2008American Chemical Society.R.Mukherjee et al.520ions have been investigated by researchers in theﬁeld.As a replacement for carbon based anode materials,some of the potential materials include Tin,Aluminum,Germanium and Silicon either in its pure form or as composites. AluminumAluminum has been envisioned as a suitable candidate for anode materials.The Al-Li phase diagram suggests that aluminum can form three possible alloys with lithium, namely AlLi,Al2Li3and Al4Li9[44].The Al–Li alloy corre-sponds to a capacity of993mA h gÀ1while Al4Li9alloy can deliver a gravimetric capacity as high as2234mA h gÀ1, owing to its light weight.However,the use of Al anodes in high power lithium ion batteries is largely restricted owing to a very slow lithium ion diffusion in aluminum.High power batteries require rapid insertion and extraction of lithium ions into the anode material.The diffusivity of lithium ions in aluminum is$6Â10À12cm2sÀ1,which results in only partial lithiation of the aluminum alloy[45].Moreover, aluminum also experiences as much as90%expansion in its volume,associated with the formation of Li–Al alloys. This rapid and extensive volume expansion leads to cracking and pulverization of the anode.Furthermore,it leads to delamination of the anode,thereby causing a loss in electrical contact.As a result,aluminum anodes have generally shown rapid capacity fading during subsequent cycling[46,47].The aforementioned factors have therefore prevented a full scale entry of aluminum anodes in high power lithium ion batteries.TinTin,on the other hand,offers several advantages over graphite as well as aluminum as anode materials in lithium ion batteries.Tin has a theoretical charge capacity of 994mA h gÀ1corresponding to the formation of Li4.4Sn alloy as shown in Eqs.(5)and(6).Moreover,it has a higher operating voltage,thereby providing a safer operation during rapid cycling[48].LiþþSnþeÀ-LiSnð5Þ3:4LiþþLiSnþ3eÀ-Li4:4Snð6ÞDiffusivity of lithium ions in tin is also signiﬁcantly high at 5.9Â10À7cm2sÀ1which allows for rapid charge/discharge cycles[49].Tin however,suffers from poor cycle life owing to tremendous volume expansion of as much as300%during alloying/dealloying reactions[50].There has however been a considerable effort aimed at addressing the aforemen-tioned issue.Recently,a novel three dimensional porous Sn–Cu alloy anode had been suggested that demonstrated stable capacities and high rate capabilities[51,52].Copper is an inactive material in the anode and does not participate directly in the lithium intercalation reaction.However, reaction involving intermetallic Cu6Sn5forms a ductile matrix that allows large volume expansion associated with Sn–Li alloying reactions,thereby improving the cycle life of the anodes.Foam based support matrices have especially helped improve the cycle life by relieving the induced strain during lithiation and delithiation cycles as demonstrated by the performance of Sn–Co alloy anodes supported on nickel foams[53].These anodes showed appreciable stability over 60cycles and a speciﬁc capacity as high as663mA h gÀ1.Tin oxides have also been considered as attractive candidates for anodes due to their high theoretical capacity of1491mA h gÀ1,good cycle ability and high coulombic efﬁciency[54–56].Li et al.demonstrated excellent capa-cities with SnO2nanoﬁbers(Fig.3(a)),prepared using a sol-gel synthesis method,when cycled at very high C-rates[57]. These nanoﬁbers performed better than thin-ﬁlms owing to higher surface area.Although,volume expansion was still visible in the nanoﬁbers,extensive disintegration of the electrode was prevented due to the superior ability of these nanoﬁbers to withstand volume changes(Fig.3(b)).Excel-lent cycle ability due to the aforementioned attributes was thus observed even at a rate as high as58C,as the cell delivered steady capacities over500mA h gÀ1through1400 charge/discharge cycles.Hybrid Sn nanoclusters–SnO2nano-ﬁbers have further boosted the capacities at moderate charge/discharge rates(0.1A gÀ1)[58].The presence of tin induces a reversible reaction between Sn and Li2O, produced during the reduction of SnO2nanowires,to further enhance the capacity and coulombic efﬁciency.Carbon nanotube(CNT)encapsulated SnO2composite anodes have also demonstrated a promise in improving the power densities of lithium ion batteries,whereby the strong,tubular CNT structures are incorporated to buffer the volume changes during lithium insertion and extraction. This has permitted incorporation of high C-rates with impressive cycle life due to the enhanced stability of the composite anode[59,60].With the advent of graphene, there has been an interest in SnO2-graphene composites with the hope that the high surface area and superior electrical conductivity of graphene could enhance the rate capability[61,62].CNT encapsulated structures also boost the cycle life of such anodes by providing a stable solid electrolyte interphaseﬁlm.Furthermore,graphene would prevent the agglomeration of SnO2particles and mitigate the volume changes,thereby enhancing cycle life. GermaniumGermanium and silicon have both received special interest as anode materials owing to their high theoretical capacities.The theoretical capacity of germanium is 1600mA h gÀ1while that of silicon is the highest reported capacity till date at4200mA h gÀ1,more than a10-fold increase from the theoretical capacities of graphite anodes. However,both Si and Ge suffer from three major draw-backs.First,being semi-conductors,the electrical resistiv-ities of both Si and Ge are signiﬁcantly higher at103O m and1O m,respectively.The poor conductivity leads to an inefﬁcient charge transfer during charge/discharge cycles. Second,the diffusivity of lithium ions is signiﬁcantly slower in Ge and Si.Lithium ions have very slow diffusion in silicon with a diffusion coefﬁcient as low as$10À12cm2sÀ1, thereby limiting the charge/discharge rates[63].Germa-nium,on the other hand,offers a$400times increase in the diffusivity of lithium ions as compared to silicon at room temperature[64],although the diffusion of lithium ions is still signiﬁcantly slower than that in carbon and tin basedNanostructured electrodes for high-power lithium ion batteries521anodes.Finally,both silicon and germanium have tremen-dous volume expansions between 300–400%.In order to develop high power lithium ion batteries incorporating Si or Ge anodes,it is necessary to overcome the aforementioned inherent characteristics of the two materials.Doping of germanium and silicon with phosphorus or boron has been established as a successful means to over-come the conductivity limitations in these materials [65–67].Higher conductivity induced in the anodes as a result of their doping in turn ensures efﬁcient charge transport.The poor diffusivity of lithium ions in germanium and silicon can be overcome by incorporating nanostruc-tured anodes.Nanostructuring the anodes signiﬁcantly reduce the diffusion distance for lithium ions,thereby permitting complete lithiation and delithiation even at high charge/discharge rates.Nanostructured anodes are often fabricated by means of physical vapor deposition methods such as sputtering.Addition of dopants to Si and Ge,enables simple dc sputtering of the target materials,thereby avoiding the need to adopt more complex r .f.sputtering or pulsed dc sputtering techniques.Not only does this reduce the cost of deposition,it also provides a relatively simple method for fabrication of such nanostruc-tured anodes.In spite of the relatively high cost of germanium,there has been an active interest in this material as anodes,especially since their applications are restricted to nanos-tructures or thin ﬁlms and the cost incurred at such scales is signiﬁcantly low.Germanium follows a three step electro-chemical reaction during its initial charge cycle as shown in Eq.(7)below [68].Ge -Li 9Ge 4-Li 7Ge 2-Li 15Ge 4þLi 22Ge 5ð7ÞFormations of Li 15Ge 4and Li 22Ge 5alloys induce a volume expansion of $270%,thereby affecting the structural integ-rity of the anode and its cycle ability.In order to accom-modate this rapid volume change,Park et al.fabricated novel hollow 0-D and porous 3-D Ge nanoparticles [69].The 3-D Ge nanoparticles in particular showed impressive integrity after 100cycles at a rate of 1C.The excellent structural integrity also ensured that there was very little capacity fading associated with pulverization or loss of electrical contact of the germanium nanoparticles.Germanium nanocrystals and amorphous thin ﬁlms have also demonstrated superior rate capabilities [70].Germanium thin ﬁlms,in particular ,demonstrated excellent capacity retention at lithiation rates of 1C and delithiation rates as high as 1000C.Bulk germanium on the other hand is incapable of charging at such high rates due to long diffusion distances and shows a rapid degradation in capacity over as few as 10cycles.Although,extremely high charging rate was obtained in this work,it is important to demonstrate high rate capability for both charge and discharge cycles.Butyl capped amorphous germanium nanoparticles have shown impressive performance at high C-rates in this regard [71].At low rates,amorphous germanium nanoparticles displayed a near theoretical capa-city of $1470mA h g À1.Interestingly ,even at a 5C rate,the anodes still delivered capacities of $1400mA h g À1.However ,amorphous Ge nanoparticles are severely affected by volume expansion.The electrodes fail to withstand such severe stress,thereby causing rapid loss in capacity .This necessitates the need to develop stable and structurally robust germanium anodes with an improved cycle life.Rapid charging and discharging with germanium anodes over 50cycles was recently demonstrated by incorporating germanium nanotubes (Ge-NT)[72].Ge-NTs,showninFigure 3(a)SEM image of SnO 2nanoﬁbers before cycling and (b)SEM image of SnO 2nanoﬁbers after 800charge/discharge cycles at 58C.Volume expansion due to lithiation is evident from the larger ﬁbril diameter post cycling.Formation of bulb-like structures at the ﬁbril tips suggest larger volume expansion at these sites due to higher lithiation and was attributed to availability of a larger surface area for lithium insertion here.At the tip,lithium can diffuse not only through the walls but also through the tip itself in a direction parallel to the ﬁber axis.In comparison,lithium diffusion is restricted only through the walls along the remaining length of the ﬁbers.Reprinted with permission from Ref.[57].Copyright 2001The Electrochemical Society.R.Mukherjee et al.522Fig.4(a),proved to be extremely effective in accommodat-ing volume expansion associated with lithiation of germa-nium.They demonstrated impressive rate capabilities as well,as shown in Fig.4(b).Germanium nanowires (Ge-NW)which allow for extensive volume changes without struc-tural degradation and offer improved charge transport efﬁciency by providing 1-D conductivity have also shown to be effective in allowing high rate cycling [73].At a discharge rate of 2C,Ge-NWs demonstrated capacities of almost 600mA h g À1with enhanced capacity retention.The improved capacity retention may however be attributed to the fact that at such high rates,there is only partial lithiation and delithiation which causes lower volume expansion in the anode as compared to volume expansion following complete lithiation cycles.Germanium and carbon based nanocomposite materials have also managed to boost the rate capabilities of lithium ion batteries by improved structural integrity in the anodes attributed to the inclusion of carbon [74,75]as carbon is far more stable towards charge/discharge cycles,experiencing a volume expansion of only $10%.SiliconSilicon remains one of the most attractive choice for anodes in lithium ion batteries,primarily due to its exceptional high theoretical capacities of 4200mA h g À1corresponding to the formation of Li 22Si 5as shown in Eq.(8).22Li þþ5Si þ22e À-Li 22Si 5ð8ÞHowever ,there are numerous challenges confronting the commercialization of Si anodes in lithium ion batteries.First,an extremely slow diffusivity rate of lithium ions in Si has prevented its use in high C-rate applications.On the other hand,use of ultra-thin Si ﬁlms restricts the mass scalability of Si anodes [76,77].Second,Si undergoes tremendous volume changes associated with insertion and extraction of lithium ions [78–80].This leads to cracking (Fig.5(a))anddelamination (Fig.5(b))of silicon from the substrate,all contributing towards a very poor cycle life of silicon anodes [81],as shown in Fig.5(c).Third,Si anodes suffer from poor electron transfer characteristics due to their inherently low conductivities.Finally ,the cycle ability of amorphous Si is also limited by the operating voltage.In situ XRD tests have shown evidence that discharging below 30mV leads to the formation of a new Li 15Si 4crystalline phase with poor capacity retention for ﬁlms with critical thickness of about 2.5m m [82,83].In spite of these limitations,the great promise that silicon holds due to its remarkable theoretical capacity have inspired researchers in the ﬁeld to actively seek solutions to overcome these challenges.The nanostructuring of silicon electrodes has helped over-come the aforementioned limitations to a considerable extent.Short diffusion distances in nanostructured electrodes permit nearly complete lithiation even at reasonably high charge/discharge rates.High rate capabilities due to short diffusion distances have already been demonstrated in silicon thin ﬁlms,although such thin ﬁlms still present a considerable challenge with respect to commercialization due to their very low mass loading [84,85].Silicon nanospheres developed by a solvo-thermal method could deliver high capacities at a moderate cycling rate of 0.5C [86].However ,it demonstrated very low capacity retention of $35.9%of the initial capacity over repeated cycling.Kim et ter showed an improve-ment in high rate capabilities and cycle life with 3D porous carbon coated Si particles [87].Even at a rate of 3C,the capacity retention was 72%while at 1C the capacity retention was an impressive 90%(Fig.6).Cui and co-workers further demonstrated the unique capabilities of nanostructured silicon by using crystalline silicon nanowires as anodes [88].These nanowires have a three-fold advantage over other nanoscale morphologies such as ﬁlms and rods.First,the nanowires had an average diameter of $89nm that could accommodate the volume expansion without undergoing fracture.Second,the nano-wires were directly grown onto the current collecting substrate that ensures participation of all thenanowiresFigure 4(a)TEM image of Ge-NT and (b)capacity vs.cycle number as a function of C-rate.The drop in capacity associated with high C-rates is attributed to low diffusivity of lithium ions.Higher rates allow lesser time for the lithium ions to insert into and extract from the host material,resulting in partial lithiation and lower than theoretical capacities.Reprinted with permission from Ref.[72].Copyright 2011WILEY-VCH Verlag GmbH &Co.KGaA,Weinheim.Nanostructured electrodes for high-power lithium ion batteries 523。

尖晶石LiMn2O4纳米棒的合成和储锂性能

第51卷第11期2020年11月中南大学学报(自然科学版)Journal of Central South University (Science and Technology)V ol.51No.11Nov.2020尖晶石LiMn 2O 4纳米棒的合成和储锂性能周钢1，陈月皎2，范洪波1(1.东莞理工学院生态环境与建筑工程学院，广东东莞，523808；2.中南大学粉末冶金国家重点实验室，湖南长沙，410083)摘要：为了改善LiMn 2O 4材料在锂离子电池应用中倍率放电性能和循环性能较差的问题，利用纳米化尺寸改善方案，对二维LiMn 2O 4纳米棒的储锂应用进行研究。

通过水热法合成的α-MnO 2纳米线与LiOH·H 2O 进行固相烧结制备尖晶石结构的LiMn 2O 4纳米棒，探讨水热反应时间、温度和反应物物质的量比对前驱体α-MnO 2形貌结构和性能的影响。

研究结果表明：制备的LiMn 2O 4纳米棒比表面积大，有利于缩短Li +扩散距离，较大的长径比可减少循环过程中材料的团聚，从而赋予LiMn 2O 4优异的循环性能。

LiMn 2O 4纳米棒在1C 倍率充放电时初始放电容量为104.8mA·h/g ，循环150圈时依旧可以保持初始容量的84.4%，而在3C 倍率充放电时可以保持初始容量的75%，显示了较高的电化学储锂性能。

关键词：水热合成；固相烧结；锰酸锂纳米棒；锂离子电池；储锂性能中图分类号：TB321文献标志码：A文章编号：1672-7207（2020）11-3211-09Synthesis and lithium storage performance of spinelLiMn 2O 4nanorodsZHOU Gang 1,CHEN Yuejiao 2,FAN Hongbo 1(1.School of Environment and Civil Engineering,Dongguan University of Technology,Dongguan 523808,China;2.State key Laboratory of Powder Metallurgy,Central South University,Changsha 410083,China)Abstract:To improve attenuation of electrical performance for LiMn 2O 4,the application of two-dimensional LiMn 2O 4nanorods in lithium storage was studied by using the nano-size improvement scheme.LiMn 2O 4nanorods were synthesized by solid phase sintering method using LiOH·H 2O and single-crystal α-MnO 2nanowires obtained by hydrothermal method.And the effect of hydrothermal time,temperature and the ratio of reactant and salt on the morphology and structure were investigated.The results show that the prepared LiMn 2O 4nanorods have large specific surface area and large aspect ratio,which benefits the diffusion distance of Li +,and reduces the agglomeration of materials during cycling,thus endowing the superior cycling performance of LiMn 2O 4.The electrochemical test demonstrates that LiMn 2O 4nanorods can maintain 84.4%of the initial discharge capacity which is 104.8mA·h/g after 150cycles at 1C rate,and 75%of the initial discharge storage can be maintained at 3C rate,which displays a better performance of electrochemical lithium storage.DOI:10.11817/j.issn.1672-7207.2020.11.024收稿日期：2020−08−01；修回日期：2020−09−14基金项目(Foundation item)：国家自然科学基金资助项目(51901043)(Project(51901043)supported by the National Natural ScienceFoundation of China)通信作者：范洪波，教授，从事材料科学研究，E-mail ：*****************第51卷中南大学学报(自然科学版)Key words:hydrothermal synthesis;solid phase sintering;lithium manganese oxide nanorods;lithium ion battery;lithium storage performance锂离子电池具有能量密度、功率密度较高，使用寿命长以及环境友好的特点，被广泛应用于电子产品中[1−3]，而尖晶石结构的锰酸锂(LiMn 2O 4)成本低、资源丰富、环境友好，是最具潜力的锂电正极材料之一[4]。

多通道碳阴极活化过一硫酸盐降解水中有机物的性能

大连理工大学硕士学位论文摘要活化的过硫酸盐氧化，作为一种新兴的高级氧化技术，是一种矿化难降解有毒污染物的有效方法。

在众多的活化方法中，过硫酸盐通过接受电子完成的电化学活化，具有容易操控和环境友好的特点，被认为是一种有前景的活化技术。

但在电化学活化的过程中，由于静电斥力阻碍了过硫酸盐阴离子和阴极之间的接触，导致过硫酸盐低的分解率和随后低的自由基的产生量，从而使污染物的降解效果变差。

针对此问题，本文使用天然木材衍生的碳化木（CW）制备了具有多通道的流通式阴极（FTC），通过将过一硫酸盐（PMS）阴离子限制在阴极的微通道中，能够显著地强化其与阴极的碰撞与接触，提高电化学活化的效率并增强对污染物的降解。

主要的研究成果如下：（1）通过天然松木的一步碳化制备并组装了具有丰富的介孔，良好的导电性，较高的机械强度，大量有序的微通道以及对PMS有良好的电催化活性的FTC。

以苯酚为目标污染物，探究了不同的反应条件（PMS浓度、电流密度和停留时间）对FTC电活化PMS降解苯酚性能的影响。

结果表明，在苯酚进水浓度为20 mg/L, 进水TOC=18 mg/L，进水PMS浓度为6.51 mM，背景Na2SO4为0.05 M，电流密度为2.75 mA/cm2，进水pH 2.87，停留时间10 min以及常温的条件下，通过FTC电活化PMS，PMS的分解率达到了71.9%。

苯酚和TOC的去除率分别达到了97.9%和39.6%。

EPR实验结果表明，在FTC电活化PMS的过程中，产生了大量的·OH和SO4•-。

同时，自由基淬灭实验也表明，·OH和SO4•-均参与了对苯酚的降解，且·OH对降解的贡献更大。

此外，五次循环实验的结果证明了本研究组装的FTC具有很好的稳定性。

（2）通过封闭CW的微通道，获得了流过式阴极（FBC）。

在相同的优化条件下，详细对比了在FTC中和FBC上的PMS的分解、自由基的产量以及电活化PMS降解三种酚类有机物（苯酚、双酚A和4-氯苯酚）的性能。

电化学脱合金的英文

电化学脱合金的英文Electrochemical Dealloying: Principles, Applications, and Challenges.Introduction.Electrochemical dealloying is a process that involves the selective removal of one or more constituent metalsfrom a multicomponent metallic alloy by electrochemical means. This process, often referred to as "dealuminization" in the context of aluminum-based alloys, has found widespread applications in materials science, nanotechnology, and energy conversion and storage systems. The primary advantage of electrochemical dealloying lies in its ability to create nanostructured materials with unique physical and chemical properties, such as high surface area, porosity, and conductivity.Principles of Electrochemical Dealloying.The electrochemical dealloying process occurs when an alloy is immersed in an electrolyte solution and apotential is applied between the alloy and a counter-electrode. The applied potential drives the electrochemical reactions at the alloy surface, resulting in thedissolution of one or more constituent metals. The dissolution rate of each metal depends on its electrochemical properties, such as the redox potential and electrochemical activity in the given electrolyte.During the dealloying process, the alloy is typically the anode, and the counter-electrode is the cathode. The anode is connected to the positive terminal of the power source, while the cathode is connected to the negative terminal. When the potential is applied, the alloy begins to dissolve, and the dissolved metal ions migrate towards the cathode. At the cathode, the metal ions are reduced and deposited on the surface, forming a new metal layer.The rate of metal dissolution during electrochemical dealloying is controlled by several factors, including the electrolyte composition, applied potential, temperature,and alloy composition. By optimizing these parameters, researchers can precisely control the morphology, porosity, and composition of the resulting nanostructured materials.Applications of Electrochemical Dealloying.Electrochemical dealloying has found numerous applications in materials science and engineering. Some of the key applications are discussed below:1. Nanoporous Metals: Electrochemical dealloying is widely used to create nanoporous metals with high surface area and porosity. These materials exhibit unique physical and chemical properties that are beneficial in various applications, such as catalysis, sensors, and energy storage.2. Battery Materials: Nanoporous metals produced by electrochemical dealloying have been explored as anode materials for lithium-ion batteries. The high porosity and surface area of these materials enhance the lithium storage capacity and improve the battery's performance.3. Fuel Cells: Electrochemical dealloying has also been used to create nanostructured catalysts for fuel cells. These catalysts exhibit enhanced activity and durability, which are crucial for efficient fuel cell operation.4. Biomedical Applications: Nanoporous metals produced by electrochemical dealloying have potential applicationsin biomedicine, such as drug delivery, tissue engineering, and implant materials. The porous structure of these materials allows for controlled drug release and improved cell adhesion and growth.Challenges and Future Directions.Despite the significant progress made inelectrochemical dealloying, several challenges remain to be addressed. One of the primary challenges is the control of the dealloying process at the nanoscale, as it is crucialfor achieving the desired material properties. Additionally, the development of new electrolytes and optimization of dealloying parameters are ongoing research efforts.Future research in electrochemical dealloying could focus on exploring new alloy systems, optimizing the dealloying process for specific applications, and understanding the fundamental mechanisms underlying metal dissolution and nanostructure formation. Furthermore, the integration of electrochemical dealloying with other nanotechnology approaches, such as lithography and templating, could lead to the development of even more advanced materials with tailored properties.Conclusion.Electrochemical dealloying is a powerful technique for creating nanostructured materials with unique physical and chemical properties. Its applications span multiple fields, including materials science, energy conversion and storage, and biomedicine. While significant progress has been madein this field, there are still numerous challenges and opportunities for further research and development. With the advancement of nanotechnology and materials science, electrochemical dealloying holds promise for enabling thecreation of next-generation materials with improved performance and functionality.。

助凝剂聚合电解质

助凝剂聚合电解质一、引言助凝剂聚合电解质是一种常见的功能性材料，广泛应用于电化学领域，特别是在锂离子电池和超级电容器中。

它具有良好的离子导电性能和机械稳定性，可以提高能源存储设备的性能和循环寿命。

本文将对助凝剂聚合电解质进行全面详细、完整且深入的介绍。

二、助凝剂聚合电解质的定义助凝剂聚合电解质是一种由聚合物基体和助凝剂组成的复合材料。

聚合物基体通常选择具有高分子量和良好溶解度的聚合物，如聚丙烯酸酯（PBA），聚乙烯醇（PVA）等。

助凝剂则用于增加复合材料的离子导电性能，并提供机械支撑。

三、助凝剂的种类1. 离子液体离子液体是一种具有良好离子导电性能的液态盐。

它可以作为助凝剂添加到聚合物基体中，提高电解质的离子传输性能。

离子液体的种类多样，常见的有磺酸盐、磷酸盐和四氟硼酸盐等。

2. 纳米颗粒纳米颗粒是一种具有纳米尺寸的固体颗粒。

将纳米颗粒添加到聚合物基体中可以增加电解质的导电性能，并提高材料的机械强度。

常用的纳米颗粒包括二氧化硅、氧化铝和氧化锡等。

3. 导电聚合物导电聚合物是一种具有良好导电性能的聚合物材料。

将导电聚合物作为助凝剂添加到聚合物基体中，可以显著提高复合材料的离子传输性能。

常见的导电聚合物有聚苯胺、聚噻吩和聚丙烯酸银等。

四、助凝剂对复合材料性能的影响1. 离子传输性能助凝剂可以提高复合材料的离子传输性能，促进离子在材料中的扩散和迁移。

这对于电化学设备的性能至关重要，可以提高储能设备的充放电效率和循环寿命。

2. 机械稳定性助凝剂可以增加复合材料的机械强度和稳定性，提高材料的耐久性和抗变形能力。

这对于锂离子电池等需要经受长时间循环充放电的设备来说非常重要。

3. 界面相容性助凝剂可以改善复合材料与其他组件之间的界面相容性，降低界面阻抗，提高整体系统的集成度和效率。

五、应用领域助凝剂聚合电解质广泛应用于各种电化学设备中，包括锂离子电池、超级电容器、燃料电池等。

它们可以提高这些设备的能量密度、功率密度和循环寿命，推动可再生能源技术和电动交通技术的发展。

sandwich-like,st...

S andwich-Like, Stacked Ultrathin Titanate Nanosheets for Ultrafast Lithium StorageJ iehua L iu ,J unN anomaterials in architecture for green energy conversion and/or storage provide one of the most desirable approaches to alle-viate environmental and energy issues.increasing interest in developing high-power anode materials,which can match with the state-of-the-art high-power cathodematerials, for next generation high-performance rechargeableLi-ion batteries. [5–7]Titanium dioxide is regarded as one of theideal candidates for high-rate anode materials, owing not onlyto its structural characteristics and special surface activity,but also to its low cost, safety, and environmental benignity. Thelack of open channels in bulk TiOrestricts its capacity and rate capability for reversible lithiuminsertion and extraction. A reduction in the effective size andconstruction of open channels in the material are the mainstrategies currently employed to increase the rate perform-ance. [1,4]The capacity could also be improved by reducing thepath length of lithium ion migration and improving electrontransport at the surface or in the bulk of the material.With these strategies, the capacity of ultraﬁtals and nanotubes, for example, is signiﬁlower rates. However, their capacity and cycle life deterioratedramatically at higher rates.efforts have recently been made on the fabrication of anataseTiO 2nanosheets with exposed highly reactive (001) facets.These TiO 2nanosheets are shown to be an excellent host struc-ture for lithium insertion and extraction due to the presence ofexposed (001) facets and short path along the [001] direction forlithium ion diffusion.A lthough the anatase framework undergoes insigniﬁstructural distortion during lithium insertion and extraction,the rate of lithium diffusion is still limited by the narrow spaceof the host Ti–O lattice. Also, strongly caustic NaOH and cor-rosive HCl or HF are commonly used for the synthesis ofTiO 2nanomaterials. [danger in the high-temperature and high-pressure processthe as-synthesized sample consists of LTNSs with a lat-, and the magniﬁ ed image displays the layered structure with an interlamellar spacing of nm. The high-resolution (HR) TEM image (Figure 2b ) of the gray sample obtained by annealing at 350 ° C for 2 h, clearly shows that the disc-like nanostructures are formed by stacking of several ultrathin nanosheets. The single layer thickness of .4 nm is consistent with the single-unit-cell thickness along the [010] direction. Therefore, the nanosheets should be bound by (010) facets, which is also revealed by the 2D lattice fringes observed via HRTEM (Figure S1, Supporting Information). The exposed (010) facets are considered the ideal facet, possessing empty zigzag channels with large Ti–Ti distances for lithiumhe layered structure of the as-synthesized sample was also observed by powder X-ray diffraction (XRD) analysis (Figure 2c ). from the as-synthesized ned multilamellar structure with an interlayer spacing of 10.41 Å, which is consistent with the result from the TEM analysis. The peak intensity decreases sig-cantly when the annealing temperature is increased from , due to the carbonization of the intercalated organic components in the LTNSs. After calcination at 350 °C , the ordered LTNSs are superseded by the disordered ones. At phase becomes pronouncedF igure 1. T he proposed formation mechanism of CTNS with a sandwich-like multilamellar structureEM images of the as-synthesized LTNSs (a) and the CTNSs obtained after annealing at 350 ° C for 2 h (b). XRD patterns of the samplesannealed at different temperatures (100, 200, 250, 300, and 350 ° C ) (c) and elemental mapping of CTNSs (d). The insets in (a) and (b) are the mag- ed images of the corresponding samples.titanate and organic layers. This provides strong evidence for the formation of nanocarbons. Therefore, it can be concludedthat the ﬁnal structure is the CTNSs, based on the analyses from the XRD pattern, indicating partial col-lapse of LTNS to form ultrathin anatase TiO nanosheets. Further increasing the annealing temperature to 400 removal of the nanocarbon components and formation of pure anatase TiO (JCPDS no. 21 – 1272; Figure S2, Supporting Information).T he presence of C, O, and Ti was detected by elemental mapping and energy-dispersive X-ray (EDX) analysis of the sample as shown in Figure 2d . The uniformly dispersed carbon component in the CTNSs is derived by in situ carbonization of the residual organic species that stabilize the LTNSs. The corollary is also supported by thermogravimetric analysis (TGA; Figure S3, Supporting Information). The TGA curve of the LTNSs shows a total weight loss of ca. 25% which was recorded from room temperature to 500 which the organic components have been removed completely. Apparently, the sample annealed at 350 ° C still contains about 5.6%nanocarbons by weight, compared with the sample annealed at 200 carbon pillars not only strengthen the stacked ultrathin layers and prevent complete con-densation, but also offer ample space for Li ion diffusion.T he samples were further characterized by N 2adsorption–desorption isotherms and corresponding pore size distributions (Figure S4 and S5, Supporting Information). It is interesting to observe the abrupt increase of the pore fraction with size above ≈ 1 nm in the as-synthesized LTNSs. This observation might berelated to the uniform stacked structure. The surface area and total pore volume of the CTNSs are 109 m respectively. Notably, the CTNSs possess smaller total pore volume but larger pore diameter compared with LTNSs because of the carbonization of organic components. In addition, Fou-rier transform infrared (FTIR) and Raman spectra further sup-port our analysis of carbonization. The FTIR spectra (Figure S6, Supporting Information) shows the absorption bands of C–H and C–OH bonds in CTNSs sample are almost ﬂcrystallinity of the carbon in Raman spectra is indiscernible (Figure S7, Supporting Information). It may be due to the F igure 3. E lectrochemical measurements of the CTNSs. a) The ﬁ rst-cycle charge–discharge voltage proﬁ les at different current rates of 1 C, 2 C, 5 C, and 10 C. b) Representative cyclicvoltammograms at a scan rate of 1 mV s− 1 . c) Cycling performance of the CTNSs cycled at a constant current drain of 10 C and the corresponding Coulombic efﬁ ciency and d) cycling performance at different charge–discharge rates (2–50 C).igure 4.a) TiO 2nanosheets obtained in the IL solution without Li +ion;b) Charge–discharge curves of the TiO 2nanosheets annealed at 350 °C cycled at a constant current drain of 5 C.the promising use of this material in high-power lithium-ion batteries.E xperimental SectionS ynthesis of Ionic Liquid: The synthesis of IL was carried out in a 500-mL round-bottomed ﬂask, which was immersed in an ice-bath. Acetic acid (60.0 g, 1.0 mol) was added dropwise into the N,N dimethylethanolamine (98 g, 1.1 mol). After vigorous stirring for 2 h, the obtained protic IL was directly used to prepare the titanate material.P reparation of Samples: In a typical experiment, 11 g of tetrabutylC haracterization: XRD measurement was performed with a D8 diffractometer with Cu-KR radiation (was carried out with JEOL JEM-1400 and JEOL 2100F. Ndesorption isotherms were conducted at 77 K on a Micromeritics Tristar 3000 analyzer. The BET surface areas and pore-size distribution curves were calculated using adsorption data. Thermogravimetric analysis was determined using a thermal gravity analyzer (TGA) at a temperature rise rate of 10 °C min −1from room temperature to 600air ﬂ ow. For 13C and 1Hmeasurements, a JNM-ECA400 spectrometer was used at 100.5 and 400.0 MHz, respectively. FTIR spectra were recorded on a Shimazu IR Prestige-21 FT-IR Spectrometer. Raman spectra were collected on an R-3000HR spectrometer using a red LED laser (E lectrochemical Measurementsperformed using two-electrode Swagelok-type cells with lithium serving as both the counter and reference electrodes at room temperature. The working electrode was composed of 70 wt% of the active material, 20 wt% of conductivity agent (carbon black, Super-P-Li), and 10 wt% of binder (polyvinylidene diﬂ(CTNS) was about 1–2 mg on each electrode and the ﬁ20 μm in thickness. The electrolyte used was 1mixture of ethylene carbonate and diethyl carbonate. Cell assembly was carried out in an argon-ﬁ1 mV s −1) was performed using an electrochemical workstation (CHI 660C). Galvanostatic charge–discharge cycling was conducted using a battery tester (NEWAER) with a voltage window of 1–3 V at different current rates of 1 C,2 C, 5 C, 10 C, 20 C, 30 C, and 50 C, where 1 C 170 mA g −1.S upporting InformationS upporting Information is available from the Wiley Online Library or from the author.A cknowledgementsT he authors acknowledge ﬁUniversity (RG54/07) and the Ministry of Education (ARC24/07, T206B1218RS; AcRF Tier-1, RG63/08, M52120096), Singapore.[ 1]A. S. A rico ,P. B ruceN at. Mater.2005,4[ 2]A. Y amada ,H. K oizumiM. Y onemura ,T. N。

Nanostructured materials for advanced energy conversion and storage devices

ANTONINO SALVATORE ARICÒ1, PETER BRUCE2, BRUNO SCROSATI3*, JEAN-MARIE TARASCON4 AND WALTER VAN SCHALKWIJK5
1Istituto CNR-ITAE, 98126 S. Lucia, Messina, Italy 2School of Chemistry, University of St Andrews, KY16 9ST, Scotland 3Dipartimento di Chimica, Università ‘La Sapienza’, 00186 Rome, Italy 4Université de Picardie Jules Verne, LRCS; CNRS UMR-6047, 80039 Amiens, France 5EnergyPlex Corporation, 1400 SE 112th Avenue, Suite 210, Bellevue, Washington 98004, USA
LITHIUM BATTERIES
Lithium-ion batteries are one of the great successes of modern materials electrochemistry3. Their science and technology have been extensively reported in previous reviews4 and dedicated books5,6, to which the reader is referred for more details. A lithium-ion battery consists of a lithium-ion intercalation negative electrode (generally graphite), and a lithium-ion intercalation positive electrode (generally the lithium metal oxide, LiCoO2), these being separated by a lithium-ion conducting electrolyte, for example a solution of LiPF6 in ethylene carbonate-diethylcarbonate. Although such batteries are commercially successful, we are reaching the limits in performance using the

锂锰电池存在的问题

锂锰电池存在的问题1. 引言锂锰电池是一种常见的二次电池，具有高能量密度、低成本和良好的环境友好性等优点，因此被广泛应用于移动设备、电动车辆和储能系统等领域。

然而，锂锰电池在使用过程中也存在一些问题，这些问题包括容量衰减、循环寿命短、安全性差等。

本文将对这些问题进行全面详细的分析。

2. 容量衰减容量衰减是锂锰电池最常见的问题之一。

在充放电过程中，锂离子会在正负极之间来回迁移，但由于材料结构和化学反应的限制，锂离子在迁移过程中会发生损失。

随着循环次数的增加，正负极材料逐渐疲劳，导致容量衰减。

温度变化、充放电速率和深度等都会影响容量衰减。

3. 循环寿命短循环寿命是指锂锰电池能够进行多少次完整的充放电循环。

循环寿命短是锂锰电池的另一个问题。

循环寿命受到多种因素的影响，包括锂离子迁移、电解液的降解、正负极材料的脱附等。

这些因素导致电池内部结构和化学反应发生变化，从而影响循环寿命。

4. 安全性差锂锰电池存在一定的安全隐患。

锂金属在充放电过程中容易形成锂枝晶，导致短路和过热等问题。

由于正极材料中含有锰，当电池过充或过放时，会引发氧气释放和正极材料的结构破坏，进而引发热失控和爆炸等严重安全事故。

温度升高也会加剧锂离子迁移速率，增加了安全风险。

5. 其他问题除了上述问题外，锂锰电池还存在其他一些问题。

在高温环境下使用锂锰电池容易引起自发燃烧；充放电速率过快会导致电池内部温度升高，进而影响电池的性能和寿命；锂锰电池的能量密度相对较低，无法满足某些高能量需求的应用等。

6. 解决方案针对锂锰电池存在的问题，可以采取一些解决方案来改善其性能和安全性。

可以通过优化电极材料和结构设计来提高容量衰减和循环寿命。

引入新型材料、改变电极结构和增加表面积等措施可以减缓容量衰减和延长循环寿命。

可以通过优化电解液组分和添加剂来提高安全性。

添加抑制锂枝晶生长的添加剂、控制过充过放等措施可以改善安全性能。

还可以通过改进制造工艺和严格质量控制来提高产品的一致性和可靠性。

高氮掺杂的手风琴状碳用于快速稳定的储钾

高氮掺杂的手风琴状碳用于快速稳定的储钾=钾离子电池(PIB)是未来大规模储能的潜在候选者。

一个关键的挑战是石墨碳负极的(脱)钾稳定性受到有限的(002)层间距的阻碍。

具有层次结构的非晶碳可以缓冲重复(脱)钾过程中的体积变化病实现稳定循环。

近日，广东工业大学张文礼教授等研究人员采用直接热解的方法合成了高氮掺杂(26.7 at.%)的类手风琴碳负极，该负极由薄碳纳米片和涡轮层状晶体结构组成。

手风琴状碳的层次结构是由于热解炭化过程中的自组装过程而形成的。

层次化的氮掺杂手风琴结构使其具有346 mAh g−1的高可逆容量和优越的循环稳定性。

相关研究结果以“Accordion-Like Carbon with High Nitrogen Doping for Fast and Stable K Ion Storage”为题，发表在Advanced Energy Materials 上。

本工作为利用直接热解工艺设计具有层次结构和富缺陷晶状结构的碳负极开辟了新的途径。

图文解析：图1. NvGn 的制备过程及其微观结构。

a) 用于制备NvGn 的RRS 流程图，以及在该策略的每个阶段获得的产品结构示意图。

b) 表面和 c,d) 典型 3DG 横截面的 SEM 图像。

e) PGPs 的 SEM 图像和 f) TEM 图像。

g) 高倍率下 PGP 的 TEM 图像，其中 PGP 上的孔显示在黄色虚线圆圈中。

h,i) NvGII 的 SEM 图像。

j) NvGII 微观结构细节的SEM 图像。

k,l) NvGII 横截面的 SEM 图像。

NvGII 中石墨烯层状结构的 m) 表面和 n) 层到层的微观结构。

o,p) NvGII 的 TEM 图像。

q) 典型的 3DG 仅作为人体骨骼，而r) NvGn (n ≥ 2) 就像肌肉和韧带相连的人体骨骼。

图 2. NvGII 的表征。

a) NvGII 可以制造成所需的确切形状和尺寸。

High capacity and stable cathode materials

专利名称：High capacity and stable cathode materials发明人：Jianming Zheng,Jiguang Zhang,PengfeiYan,Chongmin Wang,Wengao Zhao,ShuruChen,Wu Xu申请号：US15597025申请日：20170516公开号：US10243206B2公开日：20190326专利内容由知识产权出版社提供专利附图：摘要：High energy density cathode materials, such as LiNiMnCoO2 (NMC) cathode materials, with improved discharge capacity (hence energy density) and enhanced cyclelife are described. A solid electrolyte, such as lithium phosphate infused inside of secondary particles of the cathode material demonstrates significantly enhanced structural integrity without significant or without any observable particle cracking occurring during charge/discharge processes, showing high capacity retention of more than 90% after 200 cycles at room temperature. In certain embodiments the disclosed cathode materials (and cathodes made therefrom) are formed using nickel-rich NMC and/or are used in a battery system with a non-aqueous dual-Li salt electrolytes.申请人：Battelle Memorial Institute地址：Richland WA US国籍：US代理机构：Klarquist Sparkman, LLP更多信息请下载全文后查看。

FeS2钠离子电池正极材料的研究进展

FeS2钠离子电池正极材料的研究进展孙瑞军;董亚浩【摘要】由于黄铁矿(FeS2)具有资源丰富、廉价、无毒及理论容量高等优点,使其成为一种极具前景的钠离子电池的正极材料,引起了人们的广泛关注和研究.本文总结了近年来FeS2作为钠离子电池电极材料的研究进展,就其储钠机理、微观结构的设计与合成、电解质体系的选择及它们对储钠性能的影响等进行了介绍,并对其存在的问题和挑战及解决思路进行了探讨.%Owing to the advantages of abundance in the earth crust, low cost, nontoxicity and high theoretical specific capacity, FeS2 has received great attention as cathode materials in rechargeable batteries.The recent progress of FeS2 as cathode materials for sodium-ion batteries, including storage mechanism, materials design and synthesis, optimization of electrolyte systems and electrochemical performance was summarized.The challenges and solutions were also discussed.【期刊名称】《广州化工》【年(卷),期】2017(045)006【总页数】5页(P5-8,37)【关键词】FeS2;钠离子电池;正极材料;储钠机理【作者】孙瑞军;董亚浩【作者单位】兰州交通大学,光电技术与智能控制教育部重点实验室,甘肃兰州730070;天津大学理学院化学系,天津 300354【正文语种】中文【中图分类】O61.481+1钠离子电池(Sodium-ion Batteries)和锂离子电池(Lithium-ion Batteries)的研究都起始于上世纪中后期，但是随着锂离子电池的成功应用，钠离子电池的研究一度停滞[1-2]。

有机材料能量密度

有机材料能量密度概述能量密度是衡量材料储存和释放能量能力的重要指标。

有机材料是指由碳元素构成的化合物，它们通常具有较低的能量密度。

然而，随着科学技术的不断发展，人们对于提高有机材料的能量密度进行了广泛研究。

本文将探讨有机材料能量密度的相关概念、影响因素以及目前取得的进展。

什么是能量密度？能量密度是指单位体积或单位质量中所含有的储存或释放的能量数量。

在化学领域中，常用单位为焦耳/升（J/L）或焦耳/克（J/g）。

较高的能量密度意味着更多的储存或释放能量。

有机材料的特点有机材料主要由碳元素构成，具有许多独特的特点：1.轻质：与金属和无机材料相比，有机材料通常具有较低的密度，因此在相同质量下其体积较大。

2.易加工：由于其分子结构相对简单，有机材料易于合成和加工，可以通过调整分子结构来改变其性质。

3.稳定性差：有机材料通常不具备较高的熔点和热稳定性，容易受到高温、氧气和光照等外界条件的影响而发生分解。

影响有机材料能量密度的因素提高有机材料的能量密度是一个复杂而具有挑战性的任务。

以下是影响有机材料能量密度的关键因素：1.分子结构：有机材料的分子结构直接影响其储存和释放能量的能力。

一些特定的化学键或官能团可以提供更高的能量密度。

2.化学反应：有机材料在储存和释放能量过程中通常会经历化学反应。

通过优化反应条件和选择适当的反应物，可以提高能量密度。

3.能源转换效率：有机材料作为电池、超级电容器等能源储存设备中的活性物质，其转换效率对于提高能量密度至关重要。

目前取得的进展近年来，科学家们在提高有机材料能量密度方面取得了一些重要进展。

以下是一些具有代表性的研究成果：1.高能量密度聚合物：研究人员通过合成新型聚合物材料，成功实现了较高能量密度的有机材料。

这些聚合物通常具有特殊的结构，可以在储存和释放能量过程中发生可逆的氧化还原反应。

2.纳米复合材料：将纳米颗粒引入有机材料中，可以增强其电导性和储能性能，从而提高能量密度。

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Nanostructured high-energy cathode materials for advanced lithium batteriesYang-Kook Sun 1,2*†,Zonghai Chen 3†,Hyung-Joo Noh 1,Dong-Ju Lee 1,Hun-Gi Jung 1,Yang Ren 4,Steve Wang 4,Chong Seung Yoon 5,Seung-Taek Myung 6and Khalil Amine 3*Nickel-rich layered lithium transition-metal oxides,LiNi 1−x M x O 2(M =transition metal),have been under intense investigation as high-energy cathode materials for rechargeable lithium batteries because of their high speciﬁc capacity and relatively low cost 1–3.However,the commercial deployment of nickel-rich oxides has been severely hindered by their intrinsic poor thermal stability at the fully charged state and insufﬁcient cycle life,especially at elevated temperatures 1–6.Here,we report a nickel-rich lithium transition-metal oxide with a very high capacity (215mA h g −1),where the nickel concentration decreases linearly whereas the manganese concentration increases linearly from the centre to the outer layer of each ing this nano-functional full-gradient approach,we are able to harness the high energy density of the nickel-rich core and the high thermal stability and long life of the manganese-rich outer layers.Moreover,the micrometre-size secondary particles of this cathode material are composed of aligned needle-like nanosize primary particles,resulting in a high rate capability.The experimental results suggest that this nano-functional full-gradient cathode material is promising for applications that require high energy,long calendar life and excellent abuse tolerance such as electric vehicles.In the past decade,major efforts have been devoted to search-ing for high-capacity cathode materials based on LiNi 1−x M x O 2,mostly on account of their very high practical capacities (220–230mA h g −1)at high voltages (4.4–4.6V).However,at such high operating voltages,these materials react aggressively with the elec-trolyte owing to the instability of tetravalent nickel in the charged state,leading to very poor cycle and calendar life.Therefore,these materials operate reversibly only at a potential range below 4V,resulting in low capacities of 150mA h g −1.To improve the stability of these materials,several researchers have investigated the effect of Mn substitution on cycle and calendar life.The introduction of Mn to the transition-metal layer can help stabilize the transition-metal oxide framework,because part of the Mn does not change valence state during charge and discharge 7–9.Recently,we reported several approaches to improve both the life and safety of nickel-rich cath-ode materials for potential use in plug-in hybrid electric vehicles 10.For instance,a core–shell approach 11resulted in a nickel-rich LiNi 0.8Co 0.1Mn 0.1O 2core that delivered high capacity at high volt-age,and a manganese-rich LiNi 0.5Mn 0.5O 2shell that stabilized the surface of the material.However,owing to the structural mismatch and the difference in volume change between the core and the shell,a large void forms at the core/shell interface after long-term cycling,1Departmentof WCU Energy Engineering,Hanyang University,Seoul 133-791,South Korea,2Department of Chemical Engineering,Hanyang University,Seoul 133-791,South Korea,3Chemical Sciences and Engineering Division,Argonne National Laboratory,9700South Cass Avenue,Lemont,Illinois 60439,USA,4Advanced Photon Source,Argonne National Laboratory,9700South Cass Avenue,Lemont,Illinois 60439,USA,5Department of Materials Science and Engineering,Hanyang University,Seoul 133-791,South Korea,6Department and Institute of Nano Engineering,Sejong University,Seoul 143-747,South Korea.†These authors contributed equally to this work.*e-mail:yksun@hanyang.ac.kr;amine@.leading to a sudden drop in capacity 12,13.We also demonstrated that this structural mismatch could be mitigated by nano-engineering of the core–shell material,where the shell exhibits a concentration gradient 14–16.However,because of the short shell thickness,the manganese concentration at the outer layer of the particle is low;therefore,its effectiveness in stabilizing the surface of the material is weak,especially during high-temperature cycling (55◦C).The nickel-rich lithium transition-metal oxide investigated here has a nominal composition of LiNi 0.75Co 0.10Mn 0.15O 2,and the concentration gradient of transition metals shown in Fig.1;the con-centration of nickel decreases gradually from the centre towards the outer layer of the particle,whereas the concentration of manganese increases gradually so that the manganese-rich and nickel-poor outer layer can stabilize the material,especially during high-voltage cycling.The full concentration gradient (FCG)cathode material was prepared by a newly developed co-precipitation method involving the precipitation of transition-metal hydroxides from the precursor solutions,where the concentration ratio of Ni/Mn/Co changes continuously with the reaction time (see Methods).Figure 2shows scanning electron microscopy (SEM)images and the elemental distribution of Ni,Co and Mn within a single particle of both the precursor ((Ni 0.75Co 0.10Mn 0.15)(OH)2)and the final lithiated product (LiNi 0.75Co 0.10Mn 0.15O 2)having a concentration gradient.The atomic ratio between Ni,Co and Mn was determined by integrated two-dimensional (2D)electron probe micro-analysis (EPMA).Figure 2clearly demonstrates that the atomic percentage of Co remained constant at about 10%in both the precursor and the lithiated particles as originally designed,whereas the concentration of Ni decreased and Mn increased continuously from the centre towards the outer layer of the particle.Note that the slopes representing the metal (Ni and Mn)concentration change of the precursor are greater than those of the lithiated material because of the directional migration of the metal elements during the high-temperature calcination to increase the entropy.Hard X-ray nanotomography was used to determine the 3D distribution of Ni in a single lithiated particle.Similar to medical computerized tomography,this technique uses X-rays to obtain a 3D structure at up to 20nm resolution.Figure 3a shows the 3D volume rendering of a particle acquired with the technique.The data are imaged with the particle’s volume partially removed to reveal the central cross-section.With this 3D image,we are able to illustrate the concentration profile of Ni at any given plane (see Fig.3b for a view of a plane going through the centre of the particle).Figure 3a,b shows that the structure of the centre,with a diameter ofInsideNi-rich composition: A t o mi c r a t i o (%)N i C o M nFull gradient compositionGradual Ni decrease and Mn increase from inner part to outer part of a particle D i st a nc e (µm )SurfaceFigure 1|Schematic diagram of the FCG lithium transition-metal oxide particle with the nickel concentration decreasing from the centre towards the outer layer and the concentration of manganese increasing accordingly.A SEM image of a typical particle is shown in Fig.2b.bcLithiated2 µm80604020A t o m i c r a t i o (%)100Distance from particle centre (µm)Distance from particle centre (µm)aNiFigure 2|SEM and EPMA results.a ,b ,SEM mapping photograph of Ni,Co and Mn within a single particle for the precursor (a )and for the lithiatedmaterial (b ).c ,d ,EPMA line scan of the integrated atomic ratio of transition metals as a function of the distance from the particle centre to the surface for the precursor (c )and the lithiated material (d ).The Ni-rich particle centre and Mn-rich outer surface are clearly seen from the SEM mapping images.The Ni concentration decreased linearly towards the particle surface for both the precursor and the lithiated particle whereas the Mn concentration increased,and the Co concentration remained constant.about 2µm,was markedly different from that of the outer layer.The central core is mainly composed of bright islands with numerous voids represented by the dark background.This structure could result from the different formation kinetics of the transition-metal hydroxide seed at the beginning of the co-precipitation process,during which the seed developed along with many stacking voids.After the initial formation process,the growth of the particles reached a relatively steady state,and a denser layer developed above the loosely stacked core.Figure 3b also confirms the results of EPMA (see Fig.2b).A striking feature in Fig.3b is that the bright area with high nickel content tends to form needle-shaped spikes pointing from the centre towards the edge;this feature wasa cb d12:34:00Figure3|Hard X-ray nanotomography and transmission electron microscopy images.a,X-ray nano-computed tomography image of the3D distribution of nickel concentration in a single lithiated FCG lithium transition-metal oxide particle.b,2D distribution of nickel on a plane going through the centre of the particle.The high nickel content regions shown as bright areas tend to form needle-shaped spikes radiating outwards from a∼2µm central core.c,TEM image of the local structural feature near the edge of the particle showing highly aligned nanorods.d,TEM image of the local structural feature at the centre of the particle showing that an aligned nanorod network at the particle centre was not developed.clearly captured in the transmission electron microscopy image ashighly aligned large-aspect-ratio nanorods(Fig.3c).We believe thatthis nano-pattern was the result of the directional transition-metalmigration during the high-temperature calcination that led to thereduction in the slope of the concentration gradient(Fig.2c versusd).Owing to the pre-conditioned concentration gradient in theparticles of the precursor,the crystal growth during the high-temperature calcination was energetically preferred to forming ahighly percolated nanorod network,which minimizes the diffusionlength17between the centre and the edge for transition-metal ions.In the particle centre,the migration of the transition metal hadto rely on the limited contact of loosely packed primary particles,and the development of an aligned nanorod network was limited,as shown in Fig.3d.A potential benefit of forming such a percolated aligned nanorodnetwork is that it also provides a shorter pathway for lithium-iondiffusion during normal charge/discharge cycling at ambienttemperatures,leading to a better rate capability.Figure4a showsthe rate capability of the FCG material along with the innercomposition(IC,LiNi0.86Co0.10Mn0.04O2)and outer composition(OC,LiNi0.70Co0.10Mn0.20O2)materials,both of which were synthe-sized by the conventional constant concentration approach.Whendischarged at the C/5rate,the IC material delivered a reversiblecapacity of210.5mA h g−1;the OC material,188.7mA h g−1;andthe FCG material,197.4mA h g−1;these results are as expectedbecause the IC material has the highest nickel content andthe OC material has the lowest nickel content.However,whendischarged at the5C rate,the FCG material delivered the highestreversible capacity.As a cross-validation,the diffusion coefficientof lithium ions in the three materials was measured by thegalvanostatic intermittent titration technique18.The results showedthat the FCG material has,in general,the highest lithium-iondiffusion coefficient(Supplementary Fig.S1).Meanwhile,theelectronic conductivity was measured to be the highest for theIC material,1.67×10−4S cm−1,followed by the FCG material,3.10×10−5S cm−1,and the OC material,7.30×10−6S cm−1.There-fore,we believe that the high rate capability of the FCG material hasno strong correlation with the electronic conductivity,but mostlyoriginated from the special percolated aligned nanorod networkthat shortened the diffusion pathway of lithium ions in the particle.Figure4b shows the initial charge and discharge curve ofcoin-type half-cells based on IC,FCG and OC materials.Boththe IC material(highest nickel content)and the FCG materialdelivered a higher capacity of220.7and215.4mAh g−1,respectively,whereas the OC material(lowest nickel content)showed a lowercapacity of202mAh g−1.Note that the Coloumbic efficiency of theFCG was higher(94.8%)when compared with both the IC andOC electrodes(91%)owing to a well-developed aligned nanorodnetwork in the FCG material that facilitates Li+diffusion,and thushigh lithium utilization.Another important observation is that the reversible capacity ofthe IC material decreased markedly with cycling(Fig.4c).This rapidacD i s c h a r g e c a p a c i t y (m A h g ¬1)D i s c h a r g e c a p a c i t y (m A h g ¬1)Number of cycle Number of cycleFigure 4|Charge–discharge characteristics of IC,OC and FCG materials.a ,Comparison of rate capabilities of the FCG with the IC and OC materials (upper cutoff voltage of 4.3V versus Li +/Li).b ,Initial charge–discharge curves.c ,Cycling performance of half-cells using the FCG,IC and OC materials cycled between 2.7and 4.5V versus Li +/Li using a constant current of C/5(about 44mA g −1).d ,Discharge capacity of MCMB/FCG cathode full-cells at room and high temperature.The electrolyte used was 1.2LiPF 6in EC/EMC (3:7by volume)with 1wt%vinylene carbonate as an electrolyte additive.The cells were characterized between 3.0and 4.2V with a constant current of 1C.capacity fade was mainly caused by the direct exposure of a high content of Ni(IV)-based compound to non-aqueous electrolyte at a high potential;this exposure led to the chemical decomposition of both the surface of the electrode material and the electrolyte.In contrast,the OC material had a higher manganese content and a lower oxidizing capability towards non-aqueous electrolyte.Therefore,this material had a lower reversible capacity,but a much better capacity retention.Figure 4c shows that the FCG material had combined advantages of a high capacity from the high nickel content in the bulk and a high electrochemical stability from the high manganese content on the surface.More detailed investigation by varying the upper cutoff voltage of the test cells consistently led to the same conclusion (Supplementary Fig.S2).We assembled a pouch cell using the FCG material as the cathode and mesocarbon microbeads (MCMB,graphite)as the anode.This cell was cycled between 3.0and 4.2V with a constant current of 1C (33mA).The full-cell showed an outstanding capacity retention after 1,000cycles both at room and high temperature (Fig.4d).We also fabricated pouch-type full-cells;the cells were cycled to 4.3,4.4and 4.5V at 1C rate.In all cases,the cells exhibited excellent cycling performance (Supplementary Fig.S3).The capacity of the cells increased with cutoff voltage owing to the higher lithium utilization at high voltage.The cell cycled to 4.5V showed very minor capacity fade at 55◦C possibly caused by a limited reactivity between the charged cathode and the electrolyte.To investigate the safety of the FCG approach,we developed an in situ high-energy X-ray diffraction (HEXRD)technique 19,20and used it to study the thermal decomposition of delithiated cathode materials in the presence of the electrolyte.The delithiated cathode material was recovered from the charged cell at 4.3V and mixed with an equivalent amount of non-aqueous electrolyte,and the mixture was placed in a stainless-steel high-pressure vessel for differential scanning calorimetry (DSC).The sample was then heated from room temperature to 375◦C with a heating rate of10◦C min −1.During the thermal ramping,a high-energy X-ray beam (∼0.1Å),which is able to penetrate through a 4-mm-thick stainless-steel block,was deployed to continuously monitor the structural change of the delithiated material.High-quality X-ray diffraction data were collected at an interval of 20s per spectrum (see Supplementary Fig.S4for the full index of the layered material).Figure 5a,b shows zoomed (2.40◦–2.80◦)contour plots of the in situ HEXRD profiles of delithiated IC and FCG materials during thermal ramping.In these profiles,red represents a high intensity;blue,a low intensity.The complete spectra of the in situ data can be seen in Supplementary Figs S5and S6.Three diffraction peaks can be seen in Fig.5a,b;the left one starting at 2.57◦is the (101)peak for layered transition-metal oxides,and the right weak one starting at 2.68◦is the (012)peak for layered oxides.The one in the middle (starting at 2.62◦)is the diffraction peak from the DSC vessel and can be used as a semi-reference.Figure 5a shows that the delithiated IC material (Li 1−x Ni 0.86Co 0.10Mn 0.04O 2)starts converting to a new phase at around 100◦C;the (101)and (012)peaks shift towards a smaller angle (more details can be seen in Supplementary Fig.S5).Figure 5b shows that the low-temperature phase transformation occurred at about 140◦C for the FCG material.Figure 5c shows the DSC profiles of delithiated cathodes in the presence of non-aqueous electrolyte.No heat flow was detected with DSC within the temperature range between 100◦C and 150◦C for both samples.Thus,this phase transformation is not related to the safety of the cathode materials,but can cause the degradation of the electrochemical performance of high-nickel-content cathode materials.It was previously reported that Li(Ni 0.9Co 0.1)O 2markedly loses its reversible capacity when aged at 90◦C (ref.21).Therefore,we believe that the better capacity retention of the FCG material (as shown in Fig.4b)can be attributed to the suppressed kinetics of detrimental phase transformation at temperatures around 100◦C.The in situ HEXRD data also showed that the newly formed phase started to disappear at about 200◦C for the delithiated IC material,T e m p e r a t u r e (°C )150350300250200100502.752.702.652.602.552.502.452 2.802.40T e m p e r a t u r e (°C )150350300250200100502.752.702.652.602.552.502.452 2.802.40(012)(101)H e a t f l o w (W g ¬1)024********Temperature (°C)cabθθFigure 5|Contour plots of in situ HEXRD proﬁle.a ,Delithiated IC material.b ,Delithiated FCG material during thermal ramping from room temperature to 375◦C with a scanning rate of 10◦C min −1.c ,DSC proﬁles of the delithiated FCG material,the delithiated IC and the delithiated OC(Li 1−x Ni 0.70Co 0.10Mn 0.20O 2)with a scanning rate of 1◦C min −1.The cells were constant-voltage charged to 4.3V versus Li +/Li before disassembling.and the DSC data indicated a significant exothermal reaction at about 210◦C (Fig.5c).The corresponding phase transformation for the FCG material was much slower:the new phase disappeared at about 250◦C,and the DSC data showed an exothermal reaction starting at about 250◦C (Fig.5c).Thus,the FCG material shows better safety characteristics than the IC material by shifting its exothermal reaction to a higher temperature.We have developed a high-performance cathode material composed of lithium transition-metal oxide with FCG within each particle.The structure takes advantage of the high capacity from nickel-rich materials,the high thermal stability of manganese-richmaterials and the high rate capability of highly percolated and aligned nanorod morphology.This newly developed material can deliver a specific capacity of up to 215mA h g −1with outstanding cycling stability in a full-cell configuration,maintaining 90%capacity retention after 1,000cycles.This material based on the full gradient approach can lead to the rational design and development of a wide range of functional cathodes with better rate capability,higher energy density and better safety characteristics.MethodsSynthesis of Li(Ni 0.86Co 0.10Mn 0.04)O 2and Li(Ni 0.70Co 0.10Mn 0.20)O 2.To synthesize spherical constant-concentration layered oxide cathodes,NiSO 4·6H 2O ,CoSO 4·7H 2O and MnSO 4·5H 2O (0.86:0.1:0.04,molar ratio for Li(Ni 0.86Co 0.10Mn 0.04)O 2and 0.70:0.10:0.20molar ratio forLi(Ni 0.72Co 0.10Mn 0.18)O 2)were used as the starting materials for the co-precipitation process 22.The obtained spherical precursors were mixed with LiOH ·H 2O (Li /(Ni +Co +Mn)=1molar ratio)and calcined at 750◦C for 20h in air.Synthesis of FCG material.To prepare the FCG cathode material,NiSO 4·6H 2O ,CoSO 4·7H 2O and MnSO 4·5H 2O (0.9:0.1:0.0molar ratio)were used as the starting materials for the co-precipitation process.During the reaction,a manganese-rich aqueous solution (Ni/Co/Mn,0.64:10:26molar ratio)was continuously pumped into the stock solution tank containing the starting nickel-rich solution,after which the homogeneously mixed solution was continuously fed into a continuously stirred tank reactor.The obtained FCG hydroxide was mixed with LiOH ·H 2O (Li /(Ni +Co +Mn),1molar ratio)and calcined at 750◦C for 20h in air.Morphology characterization.The morphology of the prepared powders was characterized by SEM (S4800,HITACHI)and transmission electron microscopy (2010,JEOL).Element mapping was carried out with an electron-probe micro-analyser (JXA-8100,JEOL).Chemical composition characterization.The chemical composition of the resulting powders was analysed by atomic absorption spectroscopy (Vario 6,Analyticjena).Hard X-ray 3D nanotomography.This analysis was carried out at the beamline 32-ID of the Advanced Photon Source,Argonne National Laboratory.The instrument uses a Fresnel zone plate lens to magnify X-ray images to achieve up to 20nm resolution.The X-ray energy can be continuously tuned between 8and 30keV with 0.01%energy resolution.A spherical particle with a diameter of about 6µm was selected and mounted on the tip of a sharp sample pin.A series of 2D X-ray images was collected while the particle was rotated by 180◦.A differential absorption contrast technique was used to map the Ni concentration in three dimensions.Two data sets were acquired,above the Ni K-edge (8,350eV)and below the edge (8,320eV).These data sets were reconstructed by computed tomography techniques to produce the 3D volume data 20,and the image intensity changes within each voxel can be used to calculate the Ni concentration.Electrochemical test.For fabrication of the cathodes,the prepared powders were mixed with carbon black and polyvinylidene fluoride (80:10:10)in N -methylpyrrolidinon.The obtained slurry was coated onto Al foil androll-pressed.The electrodes were dried overnight at 120◦C in a vacuum before use.Preliminary cell tests were done with a 2032coin-type cell using Li metal as the anode.The cycle-life tests were performed in a laminated-type full-cell wrapped with an Al pouch.MCMB graphite (Osaka Gas)was used as the anode.The electrolyte solution was 1.2M LiPF 6in ethylene carbonate-ethyl methyl carbonate (3:7in volume,PANAX ETEC).The cells were cycled between 3and 4.2V at a very low rate of 0.1–0.5C (0.33–16.5mA)during the initial formation process.The cells were charged and discharged between 3.0and 4.2V by applying a constant 1C current (1C corresponds to 33mA)at 25◦C.Electric conductivity measurement.The d.c.electrical conductivity was measured by a direct volt–ampere method (CMT-SR1000,AIT),in which disc samples were contacted with a four-point probe.DSC.For the DSC experiments,the cells containing the cathode materials were constant-voltage charged to 4.3V versus Li,and disassembled in an Ar-filled dry box.A 30-µl high-pressure stainless-steel DSC vessel with a gold-plated copper seal was used to host 3–5mg samples,including solids and electrolyte.The measurements were carried out in a Pyris 1differential scanning calorimeter (Perkin Elmer)using a scanning rate of 1◦C min −1.In situ HEXRD.The experimental set-up for the in situ experiment was similar to that reported in previous publications 19,20.The experiment was carried out at beamline 11-ID-C of the Advanced Photon Source,Argonne National Laboratory;the X-ray wavelength was0.10798Å.A DSC sample contained in a high-pressure stainless-steel vessel was placed vertically in a programmable thermal stage,and the sample was heated up to350◦C with a constant heating rate of10◦C per minute. During the course of thermal ramping,high-energy X-rays penetrated through the sample horizontally,and a Perkin Elmer area detector was used to collect the X-ray diffraction patterns in the transmission geometry with a spectrum data collection rate of one pattern every20s.The collected2D patterns were then integrated into conventional1D data(intensity versus2θ)using the fit2d program23. 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AcknowledgementsThis work was supported by the Human Resources Development of the Korea Institute of Energy Technology Evaluation and Planning(KETEP)grant funded by the Korea government Ministry of Knowledge Economy(No.20114010203150)and the National Research Foundation of KOREA(NRF)grant funded by the Korea government(MEST; No.2009-0092780).Research at Argonne National Laboratory was funded by theUS Department of Energy,EERE Vehicle Technologies Program.Argonne National Laboratory is operated for the US Department of Energy by UChicago Argonne, LLC,under contract DE-AC02-06CH11357.The authors also acknowledge the use of the Advanced Photon Source of Argonne National Laboratory supported by the US Department of Energy,Office of Science,Office of Basic Energy Sciences.Author contributionsY-K.S.and K.A.proposed the concept.Z.C.,H-J.N.,D-J.L.,H-G.J.,S-T.M.,C.S.Y., Y.R.and S.W.performed the experiments and acquired the data.Y-K.S.,Z.C.andK.A.wrote the paper.Additional informationSupplementary information is available in the online version of the paper.Reprints and permissions information is available online at /reprints.Correspondence and requests for materials should be addressed to Y-K.S.or K.A.Competingﬁnancial interestsThe authors declare no competing financial interests.。