Porous Electrode Materials for Lithium-Ion Batteries – How to Prepare Them and What Makes Them Special
Numerous benefi ts of porous electrode materials for lithium ion batteries (LIBs) have been demonstrated, including examples of higher rate capabili-ties, better cycle lives, and sometimes greater gravimetric capacities at a given rate compared to nonporous bulk materials. These properties promise advantages of porous electrode materials for LIBs in electric and hybrid electric vehicles, portable electronic devices, and stationary electrical energy storage. This review highlights methods of synthesizing porous electrode materials by templating and template-free methods and discusses how the structural features of porous electrodes infl uence their electrochemical prop-erties. A section on electrochemical properties of porous electrodes provides examples that illustrate the infl uence of pore and wall architecture and inter-connectivity, surface area, particle morphology, and nanocomposite formation on the utilization of the electrode materials, specifi c capacities, rate capa-bilities, and structural stability during lithiation and delithiation processes. Recent applications of porous solids as components for three-dimensionally interpenetrating battery architectures are also described.
1. Introduction
Since they were fi rst commercialized about twenty years ago, rechargeable lithium-ion batteries (LIBs) have become ubiqui-tous power sources for portable devices used in a wide range of consumer, health, and military applications. They are now beginning to enter the market in the transportation sector, and for load leveling and large-scale storage of electrical energy from alternative power sources, such as wind and solar energy. Batteries based on the Li + /Li redox couple are attractive for high voltage and high capacity applications, lithium being the most electropositive metal (−3.04 V vs. a standard hydrogen electrode) and having a very low atomic mass. Over two dec-ades of development, cell capacities have slowly improved, but the main components, including cathode and anode materials, have not changed signifi cantly. A typical modern LIB cell con-sists of a cathode made from a lithium-intercalated layered oxide (e.g., LiCoO 2 , LiNi x Co 1-x O 2 , or LiNi x Mn y Co 1-x-y O 2 ) and a graphite anode with an organic electrolyte. The battery operates
A. Vu, Y. Qian, Prof. A. SteinDepartment of ChemistryUniversity of Minnesota207 Pleasant St. SE., Minneapolis, MN 55455 E-mail: [email protected]
DOI: 10.1002/aenm.201200320
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following a “rocking chair” concept with an initial charging step, during which lithium ions are extracted from LiMO 2 (M = Co, Ni, Mn) and intercalate into graphite, and a subsequent discharging step, during which lithium ions deintercalate from the graphite particles and are reintroduced into the layered Li 1-x MO 2 structure. [ 1 ] Both electrodes are composed of micrometer-sized active particles, mixed with conduc-tive carbon particles to lower the electrical resistance with the electrodes, and held together by a binder phase. Each cell can be discharged and recharged for many cycles as long as the cathode and anode do not undergo signifi cant structural changes.
Numerous applications are now putting increased demands on electrical energy storage devices, calling for higher specifi c capacities and faster rate performance. [ 2 ] These requirements are particularly
important for electric vehicles (EVs) and plug-in hybrid electric vehicles (HEVs), where energy storage devices must supply suf-fi cient power to accelerate a vehicle quickly and recover energy during braking. In transportation applications, shorter recharge times would increase the consumer acceptance rate of EV and HEV technology. Improved energy densities and rate perform-ance are also required for devices storing charge from renew-able energy sources, such as solar and wind energy, which are seasonal and intermittent.
One commonly used solution for high power requirements is to employ capacitors or supercapacitors that can deliver or take up charge at faster rates than batteries. In capacitors, charge is stored on the electrode surface, allowing fast delivery but only relatively low energy density. In contrast, a battery stores charge in the bulk, thereby providing greater energy density. Unfortu-nately, the rate of charge delivery is limited by Li-ion diffusion through the bulk particles. Recent research has led to improve-ments for both capacitors and batteries. By adding redox charge storage (pseudocapacitance) to double-layer storage, energy densities of supercapacitors are now being increased. On the other hand, rate capabilities of batteries can be improved by structuring the electrodes appropriately. The characteristic time for lithium ions to diffuse through an electrode material ( τ eq ) varies as the square of the characteristic diffusion length ( l ) according to the relation τeq ∼ l2
D , where D is the diffusion coeffi cient. [ 3 ] To decrease the diffusion time, one can either increase the diffusion coeffi cient (by synthesizing better Li +
Anh Vu received his B.S. degree in Analytical Chemistry from Hanoi University of Science, Hanoi, Vietnam, in 2005. He is currently pursuing his Ph.D. degree under the supervision of Prof. A. Stein. His research focuses on nanoporous and nanostructured materials for energy storage and sensor applications.
Yuqiang Qian was born in Hebei province, China. He graduated from Fudan University, Shanghai, in 2008 with a B.S. in Chemistry. He is currently a doctoral candidate in Chemistry at the University of Minnesota, working on nanomaterials for energy storage and polymer nanocom-posite applications under the supervision of Prof. A. Stein.
Andreas Stein is a Distinguished McKnight University Professor at the University of Minnesota. He received his Ph.D. in Physical Chemistry from the University of Toronto in 1991 and car-ried out postdoctoral research at Bayer A.G., Germany, the University of Texas, and Penn State University. Professor Stein’s research interests are
in the fi eld of materials chemistry, in particular porous materials, templating methods, and nanocomposites.
conductors) or decrease the diffusion length (by using electrode components with nanometer dimensions). The latter approach avoids changes in the battery chemistry, has a larger impact on the diffusion time because of the square relationship, and has therefore attracted the attention of many researchers in recent years. A particularly effective way of achieving critical dimen-sions of active material on nanometer length scales is to employ porous electrodes, the focus of this review.
Porous materials have been widely used in areas such as heterogeneous catalysis, adsorption, separation, gas storage, sensing, etc. [ 4 ] Electrical energy storage systems, including bat-teries and supercapacitors, are among the applications that can profi t from porous materials. Porous electrodes offer important benefi ts that are listed briefl y below and will be discussed in detail later in this review.
• Pores provide good access of the electrolyte to the electrode surface.
• The surface area in a porous material is relatively large, there-by facilitating charge transfer across the electrode/electrolyte interface.
• The walls of active material surrounding the pores can be very thin (nanometers to tens of nanometers), reducing path lengths for ion diffusion.
• The small feature sizes permit increased utilization of active material (more utilized volume, deeper cycling), so that spe-cifi c capacities can be increased, particularly at high charge/discharge rates.
• The walls and pores in a porous electrode can be bicontinu-ous, thereby providing continuous transport paths through the active phase (walls) and the electrolyte phase (pores).
• Although porosity usually decreases the capacity per unit vol-ume (volumetric capacity) of an electrode, a small number of examples of porous electrodes have been shown to pro-vide increased volumetric capacity compared to packed nanoparticles.
• The nanosized features in a porous solid are available within a material that has larger dimensions and can therefore be han-dled and processed more easily than discrete nanoparticles.
• In some cases, less or no binder is needed to hold the active phase together.
• The void spaces separating particles of active material can help to constrain growth of active material during cycling.
• In nanosized particles, irreversible phase transformations that occur in microcrystalline anodes can be suppressed. Such particles are therefore better able to accommodate volume changes due to fi rst order phase changes during cycling.
• It is also possible to synthesize porous composite electrodes in which a supporting structure stabilizes active components with short cycle lives (e.g., components that disintegrate due to large volume changes during cycles).
• Porous composites can incorporate a secondary conductive phase to improve conductivity and high rate capacities of ac-tive phases with low intrinsic conductivity. Inclusion of the conductive phase eliminates or reduces the amount of con-ductive carbon additive needed in the fi nal electrode.
At this point, porous carbon materials are among the most prominent porous electrode materials, being found in
capacitors, [ 5 ] and different battery systems (lithium-ion, lithium-air, and lithium-sulfur batteries), where they help to enhance the system performance when high power or high current rates are required. [ 6–8 ] However, other porous materials, such as metal oxides, metal phosphates, metals, alloys, and composites have been investigated as electrodes, solid electrolytes, or cur-rent collector materials over the past years for both capacitor and battery applications. In this review, we will concentrate on porous materials as electrodes for LIBs, since progress in supercapacitors relying on porous carbon materials has recently been reviewed. [ 9 ] We will fi rst discuss methods of synthesizing porous electrodes with various pore sizes. We will then review performance parameters that can benefi t from porous electrode
architectures and provide specifi c examples from the literature that address these parameters. We will conclude with a brief discussion on the future role of porous electrodes in LIBs. The focus of the review will be largely on active materials other than carbon, although some examples of carbon composites with active materials will be provided.
2. Synthesis of Porous Electrodes for LIBs
Porous materials can be classifi ed based on their pore sizes (micropores < 2 nm, mesopores 2–50 nm, and macropores > 50 nm), their morphologies (ordered and non-ordered), and methods of their synthesis (templated and non-templated). One also distinguishes between “textural” porosity (voids created by the packing of particles), and “true” or “integral” porosity (from pores that are an integral part of a continuous solid frame-work). The pore sizes in texturally porous materials are usually correlated to particle sizes, whereas in materials with integral porosity they are independent of particle size. Textural porosity exists in conventional electrodes that are obtained by casting and drying of a slurry of active material, binder, and conductive
Figure 1 . (A) Schematic representation of the synthesis of mesoporous stemplating. Reproduced with permission. [ 11 ] Copyright 2005, American ChMechanism of Evaporation-Induced Self-Assembly (EISA) of mesostructuredduced with permission. [ 14 ] (C–E) Morphology of KLE-templated α -Fe 2 O 3 t550 ° C to remove the template. (C) Cross-sectional SEM image. (D) High magheight image. (E) TEM image and electron diffraction pattern. The imagescubic network of pores averaging 14–15 nm in diameter. The fi lm is crack-freelength scale, and pores at the top surface are open. Reproduced with permi
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additive particles on a current collector. The textural porosity then permits wetting of active material by the electrolyte. This article will concentrate mostly on porous materials with integral porosity. In the following discussion of synthetic methods for porous solids, we will emphasize the infl uence of the synthetic steps on the pore sizes, pore architecture, and dimensions of the walls in the synthesis products, as these parameters infl u-ence the electrochemical properties of the materials.
2.1. Soft Templating of Mesoporous Electrodes
Soft templating approaches toward porous electrodes typically involve surfactants as structure-directing agents (SDAs). The necessary chemistry has been well developed for silica systems, where surfactants direct the formation of mesoporous silicas with pore diameters covering most of the mesopore size range ( Figure 1 A). [ 10–12 ] The pore architecture of surfactant-templated materials can be controlled by the choice of surfactants, sol-vents, and synthesis conditions. It includes disordered mes-opores as well as ordered mesopores with hexagonal, cubic, lamellar, or other symmetries. Syntheses are generally carried
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hin fi lm heated to nifi cation 3D-AFM reveal a distorted at the micrometer ssion. [ 15 ] .
out in aqueous solution under hydrothermal or lower temperature conditions. Alterna-tively, non-aqueous solvent systems may be used. Particularly for thin fi lm formation, a process called evaporation-induced self-assembly (EISA) is widely employed (Figure 1 B). [ 13 , 14 ] These syntheses are based on self-assembly of the inorganic precursors either with the surfactant or around organized sur-factant assemblies, forming mesostructured products. Removal of the surfactant phase by extraction or calcination produces the mes-oporous material.
Various soft-templating processes have also been applied to the synthesis of mesoporous oxides, phosphates, and metals suitable for LIB electrodes. The versatility of the soft tem-plating approach can be illustrated by the dif-ferent types of SDAs that have been used to synthesize mesoporous TiO 2 , an inexpensive, promising anode material for LIBs with rea-sonably high capacity. Ordered mesoporous TiO 2 electrodes have been synthesized through soft-templating with cationic alkyl trimethy-lammonium surfactants (C 8 –C 16 TA + ), [ 16 , 17 ] anionic sulfonate surfactants, [ 18 ] the noni-onic surfactants Brij56 (C 16 EO 10 ) [ 19 ] and dodecylamine, [ 20 ] and Pluronic type block copolymers H(–OCH 2 CH 2 ) x [–OCH(CH 3 )CH 2 ] 70 –OCH 2 CH 2 ) x OH with x = 20 (P123) [ 21 ] or x = 106 (F127). [ 22 ] The infl uence of the sur-factant type on pore structure of TiO 2 and other electrode materials will be discussed with a few more examples.
Low molecular weight cationic surfactants of the alkyltrimethylammonium type have been employed to prepare both mesoporous
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cathode (Li 3 Fe 2 (PO 4 ) 3 , V 2 O 5 ) [ 23 ] and mesoporous anode (TiO 2 , SnO 2 , Sn 2 P 2 O 7 ) [ 16 , 17 , 24 , 25 ] structures. With a single alkyltrime-thyl ammonium surfactant, mesopore diameters in the cal-cined materials tend to be small (2.3–4 nm), but if a mixture of these cationic surfactants is used, a broader mesopore size dis-tribution with a larger average mesopore diameter is obtained (3–7 nm). [ 24 ] Somewhat larger mesopores were also obtained in the case of a synthesis of mesoporous anatase TiO 2 using C 8 –C 16 TA + cationic surfactants, where pore sizes increased from 5.7 to 7.0 nm with increasing chain length and surface areas followed a similar trend. [ 17 ] A more complex mixture of didodecyldimethylammonium bromide, hexane or tetradecane, and water was used in another synthesis of mesoporous ana-tase TiO 2 . [ 26 ] The bicontinuous microemulsion resulting from this water/surfactant/oil ternary system templated disordered mesopores with mean diameters in the 2.4–3.2 nm range. With increasing calcination temperature in the range from 300 to 380 ° C, the surface area and mesopore volume decreased and the average mesopore size increased. Mesoporous rutile TiO 2 with high crystallinity, another anode material, was synthesized using anionic sulfonate surfactants as the template with TiCl 3 as a precursor. [ 18 ] The multi-step synthesis included an oxidation step with hydrogen peroxide before a low temperature reaction (60 ° C) and calcination (400 ° C). The mesoporous product con-sisted of assemblies of rutile nanorods. Mesopore sizes were controllable in the range from 2.2 to 3.8 nm by modifying the length of hydrocarbon chains in the anionic surfactant. The surfactant appeared to play a role in promoting the rutile phase at the relatively low calcination temperature, although the exact role was not fully understood. Nonionic surfactants of the type octaethylene glycol monohexadecyl ether (C 16 EO 8 ) and related surfactants (C 18 EO 10 , C 18 EO 20 ) have been employed to syn-thesize mesoporous Sn fi lms (1.2 μ m thick) with disordered mesopores by electrodeposition on a copper foil substrate. [ 27 , 28 ] Although pore sizes were not specifi ed, repeat distances (which include the diameter of the mesopores and thickness of the wall) were in the 6–8.5 nm range. [ 28 ]
When higher molecular weight triblock-copolymer sur-factants are used to prepare mesoporous electrodes, the mes-opore size depends on the choice of the polymer surfactant and the particular materials composition. Mesoporous vana-dium oxide electrodeposited from a water/ethanol solution containing VOSO 4 and the nonionic block copolymer Pluronic P123 exhibited a wormlike mesopore structure with 3–4 nm mesopores. [ 29 ] Some macropores were also present in this material. Mesoporous materials with similar dimensions are obtained in most other syntheses employing P123. This is true even for mesoporous electrodes with much more complex compositions, such as a nanocomposite of anatase nanopar-ticles embedded in a glassy P 2 O 5 phase that was also loaded with a conducting material (CuO or SnO 2 ). [ 30 ] In this material, pore diameters were ca. 4 nm and wall thicknesses ca. 5 nm. Larger mesopores were obtained in syntheses of FePO 4 tem-plated with P123. [ 31 ] By increasing the surfactant to inorganic precursor ratio, the average pore diameter in the disordered mesopore system expanded from 6.2 to 10 nm. Pore volumes increased at the same time. Wall thicknesses of ca. 5 nm are typically observed when Pluronic-type surfactants (P123, F127) are used. [ 32 ]
Diblock copolymers of the KLE type (poly(ethylene- co -butylene)- block -poly(ethylene oxide) diblock copolymer) produce much larger mesopores and thicker walls. [ 15 ] Mesoporous hem-atite Fe 2 O 3 fi lms synthesized with a KLE surfactant by the EISA method contained mesopores with diameters of 14–15 nm and walls 10–16 nm thick (Figure 1 C–E). [ 15 ] In a synthesis of mes-oporous titania doped with niobium to increase the electrical conductivity of titania, a KLE surfactant produced disordered mesopores with sizes in the 10–20 nm range. [ 33 ] Mesoporous Li 4 Ti 5 O 12 (a “zero-strain” anode material) was synthesized by templating with KLE and contained mesopores with an average diameter of ∼ 18 nm and a nanocrystalline framework of phase-pure spinel nanoparticles. [ 32 ] In the calcined mesoporous mate-rial the nanocrystallites in the walls were 11 ∼ 15 nm in size. The KLE surfactant has the additional advantage of higher thermal stability than the Pluronic-type surfactants, allowing calcina-tion conditions that result in more complete condensation of the inorganic components. [ 34 ] Another amphiphilic diblock sur-factant with high thermal stability is polystyrene- b -polyethylene (PS- b -PEO). This surfactant was employed in the synthesis of mesoporous NbVO 5 thin fi lm electrodes with wormlike porous networks whose pore sizes and wall thicknesses could be tuned in the range from 15 to 100 nm, depending on the molecular weight of the surfactant. [ 34 ] Partial crystallization of the walls was possible before the pore structure collapsed due to sin-tering of NbVO 5 crystallites.
In syntheses of electrode materials using surfactant tem-plating, the processing conditions have to be chosen appro-priately for particular precursor and surfactant combinations, as illustrated for the case of mesoporous SnO 2 anodes. Early attempts to prepare a surfactant-free mesoporous SnO 2 elec-trode were not successful due to the collapse of the mesopore walls upon surfactant removal by solvent extraction or heat treat-ment. [ 35 ] Syntheses that preserved mesopores after surfactant removal require conditions in which a rigid framework based on Sn-O-Sn bonding is formed before surfactant removal. [ 36 ] This was fi rst demonstrated for syntheses involving either a neutral primary amine (tetradecylamine) or a cationic ammonium sur-factant (C 16 TA + ) as templates. [ 36 ] Subsequently, other successful syntheses of mesoporous tin oxide were reported using the ani-onic surfactant sodium dioctylsulfosuccinate as a template, [ 37 ] and relatively large mesopores (15 nm) were obtained with the non-ionic surfactants P123 and T908. [ 38 ]
2.2. Hard Templating
Soft templating is a powerful method to prepare mesoporous materials with different compositions and morphologies. How-ever, it is sometimes diffi cult to obtain ordered mesoporous materials with highly crystalline walls when the crystallization temperatures of the inorganic phase are higher than the temper-atures at which surfactants and block copolymers are removed. Without support from the assembled surfactants during crys-tallization, sintering of crystallites can eventually destroy the mesostructure. For those systems, hard templates can be used whose rigid structures do not change through interactions with a precursor, unlike the situation with soft templates. Hard templates include preformed porous solids, anodic aluminum
oxide (AAO) membranes, and assemblies of colloidal particles, such as silica- or polymer-based colloidal crystals. When hard templates are used, inorganic precursor solutions are infi ltrated into the templates, and then the composites are treated at high temperatures. During the heat treatment, the precursor may undergo gelation and crystallization within the support of the template. After the heat treatment, ideally the inorganic frame-work is strong enough to maintain the porous structure, even when the template is removed. In the following sections, we will review the application of hard templating to syntheses of porous electrodes for LIBs, considering nanocasting and col-loidal crystal templating. Hard templating in various biological templates and in AAO membranes will be discussed separately in Sections 2.4 and 2.7.
2.2.1. Nanocasting
A particularly versatile methodology for mesoporous elec-trodes that require high temperature processing is called “nanocasting”. [ 39 , 40 ] In the nanocasting process, nanoporous molds with high thermal stability, such as mesoporous silica or mesoporous carbon, are infi ltrated with the precursor for the target electrode material. This precursor is then concentrated in the pores via solvent evaporation. Often multiple impregna-tion steps are necessary to increase fi lling of the void space in the porous mold. [ 41 , 42 ] Vacuum-assisted impregnation [ 43 ] or the application of centrifugal forces [ 44 ] can also improve the extent of pore fi lling and improve the periodicity of the product struc-ture. After thermal processing, the template is removed: silica by extraction with HF or hot NaOH or KOH solutions, carbon by combustion. The product is an inverted replica structure of the original hard template. Thus, molds with bicontinuous pore/wall structures produce mesoporous replicas, whereas tem-plates with cylindrical mesopores in a 2D hexagonal arrange-ment produce nanowire arrays, which are interconnected only
Figure 2 . Nanocasting of mesoporous electrode materials. (A) A scheme illustrating the nano-casting process in mesoporous silica KIT-6 with a double-gyroid, cubic mesopore structure. The white areas represent mesopores, and the areas shaded in grey represent the silica phase. Dark areas correspond to the precursor in the composite and to the fi nal solid phase after template removal. (B) TEM image of mesoporous anatase TiO 2 , nanocast from KIT-6. Reproduced with permission. [ 47 ] (C) The nanocasting process in mesoporous silica SBA-15 with 2D hexagonal mesopore structure. Reproduced with permission. [ 40 ] In the fi nal replica, nanowires may be interconnected by small struts resulting from seconday pores. (D) TEM image of mesopo-rous Si-C nanowire arrays that were nanocast from SBA-15. Reproduced with permission. [ 48 ] Copyright 2008, American Chemical Society.
if the template contains secondary pores between the cylindrical pores.
The most commonly used hard templates for nanocasting of porous electrodes are the mesoporous silica materials KIT-6 [ 45 ] and SBA-15, [ 46 ] as both are easily synthesized and contain mesopores that are large enough to accommodate precursor fl uids for the target electrode. Both of these structure types are templated in acidic media by the triblock copolymer P123. For KIT-6, n -butanol is included as a co-structure-directing agent. KIT-6 has a well-ordered cubic mesopore structure with Ia-3d symmetry, 4.5–10 nm mesopores (depending on synthesis param-eters), and 2–4 nm thick walls ( Figure 2 A). It is important for the nanocasting process that KIT-6 synthesized at hydrothermal treatment temperatures > 343 K contains pore channels that are well connected, allowing for complete infi ltration of the structure. SBA-15 contains a highly ordered 2D hexagonal mesopore structure with p6m symmetry, large pores (4.6–11.4 nm, again depending on synthesis
parameters), and thick walls (3.1–6.4 nm) (Figure 2 C). [ 12 ] Because the long PEO chains of the P123 surfactant can be inserted into the silica pore walls, primary mesopores may be connected by secondary tunnels, both micropores and small mesopores. These are responsible for keeping nanowire arrays connected in nanocast replica structures, so that the latter are also mesoporous.
Mesoporous electrodes with numerous compositions have been synthesized by hard templating with KIT-6, including the cathode materials β -MnO 2 (tetragonal rutile structure), [ 49 , 50 ] LiFePO 4 (olivine structure), [ 51 ] and Li 1 + x Mn 2-x O 4 (spinel struc-ture), [ 52 ] as well as the anode materials Cr 2 O 3 , [ 43 , 53 ] Co 3 O 4 , [ 41 , 54 ] anatase TiO 2 , [ 47 ] rutile TiO 2 , [ 55 ] NiO, [ 42 ] , SnO 2 , [ 56 ] WO 3-x , [ 57 ] and MoO 2 . [ 58 ] In some of these examples (e.g., SnO 2 ), nanocasting methods produce mesoporous SnO 2 with higher crystallinity and more ordered mesostructures than soft templating because the hard template protects particles from excessive grain growth by sintering. As discussed later, many of these electrodes dem-onstrated superior rate capacity compared to corresponding electrodes prepared from bulk particles. The replica structures typically exhibit similar order as the original KIT-6 template and much higher surface areas than bulk particles. Several of the studies pointed out complex pore structures in the products. For example, mesoporous WO 3-x templated from KIT-6 exhibited a bimodal mesopore size distribution (4 nm, ∼ 20 nm pores). [ 57 ] For mesoporous TiO 2 with ca. 6.5 nm-thick anatase walls, a trimodal pore size distribution was found with peaks at 5 and 11 nm resulting from inverse replication of the silica template and at 50 nm from interparticle voids (Figure 2 B). [ 47 ]
In the context of LIB electrode materials, nanocasting with SBA-15 has been applied to synthesize anatase and rutile TiO 2 , [ 55 , 59 ] Cr 2 O 3 , [ 43 ] LiFePO 4 , [ 51 ] Co 3 O 4 , [ 41 ] Si, [ 48 ] and SnO 2 , [ 56 ] where either discrete or connected mesoporous nanowire bundles were observed. As a typical example, the mesoporous nanowires in the case of SnO 2 were composed of aggregated
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Figure 3 . (A) General method to prepare 3DOM electrode materials. (B,C) SEM images of 3DOM V 2 O 5 calcined at 400 ° C. (D) SEM image of 3DOM SnO 2 calcined at 500 ° C and (E) at 800 ° C. (F) SEM image of 3DOM LiNiO 2 calcined at 600 ° C and (G) at 650 ° C. Reproduced with permission. [ 64 ] Copyright 2003, Electrochemical Society.
5–6 nm nanoparticles. [ 56 ] A similar structure consisting of packed nanowires was observed for mesoporous MoS 2 , nano-cast from SBA-15 using phosphomolybdic acid and treatment with H 2 S. [ 60 ] Although it is usually desirable to remove all of the silica template from the electrode material because silica does not contribute to lithium ion storage and is electronically insu-lating, residual silica can provide some mechanical stabilization. For instance, in the case of mesoporous SnO 2 templated from SBA-15 by melt-infi ltration of SnCl 2 ·2H 2 O, a small amount (6 wt%) of residual SiO 2 resulting from incomplete etching was believed to prop up the SnO 2 phase. [ 61 ] As a result, the mesopo-rous SnO 2 was structurally stable up to 700 ° C.
Instead of using a mesoporous silica as a host, mesopo-rous carbon can be impregnated with active material without the need to remove this matrix. The carbon scaffold provides a built-in conductive network and facilitates higher charge/discharge rates for the nanocomposite. An example involves a mesoporous carbon nanocomposite with NiO. [ 44 ] The CMK-3 mesoporous carbon was itself prepared by nanocasting in mes-oporous SBA-15 silica, using a sucrose/sulfuric acid solution as the precursor. Subsequently, CMK-3 was heated in sulfuric acid and then impregnated with an aqueous solution containing nickel nitrate and citric acid. The composite was heated in nitrogen, forming an intimate mixture of mesoporous carbon and ∼ 37 wt% NiO nanoparticles. While the mesopore volume dropped to 10% of its original value, signifi cant micropore volume remained in the nanocomposite.
Mesoporous Si is quite challenging to synthesize due to the diffi culty in preventing the formation of the insulating oxides SiO 2 and SiO 2-x while preserving the mesostructures. Most Si-containing precursors for surfactant-templated syntheses, such as alkoxides, produce SiO 2 when they are hydrolyzed, which requires a strong reducing atmosphere for reduction to Si. Such treatments usually lead to the collapse of the mesostruc-tures. Removal of the surfactant template by calcination or sol-vent extraction would also introduce a silicon oxide layer. The use of a hard template to prepare mesoporous Si does not make the situation much easier. Carbon and polymer-based hard templates can only be completely removed by calcination in an oxygen-containing atmosphere under conditions that do not prevent the formation silicon oxides, although it can be desir-able to leave carbon in the product to protect Si. Mesoporous silica as a hard template is also tricky to work with, because SiO 2 readily reacts with Si precursors to form SiO 2-x during the annealing step. [ 62 ]
To prevent a reaction between Si and a SiO 2 hard template, it is necessary to use a Si precursor terminating in carbon groups. Kim and Cho prepared mesoporous silicon-carbon nanowire arrays using SBA-15 as the template. [ 48 ] Si was fi rst capped with n-butyl groups by the reaction of LiC 4 H 9 with SiCl 4 to form Si(C 4 H 9 ) 4 . SBA-15 was mixed with the butyl-Si solution and annealed at 300 ° C several times to fully impregnate the tem-plate with the precursor. The Si-butyl-SiO 2 composite was fur-ther annealed at 900 ° C to convert it to a Si-carbon-SiO 2 com-posite, which was then treated with HF solution to selectively extract SiO 2 . Because Si is protected by thin layers of carbon formed from the decomposition of n -butyl groups (6 wt% of the composite is carbon), it is not affected by the HF treatment step. The Si-C composite of nanowire arrays showed a well-ordered
hexagonal structure (Figure 2 D) with a d -spacing of 8.8 nm, a BET surface area of 74 m 2 g − 1 , and an average pore size of 2.3 nm. It showed good electrochemical performance, which was attributed to the evenly distributed mesoporous channels that buffered against volume changes and to the carbon layer that prevented the aggregation of Si wires (see also Sections 3.7 and 3.8).
2.2.2. Colloidal Crystal Templating
Colloidal crystal templating (CCT) is a versatile method to obtain porous electrodes with well-ordered, interconnected pores of larger sizes (large mesopores and macropores). [ 63 ] The general method is described in Figure 3 A. In this method,
monodisperse spheres (polymer spheres or silica spheres), are assembled into periodic arrays, so-called colloidal crys-tals. The void space between spheres is then infi ltrated with a precursor solution for the target electrode material. Thermal processing converts the precursor to a solid skeleton that sur-rounds the templating spheres. Calcination in air above 350 ° C also removes polymer sphere templates, such as polystyrene or poly(methylmethacrylate) spheres. If pyrolysis is carried out in an inert atmosphere, polymeric templates and organic precursor components are partially converted into a carbon phase that coats the remaining skeleton. Silica spheres may be removed by etching, e.g., with hydrofl uoric acid solution. As in the case of nanocasting from mesoporous silica, the fi nal product of the CCT process is an inverted replica of the original colloidal crystal, i.e., an “inverse opal” or three-dimensionally ordered macroporous (3DOM) or mesoporous (3DOm) solid. The inverse replica typically inherits the face-centered cubic ( fcc ) symmetry of the templating sphere array. The fcc macropore array is interconnected through windows where templating spheres were in direct contact (i.e., at twelve points in an ideal and typical structure). As a result, the cubic 3DOM structure is bicontinuous, with a continuous solid phase (the active elec-trode material) and a continuous void space suitable for accom-modating, for example, an electrolyte solution.
The pore sizes can be modifi ed by adjusting the sizes of col-loidal polymer spheres or silica spheres. For large mesopores (20–50 nm) and small macropores (50–250 nm), silica spheres are the most suitable templates as they can be prepared in these sizes ranges with relatively low polydispersity. Silica-based col-loidal crystal templates are also advantageous for electrode materials that require processing at high temperatures (above ca. 600 ° C) as the silica matrix limits sintering of grains and minimizes shrinkage of the structure. For the pore size range from ca. 100–2000 nm, polymer spheres are often preferred because they can easily be synthesized in higher yield than uniform silica spheres, also with low polydispersity. Removal of polymeric templates does not require an additional etching step. However, larger shrinkage is observed with polymeric tem-plates, so that the fi nal pore spacing is often 20–40% smaller than the repeat distance between the templating spheres. Although templating spheres with greater polydispersity may also be employed in this type of hard templating, they will not produce periodic structures. Therefore, a high degree of inter-connectedness between pores and low pore tortuosity are not guaranteed in those cases.
Besides the pore architecture, an important consideration for porous electrodes is the wall structure. The morphologies of the walls, their crystallinity, and surface texture can be con-trolled by changing the conditions for template removal. Typi-cally, removing the templates at low temperatures or through etching (solution treatment) yields amorphous walls, whereas removing them at high temperatures produces crystalline sam-ples in many cases. For crystalline 3DOM materials, the macro-pore walls are composed of interconnected crystallites. The crystallite size depends on the thermal history of the material. In general, crystallites have dimensions of a few nanometers to a few tens of nanometers and overall wall thicknesses are tens of nanometers, being thinner at the narrow struts and thicker at the vertices (replicas of the octahedral and tetrahedral holes
in the fcc colloidal crystal). This provides short diffusion paths for lithium ions through these walls, similar to the situation for discrete nanoparticles, but in a macroscopic material that can be more easily handled than nanoparticles.
We will illustrate CCT syntheses with a few examples. Dunn and coworkers [ 65 , 66 ] fabricated a 3DOM V 2 O 5 aerogel cathode on an indium tin oxide-coated substrate. Polystyrene spheres (1 μ m in diameter) were fi rst assembled into an fcc array by centrifugal and capillary forces. The template was infi ltrated with a vanadyl triisopropoxide precursor, which was then hydrolyzed to transform it into a solid vanadia gel network. The polystyrene template was removed by extraction with toluene. Water produced during the condensation of the vanadia precursor was exchanged with low surface tension sol-vents through multiple solvent exchange steps to mitigate the effects of capillary forces. After removal of the fi nal solvent, the aerogel-like product contained 800 nm-diameter macropores and ∼ 150 nm-thick mesoporous walls composed of V 2 O 5 rib-bons with nanometer dimensions. Both mesopores and macro-pores were readily accessible to electrolyte solutions in this hierarchical structure.
Yan et al. developed templated syntheses for different 3DOM electrode materials, including cathode materials (Shcherbinaite-type V 2 O 5 and LiNiO 2 ) and an anode material (Cassiterite-type SnO 2 ). [ 64 ] Centimeter-sized colloidal crystals were soaked in a precursor solution. Any extra precursor was removed by vacuum fi ltration. The infi ltrated colloidal crystals were dried in air and then calcined at the desired temperature (between 400 and 800 ° C) to remove the template and form the crystal-line 3DOM material. All the samples showed periodic struc-tures with uniform pores, open windows, and interconnected walls (Figure 3 B–F), although the morphologies of 3DOM SnO 2 and LiNiO 2 depended strongly on the calcination temperature. At higher temperatures, the windows became smaller and the walls became rougher due to grain growth and sintering effects. In the case of LiNiO 2 , the periodic 3DOM structure was lost when the material was calcined at a temperature of 650 ° C or higher, but an interconnected pore structure was still main-tained (Figure 3 G). When the size of the templating PMMA spheres was increased ca. tenfold, the 3DOM structure could be maintained at 650 ° C because of the thicker walls.
Although the general steps of the synthesis of 3DOM elec-trodes are conceptually simple, details in the precursor chem-istry, solvents, and template surface chemistry greatly infl u-ence the product structures and properties. A recent example demonstrating the importance of precursor chemistry relates to 3DOM LiMn 2 O 4 . [ 67 ] In this system, a templated structure could be obtained from an ethanolic manganese nitrate precursor solution but not from a manganese acetate solution in a water/ethanol mixture. Similar precursor effects had been recognized earlier for general syntheses of 3DOM materials employing salt precursors. [ 68 ] Important considerations include the ability of the precursor to wet the template so that the templated is fully infi ltrated, the achievable precursor concentration which infl uences the solid fi lling fraction, and the thermal decomposi-tion/reaction behavior of the precursor and the template. For polymeric templates, a skeletal structure must be suffi ciently well developed before the template loses its structure. However, even if some structural order is lost due to crystallite growth,
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the resulting electrode material may still contain open, inter-connected pores, as was observed in syntheses of macropo-rous LiCoO 2 powders. [ 69 ] In these syntheses, which employed a poly(methylmethacrylate) template and a precursor con-taining lithium acetate/cobalt acetate mixtures, addition of poly(ethylene glycol) or platinum modifi ers permitted control over LiCoO 2 grain sizes in the wall skeleton.
The synthesis of macroporous Si also requires careful selection of Si precursors and annealing conditions that limit the oxidation of Si at high temperatures. Kim and coworkers prepared macroporous Si using the same approach as for the preparation of Si-C nanowires from SBA-15 mentioned earlier, but employed a silica sphere template. [ 70 ] A butyl-capped Si gel was mixed with silica spheres and then annealed under Ar for 3 h at 900 ° C. The spheres were etched out by treatment with HF, leaving macroporous Si with a protective carbon layer. The composite material showed good rate capabilities (discharge capacities of 2668, 2471, 2158 mA h g − 1 at 1C, 2C, and 3C rates, respectively) and long life-cycles (99% and 90% capacity reten-tions at 0.2C and 1C rates, respectively). After 100 cycles, the morphology of the macroporous Si was preserved, however, the wall thickness increased from 40 to 70 nm.
Electrodeposition methods have also been useful to form 3DOM metals or alloys, either as anode materials or as cur-rent collectors. Macroporous Al fi lms were synthesized by elec-trodeposition from a chloraluminate ionic liquid within a poly-styrene template on a copper support. [ 71 ] In this electrochemical deposition, the metal grows from the substrate towards the top
Figure 4 . Different methods for fabricating hollow spheres; (A) templating with colloidal spheres; (B) mesoporous silica is used as nanoreactor; (C) template-free synthesis based on an Ostwald ripening mechanism; (D) double-shelled hollow spheres based on inward and outward ripening processes. Part (B): Adapted with permission. [ 92 ] Copyright 2004, American Chemical Society. Part (C): Adapted with permission. [ 93 ] Part (D): Adapted with permission. [ 94 ] Copyright 2009, American Chemical Society.
of the template, insuring effective fi lling of the void space by the metallic product. A 3DOM Sn-Co alloy fi lm electrode (80:20 wt. ratio of Sn:Co) was prepared by electroplating within a colloidal crystal template supported on a Ni-coated Cu sheet. [ 72 ] The nickel layer on the substrate was added to prevent corrosion by the latex spheres. 3DOM Ni prepared by elec-trodeposition within a polystyrene colloidal crystal template was used as both the sup-port and current collector for a lithiated MnO 2 cathode with extremely fast rate capabilities. [ 73 ]
Over the last decade, numerous other 3DOM electrodes have been synthesized by colloidal crystal templating, including those composed of the cathode materials LiMn 2 O 4 , [ 74 ] LiFePO 4 , [ 75 ] and FePO 4 , [ 76 ] and the anode materials anatase TiO 2 , [ 77–79 ] Li 4 Ti 5 O 12 , [ 80 , 81 ] CoFe 2 O 4 , [ 82 ] and SnO 2 . [ 83 ] In addition, a 3DOM solid electrolyte of the com-position Li 1.5 Al 0.5 Ti 1.5 (PO 4 ) 3 (LATP) has been synthesized. [ 84 ] Variations of colloidal sphere templating with lower dimensionality (2D fi lms) have also been developed to prepare porous electrode structures. A 2D macropo-rous NiO fi lm was prepared on a stainless steel substrate by fi rst electrophoretically depositing a polystyrene sphere monolayer on the conducting substrate and then anodically electrodepositing nickel oxy-hydroxide. [ 85 ] Calcination at 400 ° C produced the
macroporous nickel oxide fi lm that had excellent electronic con-tact with the metallic substrate, permitting use of this electrode under high rate lithiation conditions.
2.3. Hollow Spheres as Electrode Materials
In the synthesis of 3DOM materials, the template consists of spherical colloids that are assembled into three-dimensional arrays, in which the precursor forms an extended network during gelation and calcination processes. If discrete spheres are used instead, they can act as templates for hollow spheres. Hollow inorganic spheres, particularly nanospheres, are attrac-tive electrode materials because of their high surface areas, short lithium diffusion paths, and high packing effi ciency. For alloyed anodes, the hollow nanospherical structures are particu-larly valuable since the hollow interior can act as a space buffer against volume changes that occur during electrochemical cycling, and the isotropic nature of the spheres helps to diffuse stresses evenly in all directions (see Section 3.7). Additional buffering space becomes available if the shells of the hollow nanospheres are themselves porous. Furthermore, if the pores within the walls are large enough for the electrolyte to pass through, the inner walls become accessible, increasing the use-able surface area.
A straightforward way to synthesize hollow spheres involves the use of hard templates such as polymer or silica spheres ( Figure 4 A). The templating spheres are coated with the
precursor for the target materials, which is then solidifi ed by gelation or heat treatment. The coating process may be car-ried out under conditions that keep these spheres separate from each other, for example by stirring and sonication, or by ultrasonic spray pyrolysis. The templating spheres are removed either by high-temperature treatment or by solvent extrac-tion to create the hollow interior, similar to the procedure for 3DOM materials. As an example, hollow carbon spheres (750 nm diameter, 50 nm wall thicknesses) were prepared by chemical vapor deposition (CVD) of benzene around 650 nm silica spheres, followed by treatment of the composite with 20 wt% HF solution. [ 86 ] SnO 2 nanoparticles were then coated onto the walls of the hollow carbon spheres to produce carbon-SnO 2 hollow sphere composites. The coating step was carried out by microwave-assisted hydrolysis of SnCl 4 in the presence of urea, which produced uniform SnO 2 nanoparticles (1–3 nm) on the carbon sphere surface. The sizes of the SnO 2 particles and the carbon/SnO 2 ratios were controlled by the temperature of a subsequent calcination step in air. The advantages of these composite structures for anode applications include the ability of the carbon shell to trap internal SnO 2 , even if electrochem-ical cycling removes particles from the internal surface and the short Li-ion diffusion lengths through the thin shell. A CVD method has also been applied to the synthesis of hollow silicon spheres using silica spheres as the template, but the product obtained after HF etching consisted of interconnected amor-phous silicon shells rather than discrete spheres. [ 87 ]
Solution or sol-gel processing to synthesize hollow spheres have been applied to SnO 2 , Sb, and Li 4 Ti 5 O 12 compositions. [ 88 , 89 ] To improve the conductivity and mechanical stability of hollow spheres, they can be coated with layers of carbon. Concentric SnO 2 -carbon hollow spheres were prepared by hydrothermally depositing a carbonaceous layer on a SnO 2 -SiO 2 composite, then carbonizing it before etching out the silica. [ 90 ] The products con-tained agglomerates of spheres in addition to discrete spheres.
Another scalable approach involves spray pyrolysis. Mesopo-rous TiO 2 -bronze microspheres were synthesized by ultrasonic spray pyrolysis from a mixture containing a titania precursor and colloidal silica, followed by etching away of the silica phase with NaOH solution, acid-washing and calcination. [ 91 ] The com-posite contained nanocrystalline TiO 2 -bronze grains (5–10 nm) that surrounded uniform mesopores (10–15 nm).
Hard templates such as mesoporous silica can also be used as nanoreactors to produce hollow spheres of a metal oxide such as SnO 2 . [ 92 ] The synthesis steps are illustrated in Figure 4 B. First, the precursor for the target material, for example a molten metal salt hydrate, is infi ltrated into a hollow silica sphere with a mesoporous shell. The infi ltrated nanoreactors are thermally treated to transform the precursor into an oxide that coats the inside of the shell, forming SiO 2 -metal oxide double shell hollow spheres—provided that the silica-precursor interactions are suf-fi ciently strong. If the formed metal oxide interacts poorly with the SiO 2 surface, a dense SiO 2 -metal oxide composite sphere is formed. The outer SiO 2 shell is etched away with HF solution to obtain hollow metal oxide spheres of the target materials.
Templating methods for the fabrication of hollow struc-tures are able to produce thermally stable spheres with narrow size distribution, whereas template-free methods are more advantageous in terms of cost and scalability. One particularly
interesting template-free approach relies on Ostwald ripening processes, in which small crystals dissolve and then recrystal-lize on the surface of other crystals to form bigger crystals. This process is driven by a reduction in the total surface energy of the system. An Ostwald ripening mechanism has been applied to the preparation of hollow SnO 2 structures based on the hydrolysis of potassium stannate in a water/ethanol mixture under hydrothermal conditions. [ 93 ] The mechanism is outlined in Figure 4 C. First, solid spheres are formed by the hydrolysis of stannate. These undergo inside-out Ostwald ripening during the hydrothermal treatment. The ripening process can begin just below the surface and slowly proceed towards the center of the spheres, resulting in the formation of hollow spheres with solid cores of different sizes inside (inward ripening). If the rip-ening process starts from the center of the hard spheres, only hollow spheres are obtained (outward ripening).
V 2 O 5 -SnO 2 double-shelled hollow spheres were prepared by a solvothermal synthesis, relying on the combination of both inward and outward Ostwald ripening processes (Figure 4 D). [ 94 ] During the solvothermal reaction of SnCl 4 ·5H 2 O and vanadium (IV) acetylacetone in N,N’-dimethylformamide, hard V 2 O 5 -SnO 2 composite spheres were formed, whose surfaces were composed of small crystallites that acted as nuclei for the inward ripening process. At the same time, an outward ripening process started from the core of the hard spheres, proceeding at a comparable rate. The two simultaneous processes, followed by calcination, led to the formation of double-shelled hollow nanocapsules. It was critical to include both vanadium and tin precursors, because single-shelled hollow spheres were obtained with only the vanadium precursor and solid SnO 2 spheres were obtained with only the tin precursor.
Besides hollow spheres, other hollow structures such as hollow nanotubes have been utilized as electrode materials for lithium ion batteries. SnO 2 nanotubes (200 nm diameter, micrometers length) were prepared using alumina membranes as templates and coated with a thin carbon shell by CVD. [ 95 ] Hollow nanorods of SnO 2 were prepared hydrothermally using 1D chiral silica nanorods as sacrifi cial templates. [ 96 ]
2.4. Biotemplates
As an alternative to synthetic templates, a variety of natural templates or preforms for porous electrodes have been inves-tigated in recent years. Some of these are relatively abundant, originate from renewable materials (such as diatoms and wood), and are capable of producing quite complex, often hierarchical structures. Among various approaches of using biomaterials as templates, [ 97 ] we will consider replication by nanocasting or by chemical conversion, and assembly of biological building blocks into complex structures to produce periodic porous structures or hierarchical porous structures.
Diatoms and wood are the most commonly used biotem-plates to prepare macroporous materials, including some that may potentially be used as porous electrodes for LIBs. Wood is mainly composed of cellulose, hemicellulose, and lignin, forming a cellular microstructure with high porosity and interconnectivity. The open porous system of the tracheidal cells acts as the transportation path for water and minerals in the living wood. [ 98 ] The morphology and arrangement of the
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cells on the wall depend on the types of wood. Coniferous wood cells are made up of tracheids, which are long and slender cells tapered at the end. These cells are quite uniform in size, up to 50 μ m in diameter and a few millimeters in length, depending on growth conditions. Deciduous wood cells are composed of tracheids and tracheary elements that form long tubes, which are oriented in the direction of the trunk axis. [ 98 ] These wood tissues can be transformed into carbon when treated at high temperatures in an inert atmosphere, and they can be removed by oxygen or air treatment at high temperatures, making them useful templates. Potential electrode materials such as TiO 2 , NiO, and Mn 2 O 3 have been fabricated through wood templating combined with sol-gel processes. [ 99–103 ] Porous TiO 2 was prepared by infi ltrating wood templates with titania precursors, including titanium tetraisopropoxide in isopro-panol, [ 104 ] titanium tetrachloride in ethanol, [ 99 ] or titanium(IV) bis(ammonium lactato) dihydroxide solution. [ 100 ] The pre-cursor was hydrolyzed to form an in-situ gel network around the wood template. The wood-titania gel composite was dried to strengthen the gel network, then calcined in air at high tem-peratures to remove the template and to convert the precursor into TiO 2 . Porous NiO and Mn 2 O 3 were fabricated by the same procedure using Ni(NO 3 ) 2 and Mn(NO 3 ) 2 as metal precursors, respectively. In all of these samples, the porous structures were characterized by interconnected macropores surrounded by thin walls ( Figure 5 A, B). These structures are thought to be useful for high rate performance LIBs as the open and inter-connected macropores facilitate the mass transport and the thin walls reduce the diffusion path of lithium ions. However, very limited information on the electrochemical performance of these materials has been published so far.
Diatoms are derived from one of the most common types of phytoplankton and are characterized by their unique cell walls
Figure 5 . Templating with biomaterials. SEM images of (A) Paulownia wood-templated biomorphic MnO 2 , and (B) pine-templated NiO. (C) Aulacoseira diatom frustule, (D) silicon bearing replica. Part (A): Repro-duced with permission. [ 103 ] Copyright 2006, Elsevier Ltd. Part (B): Repro-duced with permission. [ 102 ] Copyright 2006, American Ceramic Society. Parts (C) and (D): Reproduced with permission. [ 105 ] Copyright 2007, Nature Publishing Group.
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made up of silica. The silica cell walls have distinctive nanos-tructured patterns, which can be hexagonal, rod-shaped, or circular, depending on the species. [ 106 ] Diatoms can be used as templates in the same manner as silica colloidal crystals. Typi-cally, diatom templates are fi rst impregnated with an inorganic precursor, which can be hardened through gelation or condensa-tion to form a strong network around the template. This process has been employed to form thin layers of potential anode mate-rials, such as TiO 2 [ 107 , 108 ] or SnO 2 [ 109 ] on a silica diatom struc-ture. TiO 2 coatings were obtained by layer-by-layer growth of 10 nm TiO 2 colloids [ 107 ] or by atomic layer deposition [ 108 ] on the silica surface. SnO 2 coatings required amplifi cation of surface hydroxyl groups on silica diatoms by a thin interlayer of acr-ylate functionalized with glucosamine, for improved adhesion of the conformal SnO 2 coating. [ 109 ] The silica scaffold can then be removed by etching with HF or hot NaOH or KOH solu-tion. This was shown, for example in the case of porous carbon templated by diatomaceous earth, in a preparation involving sucrose and sulfuric acid as the carbon source. [ 110 ] Silica extrac-tion by refl uxing with NaOH solution formed a porous carbon product containing both macropores (2.5 μ m average diameter) and mesopores (3.7–4.1 nm average diameters, depending on the number of impregnation steps.
2.5. Pseudomorphic Conversions
Although diatom-templating is a viable approach toward LIB electrode materials, reports of the electrochemical behavior of such materials are still scarce. [ 111 ] However, their use in a pseudomorphic conversion to produce porous silicon anodes shows some promise. A pseudomorphic conversion is a reac-tion that changes the composition of a material while pre-serving the general features of its shape. [ 105 , 112 ] Silica diatoms can be directly reduced to silicon when treated with strong reducing agents at high temperatures. This concept builds on processes used in the semiconductor industry that rely on converting silica stock into silicon. However, it is chal-lenging to maintain the nanostructures of diatoms throughout the high temperature conversion step. The two traditional methods to produce silicon from silica are carbothermal reduction [ 113 ] and electrochemical reduction in molten salts or molten metals. [ 114 ] The former method requires temperatures higher than the melting point of silicon at which temperature the nanostructure is destroyed, and the latter is affected by the capillary forces of the hot liquid, which also causes collapse of silicon nanostructures. To avoid collapse due to capillary forces, a gas phase magnesiothermic reaction at relatively low temperatures may be used instead. For example, Bao et al. [ 105 ] converted nanostructured silica microshells of diatoms into nanostructured silicon with magnesium gas at 650 ° C, fol-lowing the reaction
2 Mg(g) + SiO2(s) → 2 MgO(s) + Si(s)
In this case, the resulting silicon network retained the inter-connected 3D structures of the original diatoms after magne-sium oxide was etched away with HCl solution (Figure 5 C, D).
When applied to porous solids with smaller feature sizes, such as mesoporous SBA-15, the magnesiothermic reaction
only partially preserves features of the silica preform. [ 115 ] Whereas the Si product retained the external morphology of the original SBA-15 particles, internal porosity in the product was mainly derived from textural mesopores, i.e., void spaces between 20–50 nm silicon grains. In another study, the mes-oporous Si anode obtained by a magnesiothermic reduction of mesoporous SBA-15 silica consisted of nanoparticles that surrounded mesoporous channels, indicating that parts of the SBA-15 structure were preserved. [ 116 ] The differences in prepa-ration techniques were small: an open corundum boat was used in the fi rst study and a sealed stainless steel ampoule in the second. In principle, this is an effective method for fabricating nanostructured silicon anode materials. However, it is very sen-sitive to reaction conditions, and surface oxidation of the high surface area silicon must be avoided to prevent the formation of a thin layer of nonconductive silica.
In a different type of pseudomorphic conversion, mesopo-rous Co 3 O 4 nanobelts were formed by thermal decomposition of α -Co(OH) 2 nanobelts. [ 117 ] These nanobelts possess a layered hydrotalcite structure and are ca. 20–50 nm thick, 200–500 nm wide, and tens of μ m long. The reaction involved a topotactic chemical transformation that introduced pores into the belt structures without changing the general external shape of the nanoparticles. Pore sizes were controllable by the reaction tem-perature: pores with an average diameter of 15 nm resulted from volume shrinkage during calcination at 450 ° C; mate-rials calcined at 350 ° C contained 6 nm pores. These nanobelts exhibited electron diffraction patterns typical of single crystals.
2.6. Non-Templated Porous Materials
The main focus of this review is on porous battery materials with well-controlled pore sizes or architectures, which are best achieved by using templating methods. However, numerous template-free methods of producing porous electrodes are
Figure 6 . Mesoporous anatase TiO 2 microspheres. (A)–(C) Spheres composed of isotropic TiO 2 nanoparticles. [ 124 ] (A) SEM image of an individual mesoporous microsphere. (B) The microsphere is composed of aggregated nanosized crystallites. (C) Schematic of the structure within each microsphere, showing the aggregate of TiO 2 nanocrystallites (large spheres) and Li ions (smaller spheres) after infi ltration with electrolyte. (D)–(F) Similar images for micro-spheres composed of anisotropic TiO 2 nanoparticles. [ 120 ] Parts (A)–(C): Reproduced with permission. [ 124 ] Copyright 2011, American Chemical Society. Parts (D)–(F): Reproduced with permission. [ 120 ] Copyright 2011, American Chemical Society.
available, including electrodeposition, [ 72 ] ultrasonication, [ 118 ] intercalation, [ 119 ] hydro- or solvothermal syntheses [ 120 , 121 ] and more. While these tend to produce materials with broader pore size distributions, some of the methods are simpler than those requiring templates, and they can yield materials with features and critical dimensions that ben-efi t the performance of an electrode. We will highlight only a few template-free approaches here, covering both the mesopore and the macropore range.
Aerogel syntheses are well suited to pro-duce mesoporous materials. They have been widely applied to vanadium oxide systems. V 2 O 5 ·xH 2 O aerogels are synthesized by sol-gel methods employing supercritical drying, freeze drying, or solvent exchange methods to maintain open pores after solvent removal by minimizing damaging capillary forces. [ 66 , 122 ] The resulting structures are three-dimensional networks of a solid phase surrounded by a continuous, random mesopore or macropore
network. In typical vanadia systems, the solid phase consists of ribbons composed of nanocrystalline nanoparticles (10–50 nm) whose crystalline phase and dimensions can be controlled through subsequent annealing in various atmospheres. [ 122 ] In these thermal treatments the atmosphere also determines the oxidation state of the vanadium oxide product. For example, the annealing of V 2 O 5 · xH 2 O aerogels at 400 ° C under high vacuum produced the vanadia bronze, VO 2 (B); V 6 O 13 was formed under Ar; and in air, the product was V 2 O 5 . [ 122 ] However, sin-tering during thermal treatment can result in densifi cation of the material and reduce the surface areas and pore volumes. It is possible to incorporate conductive components like single-walled carbon nanotubes (SWNT) during the synthesis of the aerogel. [ 123 ] In the case of a vanadia aerogel/SWNT composite, the SWNT formed ca. 10 nm diameter bundles intertwined between 20–30 nm vanadia ribbons. The two phases formed good electrical contact while maintaining suffi cient room for electrolyte access.
Several template-free syntheses involving hydrothermal, sol-vothermal, or simple hydrolysis reactions have been employed to produce mesoporous microspheres. The spherical mor-phology facilitates higher packing densities in an assembled electrode than individual nanoparticles. In most cases, the microspheres themselves consist of aggregates of nanoparticles either throughout the microsphere or in a shell layer around a hollow core. In either case, textural mesoporosity is present throughout the nanoparticle aggregates.
A relatively simple procedure to synthesize nanocrystalline anatase TiO 2 mesoporous microspheres involved the hydrol-ysis of tetrabutyl titanate in acetone followed by drying and annealing at 400 ° C. [ 124 ] Although the resulting microsphere size distribution was broad (200–1000 nm), the structure of the microspheres was suitable for electrolyte infi ltration, given that mesopores ( ∼ 5 nm average pore diameter) formed between aggregates of 8–15 nm anatase nanoparticles within the spheres ( Figure 6 A–C).
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A hydrothermal synthesis was applied to the production of mesoporous anatase TiO 2 microspheres. [ 120 ] The spheres were composed of self-assembled anisotropic (ca. 20 nm × 50 nm) nanocrystals whose aggregation results in textural mesopo-rosity with average mesopore diameters of ∼ 14 nm (Figure 6 D–F). Although no post-calcination step was used in this syn-thesis, this spherical particle morphology was well maintained during electrochemical cycling in a diluted ionic liquid electro-lyte. However, in other organic electrolytes, the particles aggre-gated, resulting in larger capacity losses over multiple cycles. Alternatively, solvothermal syntheses are suitable for nonoxide materials. Mesoporous Cu 2 SnS 3 spheres were prepared by sol-vothermal synthesis from tin and copper chlorides in polyeth-ylene glycol, with thiourea as the source of sulfur. [ 121 ] The poly-ethylene glycol solvent also helped to link nanoscale clusters of the precursor components during the reaction.
Carbon-containing precursors may be mixed in with the precursors for the target electrode material to obtain hollow microspheres by a different mechanism than those discussed in Section 2.3. In a synthesis of mesoporous TiO 2 microspheres by a hydrothermal reaction, resorcinol (R) and formaldehyde (F) were incorporated in the aqueous precursor solution con-taining the titanium salt. [ 125 ] During the hydrothermal reaction at 85 ° C, the RF precursors formed polymeric cores around which titania assembled and crystallized. Subsequent calci-nation in air yielded hierarchically structured hollow micro-spheres with an external diameter of a few micrometers. The product texture depended on the calcination temperature. For a sample calcined at 400 ° C, the microsphere walls were com-posed of smaller aggregated spheres with mean diameters around 60 nm. Textural mesopores formed channels between the spheres. The 60 nm spheres were themselves composed of TiO 2 nanoparticles (ca. 6 nm) and contained 3–4 nm mes-opores. Calcination at 550 ° C increased the mesopore size to 6 nm but reduced the overall surface area and pore volume. The microsphere size and shell thickness were functions of the hydrothermal reaction time, increasing with longer times. A similar approach involving hydrothermal carbonization of fur-fural in the presence of SnO 2 nanoparticle sols was employed to synthesize mesoporous SnO 2 microspheres. [ 126 ] The furfural components formed carbon spheres with a hydrophilic surface that attracted the SnO 2 nanoparticles. During carbon removal, these nanoparticles (each 3–7 nm in diameter) assembled, forming mesoporous microspheres with an average mesopore diameter of 10 nm.
Depending on experimental conditions, the carbon-con-taining additives may lead to macroporous structures rather than hollow spheres. In a synthesis of Li(Ni 1/3 Co 1/3 Mn 1/3 )O 2 from acetate precursors of the metals dissolved in an aqueous solution that also contained resorcinol and formaldehyde, RF polymer was interspersed throughout the electroactive phase. [ 127 ] Calcination at 950 ° C removed the polymer phase, leading to formation of disordered macropores surrounded by 500–1000 nm diameter particles of Li(Ni 1/3 Co 1/3 Mn 1/3 )O 2 . Although not mentioned in the publication, it is possible that the resorcinol-formaldehyde gel component helped achieve compositional homogeneity in this multicomponent electrode material, similar to the role of complexing agents in Pechini-type syntheses. [ 128 ] Irregular macropores were also generated in
an SnP 2 O 7 electrode material where sucrose was added to the reaction mixture before the hydrothermal synthesis. [ 129 ]
2.7. Electrode Materials with Hierarchical Porosity
Hierarchy in pore sizes permits better permeation of an electro-lyte through a porous electrode. Secondary mesopores within macropore walls further reduce ion diffusion paths through the electrode. They increase the interfacial area between the electro-lyte and the electrode. They also provide a way of hosting active material, an electrically conductive component, or a stabilizing material within the walls of a porous matrix while still main-taining access for the electrolyte phase.
In general, hierarchical macroporous-mesoporous materials can be produced either by the introduction of macropores into mesoporous materials (e.g., by etching) or by the introduction of mesopores into macroporous materials (e.g., by dual tem-plating). Depending on the approach, the distribution of the pores may be random or ordered in the porous substrates. If mass transport is the more important factor, it is preferable to have an ordered, interconnected macropore system with addi-tional mesopores introduced into the walls. This is usually the case when electrolytes with ions having a small diffusion coef-fi cient or viscous solvents are used (ionic liquids or high molec-ular weight solvents). An ordered, interconnected macropore system with low tortuosity facilitates electrolyte diffusion in and out of the active materials. If mass transport is less impor-tant, then the macropores can be randomly distributed in the mesoporous substrate, but now they only function as inner res-ervoirs supplying electrolyte to the surrounding mesopores.
Several of the structures discussed above possess porosity on multiple length scales if one considers textural mesoporosity between nanocrystals as one of the length scales. However, it is possible to introduce hierarchical porosity with greater control over pore architecture by employing multiple hard and/or soft templates in a synthesis. The procedures are similar to those used to prepare macroporous materials, except for the fact that surfactants, supramolecules, or nanoparticles are added into the precursors before the infi ltration of the macropore templates. Methods of synthesizing materials with hierarchical porosity have recently been reviewed, [ 130 ] and here we will provide some representative examples relating specifi cally to LIBs. Additional examples describing composite electrodes built around materials with hierarchical porosity are found in Sections 3.8 and 3.9.1.
3DOM Sn/Ni alloy fi lms with complex hierarchical structure were fabricated by depositing polystyrene spheres within pho-toresist patterns on a copper substrate, annealing the spheres to stabilize the colloidal crystals, and electroplating a Sn-Ni alloy within the colloidal crystals ( Figure 7 A–C). [ 131 , 132 ] The resulting micro-patterned macroporous Sn-Ni alloy anode exhibited a high areal capacity (i.e., a high capacity per unit area). The areal discharge capacity increased with the open pore ratio of the photoresist substrate (cylindrical pattern < square pat-tern < hexagonal pattern). Another interesting example of pat-terning a potential electrode material with hierarchical porosity involved dip-coating of a TiO 2 precursor/surfactant solution in a lithographically patterned, microstructured photopolymer resin. [ 133 ]
Figure 7 . Electrode structures with hierarchical porosity. (A)–(C) SEM images of patterned 3DOM Sn-Ni alloy fi lms at various magnifi cations, showing the lithographic patterns and the colloidal-crystal-templated macropores. [ 132 ] (D)–(F) Mesoporous titania nanotubes templated in an AAO mem-brane. [ 134 ] (D) SEM image of nanotube array. (E) TEM image showing the hexagonally ordered mesopore structure in the walls of the hollow nanotubes. (F) Schematic diagram of the transport path of lithium ions and electrons in the mesoporous titania nanotubes. Parts (A)–(C): Reproduced with permission. [ 132 ] Copyright 2010, Elsevier Ltd. Parts (D)–(F): Reproduced with permission. [ 134 ]
Figure 8 . TEM images of nanoporous carbon materials with hierarchical pore structure prepared by (A) nanocasting, [ 138 ] (B) a triconstituent syn-thesis, [ 139 ] or (C-D) a direct synthesis. [ 140 ] Macropores arise from col-loidal crystal templating. Mesopores have wormlike geometries in (A) and (B), cubic symmetry in (C) and 2D hexagonal symmetry in (D). Part (A): Reproduced with permission. [ 138 ] Copyright 2006, American Chemical Society.
Templating in anodic aluminum oxide (AAO) fi lms or mem-brane fi lters is a well developed technique to prepare nanotube structures, [ 135 ] including those of interest for high rate LIBs. [ 136 ] This type of templating can be combined with surfactant templating to obtain nanotubes containing multiple pore sizes. Titania nanotubes with mesoporous walls were grown within AAO membranes, employing an alkoxide precursor of titanium and the P123 triblock copolymer as the structure-directing agent. [ 134 ] Hollow tubes were formed as solvent evaporated due to an affi nity between the gel and the hydrophilic alumina walls, causing the gel to shrink in a direction perpendicular to the pore channels. After calcination, the membrane was removed by dissolution in NaOH solution. Arrays of well-aligned nan-otubes ( ∼ 200 nm outer diameter) formed after treatment in supercritical CO 2 (Figure 7 D–F). Without the supercritical drying step, surface tension forces between nanotubes caused the nanotubes to aggregate and entangle. The hierarchical pore structure facilitated the distribution of the electrolyte through ca. 140 nm channels and ∼ 7.5 nm mesopores, leading to excel-lent high rate performance (150 mAh/g at 240C rate) and good cycle performance.
The most common examples of hierarchical electrode tem-plating with colloidal crystals involve carbon systems, although some oxide materials with hierarchical pore structure, such as Fe 2 O 3 have also been fabricated by this approach. [ 137 ] Carbon structures can be prepared by nanocasting with hierarchically structured 3DOM/m silica prepared using both surfactants and polymeric colloidal crystals ( Figure 8 A). [ 138 ] They can also
be synthesized by more direct methods, employing a colloidal crystal template with either a triconstituent precursor (silicon alkoxide, resol, block-copolymer surfactant, Figure 8 B) [ 139 ] or a simpler mixture of phenol-formaldehyde with a block-copolymer surfactant. [ 140 ] The last approach avoids the use of caustic chem-icals for silica etching and provides some control over the mes-opore geometry (2D hexagonal vs. cubic mesopores surrounding macropores, Figure 8 C and D). The triconstituent method, on the other hand, produces the largest mesopores and provides a bimodal mesopore size distribution in the carbon skeleton around the macropores. The larger mesopores make the tri-constituent-derived or the nanocast monoliths most suitable for preparations of nanocomposite electrodes in which a precursor for active material is imbibed in the carbon host. [ 141 , 142 ] All three of these 3DOM/m carbon materials, when used as monolithic electrodes in a lithium battery, provide improved rate capabili-ties compared to 3DOM C lacking the templated mesopores.
2.8. Choice and Stability of Crystallographic Phases
The ability to control crystallographic phases of active elec-trode materials during the synthesis is critical to the elec-trochemical performance of the electrodes. In the context of porous electrodes, a few points should be considered here. The crystallographic phases obtained in the syntheses described above depend, in part, on the precursors, thermal history, and materials dimensions. Because critical dimen-sions of active material in most of the systems described here are on the nanometer length scales, the observed crys-tallographic phases are also those typically found for nano-particles. For example, in the case of 3DOM TiO 2 calcined at 400–575 ° C, the skeletal walls are composed of anatase nano-particles, [ 143 ] even though rutile is thermodynamically more stable in the bulk form.
With soft templates and even with polymeric hard templates, the range of calcination temperatures needed to achieve a specifi c crystallographic phase is limited by grain growth of the active material. For example, for active materials requiring high tem-perature syntheses above ca. 550 ° C, excessive grain growth often results in the loss of the templated structure. As noted earlier, this effect can be somewhat mitigated by employing templates that provide thicker walls of active material (Section 2.2.2). [ 64 ] Alternatively, nanocasting with hard silica or carbon templates provides access to high temperature phases, and is probably the most versatile approach (see Section 2.2.1), as long as these tem-plate materials do not affect the target composition as a result of a reaction during thermal processing or during template removal.
In syntheses of multicomponent electrode materials, impu-rity control and phase control are key items during the syn-thesis. For example, the preparation of 3DOM LiNiO 2 required excess lithium salt to compensate for loss of lithium during calcination and also required a high oxygen partial pressure to suppress the reduction from the Ni 3 + to the Ni 2 + state. [ 64 ] Pechini-type approaches that involve complexation and polym-erization of metal precursors permit precise control over product stoichiometry in templated, multicomponent porous materials [ 144 ] and are also useful for syntheses of complex elec-trode materials.
3. Electrochemical Properties of Porous Electrodes for LIBs
Given the large array of techniques available for synthesizing porous electrodes, the question naturally arises: which tech-nique is most suitable for LIB applications? To help answer this question, it is necessary to consider the infl uence of the structural features of porous electrodes on their electrochem-ical properties. For high-power applications, such as applica-tions in EVs and HEVs, it is important to increase the power ( P ) delivered by the battery, which depends on the voltage of the battery ( V ) and the current delivered by the battery ( I ) fol-lowing the relationship P = IV . With a load, the actual voltage of the battery is decreased from the open-circuit voltage (V oc ) according to V = V oc –IR b , where R b is the internal resistance of the battery. The internal battery resistance itself is a sum of several resistive components, namely R b = R el + R A + R C + R in (A) + R in (C) + R c (A) + R c (C) . [ 145 ] Here R el is the electrolyte resistance, R A and R C the resistance of anode and cathode, respectively, R in (A) and R in (C) the resistance to transport of the working ion (here Li + ) across the electrolye-anode and electrolyte-cathode interfaces, respectively, and R c (A) and R c ( C) are the current-collector resistances of each electrode, involving resistance to electron transfer from the electrode to the current collector and electron transport through the current collector. When the battery power is written in the alternate form P = V 2 /R b , it is clear that for high power appli-cations, it is critical to keep R b to a minimum. Considering each term separately, several parameters can be identifi ed that may benefi t from nanoporous electrode architectures. The electrolyte resistance to lithium ion transport, R el = L / σ i A decreases when the thickness of the electrolyte layer in the interelectrode space ( L ) is kept short and the geo-metrical area of this space ( A ) is large. The ion conductivity σ i increases with the diffusion coeffi cient of Li + through the electrolyte but decreases with greater tortuosity of the ion diffusion path, the latter of which depends on the pore geometry. The electrode resistances are functions of the elec-trode material and contact between particles making up the electrode material. In nanoporous solids interparticle contact can be quite good and nanocomposite electrode structures with an added highly conductive phase can further reduce the electrode resistance. The ion transfer resistance R in decreases with increasing interfacial area between the elec-trode and the electrolyte, which can be very high in nano-porous materials. Finally, low current collector resistances are possible if highly conducting current collectors can be placed in intimate contact throughout a porous electrode. In this section, we will consider the infl uence of structural fea-tures of porous electrodes on electrochemical properties of interest for LIBs.
3.1. Pores Provide Access of the Electrolyte to the Electrode Surface
A porous solid can be regarded from the point of view of the pores, the solid surrounding the pores, and the inter-face. Among the terms contributing to the internal battery
resistance, the pores control mostly the internal electrolyte resistance. For cylindrical pores, this resistance is Rel = ρl
3nπr 2
, where ρ is the resistivity of the electrolyte, l the pore depth, r the pore radius, and n the number of pores included in the area of a porous electrode. [ 146 ] In terms of designing the pore architecture of an electrode, this relationship suggests that the electrolyte resistance can be reduced by employing shorter pores (small l ) and larger pore radii (large r ). In general, pores in an electrode facilitate electrolyte penetration and provide diffusions paths for rapid Li + ion transport through the elec-trolyte, particularly if the pores are well interconnected and the pores are wetted by the electrolyte. Although it is diffi cult to isolate the effects of pore size completely from effects of wall thickness, some general trends have been observed, as illustrated below.
3.1.1. Effects of Pore Size
On one hand, electrolyte species cannot move effectively through pores that are too small, on the other hand, volumetric energy density is lost if the pore radius becomes too large. Thus an optimal pore size range needs to be identifi ed. On the smaller end of pore sizes, micropores provide higher surface areas but access for the electrolyte is typically too restricted. Furthermore, in micropores the Debye length of a typical elec-trolyte solution used in Li-ion batteries (ca. 0.1–0.3 nm in 1 M Li + solutions, up to 1 nm for 0.1 M solutions) [ 147 ] is on the order of the pore size. In this pore size regime, the classical descrip-tion of the double layer no longer applies and the surface may not be electrifi ed, so that Li-ion transport would be infl uenced in a negative way. Hence a large micropore concentration limits conductivity. [ 148 ]
For the larger pore size regime, an early study of macropo-rous MnO 2 cathodes in Li/MnO 2 cells already showed that the cathode utilization was sensitive to the cathode pore diameter and the volume fraction of macropores > 200 nm. [ 149 ] Cathodes with higher porosity exhibited higher electrode utilization at high discharge rates. More recent investigations refl ect this pattern. For example, a two-dimensionally ordered macropo-rous NiO fi lm (ca. 500 nm pores) on a stainless steel substrate exhibited a high rate capacity (at 15C) that was ca. 2.7 times higher than that of a nontemplated NiO fi lm. [ 85 ] In both cases the fi lm was composed of fl ake-like particles with thicknesses of ca. 12–20 nm. It was suggested that the stable open structure facilitated electrolyte transport and led to the improved high-rate performance.
In the case of a 3DOM Sn-Co alloy fi lm electrode that was electroplated within a colloidal crystal template supported on a Ni-coated Cu sheet, better cycle performance was observed for materials with smaller macropores (180 nm vs. 500 nm). [ 72 , 150 ] Although this effect was not explained, the data suggest that the mesopore-to-small-macropore size range may be most effective for achieving high rate capabilities without wasting volume in porous electrodes. The same conclusion can be made from a study of 3DOM Li 4 Ti 5 O 12 that contained macro-pores derived from a colloidal crystal template and textural mesopores between crystallites in the wall skeleton. [ 80 ] In this study, the effect of fi lling fraction of precursor in the template was evaluated. This is equivalent to changing the wall thickness
in the material. At a given calcination temperature, the number of crystallites and number of contact points between them increased as the fi lling fraction was increased. As a result, inter-particle porosity decreased and the 3DOM material behaved more like a nonporous ceramic with diminished rate perform-ance. These observations indicate that the mesoporosity is more important than the macroporosity for improving the high rate behavior. The inclusion of larger macropores is mainly advanta-geous for cells using viscous electrolytes, salts with small diffu-sion coeffi cients, or for the construction of three-dimensionally interpenetrating electrode structures (see Section 4).
A few systematic studies of the infl uence of mesopore size on specifi c capacities have been reported. In a study of mes-oporous anatase TiO 2 prepared with three different cationic surfactants, including C 8 TA + (5.7 nm pores), C 12 TA + (6.1 nm pores), and C 16 TA + (7.0 nm pores) the specifi c capacity increased with increasing mesopore size from 232 to 242 to 268 mAh/g at C/5 rate. [ 17 ] The specifi c capacity values also correlated with the surface area of the mesoporous TiO 2 samples (90, 109, and 135 m2/g, respectively). However, in the case of mesoporous rutile TiO 2 templated with anionic surfactants of different chain lengths (C 10 H 21 SO 3 Na, pore size 2.2 nm, surface area 266 m 2 /g; C 12 H 23 SO 3 Na, 3.1 nm, 300 m 2 /g; and C 16 H 31 SO 3 Na, 3.8 nm, 245 m 2 /g), the highest capacity was instead observed for the material having the highest surface area, i.e., the one templated with the C 12 -surfactant. [ 18 ] The infl uence of porosity on rate capability was also studied for mesoporous LiFePO 4 colloidal particles with spherical shape. [ 148 ] In this system, as the mes-opore volume was reduced, the discharge capacity decreased, particularly at higher rates.
Studies of geometric effects of mesoporous electrodes on the rate performance using nanocast electrodes led to the conclu-sion that the rate of intercalation depends on pore size but even more critically on the pore wall thickness. [ 50 ] A challenge in this kind of study is the fact that mesopore diameters cannot always be varied smoothly through the control of reaction conditions, and in some cases, multiple pore size ranges may be present in one material. In the case of mesoporous β -MnO 2 synthesized by hard templating with KIT-6, the rate capability decreased with increasing pore wall thickness from 5 to 8.5 nm. [ 50 ] For materials with similar wall thickness and bimodal mesopore distributions, the rate capabilities depended on the ratio of pore volumes from large mesopores (11 nm pores) to smaller mesopores (3.3–3.5 nm). The exact dependence appeared to be complex, but the effect of pore size was signifi cant for current densities in the range 100–1500 mAh/g ( Figure 9 ). Larger mes-opores provided improved Li + transport in the electrolyte phase that occupied the pores. They may also have acted as a reservoir for lithium ions.
In a system with disordered mesopores, mesoporous anatase TiO 2 synthesized via a sol-gel reaction in a bicontinuous emulsion, a clear difference in rate capability was observed for samples con-taining a mean mesopore size of 17 nm (crystallite size 14 nm) or 2.8 nm (crystallite size 5 nm). [ 26 ] At rates below 10C the sample with smaller pore- and crystallite sizes maintained higher capaci-ties, but it quickly lost capacity at higher rates. The capacity loss at rates above 10C was much less for the sample containing larger pores and crystallites. The limitation at high rates thus appears to be diffusion of lithium through the electrolyte.
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Figure 9 . First discharge capacity for different mesoporous β -MnO 2 elec-trodes as a function of rate. [ 50 ] Capacity is expressed as a percentage of the discharge capacity at 15 mA/g. The TEM images show the structures of three of the samples. Numbers in parentheses correspond to the ratios of large (11 nm) to small (3.3–3.5 nm) pores volumes. Information about the other samples can be found in the original reference. Reproduced with permission. [ 50 ] Copyright 2010, American Chemical Society.
3.1.2. Effects of Uniformity and Tortuosity of Pore Space
An important criterion for effi cient electrolyte distribution throughout a porous electrode is the need for interconnectivity between pores. This is the case, for example, in materials nano-cast with KIT-6. [ 52 ] Beyond this criterion, uniform pore sizes and control over these pore sizes permits optimization of the open volume for best mass transport without wasted space. Several advantages of ordered mesoporous electrodes have been noted in comparison to analogous electrodes formed from discrete nanoparticles. For example, in the case of mesoporous anatase TiO 2 , the packing effi ciency was greater compared to anatase nanoparticles, the uniformity and continuity of pores was greater, and less conductive carbon was needed to maintain pathways for electron transport. [ 47 ] An impressive high rate per-formance of 125 mAh/g at 12 A/g was observed and was higher than for anatase nanoparticles in electrodes prepared with sim-ilar carbon loadings. The rate capability was even higher than for a disordered mesoporous titania system whose pores were fi lled with conductive RuO 2 . [ 151 ] This observation led to the con-clusion that electron transport may not be the only rate-limiting factor in these materials, [ 47 ] but electrolyte transport is also important. In this and similar hierarchically structured systems (electrodes prepared from mesoporous hollow TiO 2 micro-spheres, [ 125 ] titania nanotube arrays with mesoporous walls [ 134 ] ), a hierarchy in pore structure permitted highly effi cient distribu-tion of electrolyte while maintaining short ion diffusion paths through the solid phase.
Electrolyte resistance is lowest for pores that offer the most direct path for lithium ions, i.e., those that have the lowest tortuosity. Tortuosity is defi ned by T =
(�ρmρb
)12 , where Φ is the
volume fraction occupied by the electrolyte in the porous elec-trode, ρ m is the resistivity of electrolyte through the porous electrode, and ρ b the bulk resistivity of the electrolyte. [ 65 ] Elec-trode polarization, i.e., an internal voltage drop resulting from limited ion mass transport increases with the square of tortu-osity, according to �Ei R = ilρb T2
thickness of the electrode, and A the area of the electrode. [ 65 ] Periodically structured porous materials with interconnected pores normally have lower tortuosity than texturally porous materials. The resulting effect on electrode polarization was studied in the case of a 3DOM vanadia aerogel on an indium-tin-oxide substrate (800 nm macropores, ∼ 150 nm-thick walls). [ 65 , 66 ] The high symmetry and high pore connectivity (12 windows surround each macropore) of the 3DOM architec-ture provided low tortuosity throughout the pore system and led to less polarization at higher discharge rates compared to bulk aerogel electrodes.
3.2. The Large Surface Area Facilitates Charge Transfer Across the Electrode/Electrolyte Interface
Porosity imparts materials inherently with a large specifi c sur-face area, i.e., a large area per unit mass. This surface area increases inversely with pore dimensions so that the largest gains are obtained for materials with micropores. However, to be useful, pores must be large enough that electrolyte can pen-etrate them. Major benefi ts of the large contact area between an electrolyte and the electrode are that the current density per unit surface area decreases, reducing electrode polarization, and that charge transfer at the interface is easier compared to bulk particles. For example, in the case of mesoporous NiO nanocast from KIT-6, AC impedance measurements revealed that the charge transfer resistance of mesoporous NiO was much lower than that of bulk NiO. [ 42 ] The activation energy for lithium ion intercalation was 20.8 kJ/mol for mesoporous NiO compared to 45.0 kJ/mol for bulk NiO.
However, the larger surface areas can also lead to unwanted side reactions involving electrolyte decomposition, particu-larly at the anode side. At the anode, higher surface areas may enhance the formation of a solid-electrolyte interface (SEI) layer. Mesoporous Cr 2 O 3 anodes nanocast from KIT-6 silica are an example where mesoporosity results in signifi cant loss of avail-able lithium ions during the early cycles. [ 53 ] The high surface area of this material favors the formation of electrolyte decom-position products, leading to large excess uptake of lithium in the fi rst discharge step, and this lithium is subsequently not released. As a consequence, the irreversible capacity of the fi rst cycle was relatively large for the mesoporous Cr 2 O 3 electrode, the capacity decayed rapidly during the fi rst 15 cycles (although it recovered slightly thereafter), and polarization changed upon cycling. In bulk Cr 2 O 3 , the electrolyte decomposition products (EDP: lithium carbonate together with organic components) tend to dissolve during charging, but in the mesoporous mate-rial signifi cant amounts remained trapped in the channels. Because the EDP layer fi lled the mesopores, it prevented the collapse of the mesopores during cycling. [ 53 ] Therefore, over all cycles, the capacity of the mesoporous material exceeded that of bulk particles and reversibility was better sustained. A similar stabilization effect was not observed in the case of mesoporous Co 3 O 4 nanocast from KIT-6. [ 54 ] This high surface area electrode also favored the formation of a polymer layer, which built up over multiple cycles. However, after ca. 50 cycles, the mes-ostructure had disappeared and turned into a nanocomposite of Co 3 O 4 and EDPs. Because the formation of an EDP layer
resulted in a large excess capacity in the fi rst cycle, the overall capacity exceeded the theoretical capacity of the actual active material. EDP or SEI layers must therefore be considered when interpreting results with unexpectedly high capacity values. Even macroporous electrodes, which have lower specifi c surface areas than mesoporous electrodes, face similar issues related to SEI interface formation. For example, 3DOM CoFe 2 O 4 with 130 nm macropores also exhibited a large irreversible capacity loss in the fi rst cycle. [ 82 ]
The EDP layer formation can result in a complex evolution of the specifi c capacity over multiple cycles, but this behavior depends on the material architecture. In mesoporous Co 3 O 4 nanocast from KIT-6 [ 41 , 54 ] or mesoporous nanowire bundles nanocast from SBA-15, [ 41 ] a large excess capacity was observed in the fi rst cycle, dropped in the second cycle, built up again over the next 20 cycles, and then gradually dropped over the following cycles ( Figure 10 A–C). This complex cycling behavior was not observed in mesoporous Co 3 O 4 nanobelts, which main-tained a relative constant capacity for cycles 2–20, after a steep drop following the fi rst cycle (Figure 10 D–F). [ 117 ] In this case, EDPs were not trapped in mesopores and were less likely to accumulate in mass.
3.3. Thin Walls Provide Short Path Lengths for Lithium Ion Dif-fusion in the Solid Phase
As mentioned above, rates of lithium ion intercalation in porous electrodes depend on both pore size and pore wall
Figure 10 . Infl uence of EDP layer trapping on electrochemical cycling behfi cation TEM image. (B) High-magnifi cation TEM image. (C) Reversible liMesoporous Co 3 O 4 nanobelts. [ 117 ] (D) Low-magnifi cation SEM image. (E) Hprepared at two different calcination temperatures. Reproduced with perm
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thickness. [ 50 ] The characteristic time constant for intercalation is τ ∼ l 2 / D , where l is the diffusion length through the solid and D the diffusion coeffi cient for lithium ions in the solid. Therefore, reducing the particle dimensions from the microm-eter range to the nanometer range reduces the time for lithia-tion or delithiation by a factor of 10 6 , so that it is possible to charge electrodes much faster or deliver larger currents during discharge. The advantages of nanostructure on rate capabilities of LiMn 2 O 4 cathodes were demonstrated quite some time ago in the case of nanotubule structures. [ 152 ] In that case, the rate capability improved with decreasing wall thickness of the nano-tubules. Similar advantages are found in porous nanostruc-tures. For example, in hollow core-shell structures of Li 4 Ti 5 O 12 with 100–200 nm thick shells, higher discharge capacities were observed at all rates for the thinner shells. [ 88 ] In 3DOM anatase fi lms, the specifi c capacity and rate capability improved when spherical templates with smaller diameters were used (100 nm) because these produced 3DOM products with thinner walls and larger specifi c surface areas. [ 78 ]
It should be noted that thinner walls are advantageous at high charge/discharge rates, but not at low rates, when kinetic limitations are less important. For example, at very slow lithiation rates, micrometer-sized bulk LiCoO 2 particles were observed to have greater storage capacities than colloidal crystal templated macroporous LiCoO 2 . [ 69 ] However, at higher cycling rates the macroporous electrodes with 70–100 nm thick walls exhibited marked improvements in specifi c discharge capaci-ties compared to bulk materials by reducing the solid-state diffusion distances through the electrode particles. A similar
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avior. (A)–(C) Mesoporous Co 3 O 4 , nanocast from KIT-6. [ 41 ] (A) Low-magni-thium storage capacity vs. cycle number at different current rates. (D)–(F)
igh-magnifi cation SEM image. (F) Capacity vs. cycle number for nanobelts ission. [ 41 , 117 ]
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Figure 11 . TEM micrograph of 3DOM Li 4 Ti 5 O 12 illustrating low angle grain boundaries between crystallites that make up the solid skeleton. Reproduced with permission. [ 80 ] Copyright 2006, American Chemical Society.
pattern was observed for mesoporous β -MnO 2 synthesized by hard templating with KIT-6, where the rate capability increased with decreasing pore wall thickness from 8.5 to 5 nm, but at low rates, the fi rst discharge capacity increased with increasing wall thickness. [ 50 ]
3.4. The Small Feature Sizes Permit Increased Utilization of Active Material, Especially at High Charge/Discharge Rates
The extent of intercalation, utilization of active material and, therefore, specifi c capacity, can increase markedly for elec-trodes with nanometer-sized features. For instance, bulk rutile can accommodate less than 0.1 Li per TiO 2 at room tempera-ture, but this value increases above 0.5 Li per TiO 2 for rutile nanoparticles. [ 153 ] For mesoporous rutile, too, a reversible capacity of 0.55 Li has been reported, which corresponds to 185 mAh/g at C/5 rate. [ 18 ] Bulk crystalline β -MnO 2 (tetragonal rutile structure) with micrometer-sized particles does not inter-calate much, if any, lithium through its narrow tunnel struc-ture (2.3 Å × 2.3 Å). [ 154 ] In contrast, mesoporous β -MnO 2 with similar external particle sizes exhibited a reversible capacity of 284 mAh/g, corresponding to the composition Li 0.92 MnO 2 ). [ 50 ] Superior rate capabilities were also demonstrated for nanocast mesoporous Li 1 + x Mn 2-x O 4 with spinel structure compared to an electrode prepared from bulk particles. [ 52 ]
Many other examples of improved specifi c capacities of porous electrodes compared to their nonporous analogs for LIBs have been reported in the literature. For example, a 2.2 μ m-thick macroporous LiMn 2 O 4 fi lm electrodeposited on a conducting substrate with polymer sphere templates exhibited a higher specifi c capacity than a nonporous fi lm with the same thickness. [ 74 ] The increased capacities are especially evident at high rates. Macroporous Li(Ni 1/3 Co 1/3 Mn 1/3 )O 2 maintained a capacity of 173 mAh/g at 12C. [ 127 ] Mesoporous vanadia deliv-ered 125 mAh/g at 50C, whereas a nonporous control sample provided only 75 mAh/g at the same high rate. [ 29 ] Mesoporous Li 4 Ti 5 O 12 synthesized by surfactant templating maintained a high capacity of ca. 150 mAh/g, even at the high rate of 64C, with excellent cycle stability. [ 32 ] The rate performance of a sur-factant-templated mesoporous nanocomposite anode consisting of a crystalline phase of active material (3–5 nm anatase parti-cles, 85 wt%) embedded in a glassy P 2 O 5 phase and CuO or SnO 2 as conductive material was also impressive, starting at a specifi c capacity of 260 mAh/g at 10 A/g and maintaining ca. 190 mA/g after 200 cycles at this rate. [ 30 ] In some instances, additional capacity at high rates can be attributed to double-layer and pseudocapacitive processes in a mesoporous material, as was the case for mesoporous α -Fe 2 O 3 (hematite), where the observed storage capacity corresponded to more than one Li per Li x Fe 2 O 3 . [ 15 ]
In the comparisons of capacities brought up in this review, it is important to distinguish between specifi c capacities (or gravi-metric capacities) and volumetric capacities. The former refer to the amount of charge stored in a given mass of electrode, and the latter specify the amount of charge stored per unit volume of an electrode. Whereas many examples of porous electrodes with improved specifi c capacities compared to bulk materials exist, particularly at high rates of charge and discharge, porosity
usually reduces volumetric capacities. As a result, for electrodes with the same thickness, the total charge (and energy) stored will be lower for a porous material than a non-porous material. A few exceptions will be discussed in section 3.6.
3.5. Bicontinuous Walls and Pores Provide Continuous Trans-port Paths through the Active Phase (Walls) and the Electrolyte Phase (Pores)
In 3DOM structures templated by fcc colloidal crystals, all spheroidal pores are interconnected by windows, so that a con-tinuous void space permeates all of the solid mass. The walls surrounding the void space are themselves completely inter-connected. Hence, the structure is bicontinuous, providing continuous paths for charge transport through the electrolyte in the pores and through the active material in the walls. In the case of mesoporous electrodes templated from KIT-6, another cubic structure, a similar bicontinuous of pores and solid phase exists. This is not necessarily the case in SBA-15-derived materials, the structure of which templates parallel arrays of nanotubes. In some cases though, these nanotubes may be con-nected through thin, perpendicular linkers originating from micropores that connect adjacent mesopores in SBA-15.
For electrode applications, it is important that the skeleton in 3DOM structures and mesoporous electrodes consists either of a continuous amorphous phase or of a nanocrystalline phase in which low-angle grain boundaries minimize grain-grain inter-face problems. The latter case was demonstrated for 3DOM Li 4 Ti 5 O 12 , where low-angle grain boundaries between wall par-ticles produced improved conduction pathways ( Figure 11 ), thereby reducing polarization and enabling high rate cycling. [ 80 ] Similarly, in mesoporous TiO 2 , close contact between ana-tase grains provided a suffi cient conduction pathway to facili-tate Li + storage at low rates even without a conductive carbon additive. [ 17 ] For higher rates, however, a carbon additive was required.
Given the relatively high interfacial area between active material and electrolyte in 3DOM electrodes and the relatively short diffusion path lengths, the rate performance of 3DOM electrodes was improved compared to the nonporous counter-parts in most cases, but not always. In the case of 3DOM ana-tase TiO 2 thin fi lms prepared on conducting glass substrates by three different methods (electrodeposition, liquid-phase depo-sition, and vapor-phase deposition), all samples showed slug-gish lithium insertion/deinsertion compared to nontemplated anatase. [ 77 ] The poor performance was ascribed to the limited number of connecting points in a highly crystalline 3DOM net-work, which signifi cantly lowered the electrical conductivity of the electrodes.
The contact between grains depends on the sintering history of the material. In a comparison of various synthesis param-eters for mesoporous LiFePO 4 colloidal particles with spherical shape, the parameter with the greatest effect on electrochemical performance was the calcination temperature, which controlled grain size of LiFePO 4 in the colloidal spheres. [ 148 ] At higher temperatures, sintering of grains lowered the charge transfer resistance and yielded greater capacities at all rates.
3.6. Better Interparticle Contact Compared to Packed Nanoparti-cles and (in Exceptional Cases) Increased Volumetric Capacity
The volumetric capacity of an electrode is important in bat-tery applications where space is limited, for example, in port-able electronics and autonomous sensors. In optimizing volumetric capacity or energy density, one has to consider a trade-off between minimizing unused void space [ 155 ] and pro-viding access of electrolyte to the active material to achieve suffi cient ionic conductivity. [ 156 ] Void spaces are present even in bulk electrodes that are prepared from nonporous particles, resulting from interparticle porosity. Especially in the case of nanoparticles, the interparticle porosity can be high, leading to a low electrode density after compaction and ultimately to low volumetric capacity in practical Li-ion cells. Mesoporous elec-trode materials provide the necessary void space for electrolyte access to the surface, and, in spite of their porosity, they can possess high volumetric capacities (although this is still more the exception than the rule). A particularly striking example is meso-WO 3-x nanocast from KIT-6 mesoporous silica. [ 57 ] This electrode exhibited a specifi c capacity of 748 mAh/g, corre-sponding to 6.5 Li/W and a remarkable volumetric capacity of ∼ 1500 mAh/cm 3 , which approaches that of Li metal ( ∼ 2000 mAh/cm 3 ). Furthermore, the external dimensions of mesopo-rous and macroporous electrode powders are typically microm-eter sized, so that packing effi ciencies are similar to those of conventional electrode materials and much better than those of true nanoparticles. For instance, the packing density of surfactant-templated mesoporous TiO 2 was estimated to be 6.6 times higher than that of a commercial TiO 2 nanopowder. [ 17 ] The external morphologies can be controlled to some extent to optimize packing densities, e.g., by synthesizing nanopo-rous microspheres. Uniform micrometer-sized spheres (e.g., mesoporous SnO 2 microspheres) [ 126 ] are ideal for packing, and processes for forming electrodes from such particles are com-patible with existing fabrication methods. Spheres pack more
uniformly with fewer vacancies than irregularly shaped parti-cles and increase tap density. [ 155 ]
An additional advantage of nanoporous materials is that their skeletal walls are typically composed of well-interconnected crys-tallites or continuous amorphous phases. As such, interparticle contact within the porous solid is much better than in packed nanoparticles, resulting in better charge transport and also cir-cumventing potential health hazards associated with discrete nanoparticles. And while nanoporous solids contain internal features that bring with them the advantages of nanoparticles (short diffusion paths, high surface areas), the nanosized fea-tures are packaged into a material that has greater macroscopic overall dimensions and is therefore easier to handle. Certain electrode materials can be synthesized as monolithic nanopo-rous structures with dimensions of millimeters or centimeters, eliminating the need for binders or conductive additives. [ 157 ]
3.7. Void Spaces Constrain Growth of Active Material, Suppress Irreversible Phase Transformations, and Accommodate Volume Changes During Cycling
Several alloy phases of lithium have attracted great interest as anode materials for LIBs because of their high theoretical specifi c capacities. These include Li-Si, Li-Al, Li-Sn, Li-Pb, Li-As and other alloys. [ 158 ] Unfortunately, most of these alloys undergo large volume changes (as large as several hundred per-cent) during alloying and dealloying with lithium. The volume changes place large stresses on the material and eventually result in pulverization of particles, loss of interparticle contact, and loss of the mechanical integrity of an electrode. This leads to rapid capacity fade during cycling. Introducing pores into an alloy electrode can alleviate some of these stresses by pro-viding some room for expansion. In the case of 1.2 μ m-thick mesoporous Sn fi lms that were electrodeposited on a copper foil substrate in the presence of nonionic surfactant-based lyotropic liquid crystals, [ 27 , 28 ] mesopores helped to accommo-date volume changes so that reasonably high capacities were maintained after 100 cycles (425 mAh/g at 1C, 320 mAh/g at 5C). The pores were believed to diminish growth of Sn clus-ters. Nonporous control fi lms (1.0 μ m thick) lost most of their capacity within 20–40 cycles as particles became progressively more disconnected. In SnO 2 , the introduction of mesoporosity with a wide mesopore size distribution also led to improved cyclability. [ 24 ] Several examples of alloy anodes based on porous SnO 2 spheres and hollow Si nanospheres with high reversible capacities and good cycle-to-cycle capacity retention have been reported. [ 87 , 159 ]
Besides providing room for expansion during phase changes, nanoporous structures play another important role in facilitating the practical use of alloy electrodes and other electrode materials that swell and shrink during cycling. The continuous skeleton in mesoporous electrodes is less prone to disconnection during expansion/contraction processes on cycling than discrete nano-particles. [ 160 ] Mesoporous electrodes, similar to nanoparticles, can better accommodate strains from lithium insertion com-pared to micrometer-sized particles and thereby suppress irre-versible phase transformations that occur in microcrystalline materials. [ 15 ] For example, microcrystalline α -Fe 2 O 3 (hematite)
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Figure 12 . Volume changes in Li-alloy-based 3DOM anodes. (A) 3DOM SnO 2 as synthesized and (B) after charging with lithium. [ 83 ] (C) TEM image of 3DOM/m C/SnO 2 nanocomposite in which SnO 2 nanoparticles are encapsulated within mesoporous carbon walls. A few SnO 2 parti-cles are also found outside of mesopores on the macropore wall surfaces. [ 142 ] (D) SEM images of 3DOM Si at different degrees of lithiation. [ 164 ] Parts (A)–(B): Reproduced with permission. [ 83 ] Copyright 2004, Royal Society of Chemistry. Part (C): Reproduced with permission. [ 142 ] Copy-right 2008, Electrochemical Society. Part (D): Reproduced with permission. [ 164 ]
undergoes an irreversible phase transforma-tion to cubic Li x Fe 2 O 3 during Li-intercalation. This transformation is suppressed in the mes-oporous framework, which is able to accom-modate expansion during Li + intercalation. [ 15 ] Mesoporous β -MnO 2 (tetragonal (rutile) structure) undergoes a two-phase intercala-tion reaction during the fi rst discharge, fi rst forming Li x MnO 2 - β with an orthorhombic structure. [ 50 ] Subsequent cycles involve a single-phase intercalation process. In spite of signifi cant volume changes and large anisot-ropy during cycling, the mesopore structure is maintained during the cycling process. The thin walls of the mesoporous material can alleviate stresses that occur during the phase change. In a study of mesoporous Sn 2 P 2 O 7 , the stress-relieving mechanism was described as “breathing” of the overall mesoporous framework, accommodating volume changes as individual grains undergo alloying/dealloying processes. [ 25 ]
In some cases, composites containing an alloy and an inactive phase allow buffering of
the volume change and act as barriers against aggregation of active material so that smaller grain sizes are maintained. [ 161 ] In the above mesoporous Sn 2 P 2 O 7 system, the phosphate may have helped to prevent aggregation of Sn particles. [ 25 ] Similarly, in a 3DOM Sn-Co alloy fi lm electrode prepared by colloidal crystal templating [ 72 ] , the formation of an alloy prevented tin aggregation, and the macroporous structure alleviated stresses during cycling. In such alloys, a larger fraction of inactive matrix material enhances cyclability but reduces the available specifi c charge. [ 161 ]
In general, the volume changes during lithiation can be buff-ered in both macroporous and mesoporous electrodes, allowing these materials to have a higher reversible capacity than their nonporous counterparts. The cycle life of mesoporous alloy-based anodes is typically better than that of macroporous ones. Because the walls in mesoporous anodes are usually thinner than those in macroporous anodes, the effect of pulverization is less severe. Amorphous alloy anodes tend to be more resistant against electrode pulverization, whereas crystalline Li-alloy anodes usually offer a higher initial capacity. In terms of volume buffering, true porous materials are more effective than textural porous materials, because in the former, pores are distributed more evenly throughout the active phase. For example, carbon-coated Si nanoparticles [ 162 ] were observed to form pulverized Si and carbon particles after a few cycles, whereas mesoporous Si@C templated from SBA-15 was stable up to 80 cycles. [ 48 ]
3.8. A Supporting Structure Stabilizes Active Components with Low Cycle Lives
If active materials that undergo large volume changes cannot be suffi ciently stabilized through nanostructuring effects, it is possible to incorporate such materials in the walls, [ 163 ] in mesopores, [ 142 ] or on the surface of a supporting structure
that provides additional stabilization. [ 157 , 164 ] 3DOM SnO 2 , for instance, lost much of its capacity in relatively few cycles of expansion and contraction ( Figure 12 A, B). [ 83 ] By attaching SnO 2 nanoparticles to the surface of a 3DOM carbon support, the specifi c capacity was increased, but only for ca. 20 cycles. [ 157 ] At that point, many nanoparticles had likely separated from the carbon walls, losing contact with each other and with the conducting support. However, by trapping SnO 2 nanoparticles within mesopores of a hierarchically structured, three-dimen-sionally ordered macroporous/mesoporous (3DOM/m) carbon, the nanoparticles remained active over more than 100 cycles, and the observed capacity losses were mainly caused by the for-mation of an SEI layer (Figure 12 C). [ 142 ] Similarly, good cycling effi ciency over 20 cycles was observed for a mesoporous carbon/tin nanocomposite in which tin nanoparticles were introduced into the mesoporous carbon host by sonochemical insertion. [ 165 ] By this method, a 33 wt% tin loading was achieved with metallic tin clusters (6.4 nm) throughout the mesopores.
Various ordered mesoporous carbons have been loaded with Sn or SnO 2 and used as electrode materials for lithium ion bat-teries. [ 165–167 ] The mesoporous carbon matrix is most benefi cial if Sn or SnO 2 is encapsulated within the mesopores. External particles are subject to pulverization during cycling, limiting the reversible capacity of the composite electrode. [ 165 ] No SnO 2 clusters outside of mesopores were observed for CMK-3 loaded with SnO 2 and prepared by sonochemical synthesis. [ 166 ] Alloy anode materials have also been encapsulated in hollow carbon tubes [ 168 , 169 ] and mesoporous carbon nanowires. [ 170 ] Tin encapsulated in hollow carbon fi bers (200 nm diameter, tens of micrometers long) were prepared via electrospining of a poly(methylmethacrylate)-polyacrylonitrile-tin octoate solu-tion. [ 168 ] A fi ber composite containing 66 wt% of Sn showed a high reversible capacity, which was attributed to the buffering effect of the hollow carbon wires and the high specifi c surface area of this structure.
A carbon coating to stabilize an alloy electrode like SnO 2 can be derived from the template, for example, when carbonaceous spheres are used as the template for macroporous SnO 2 [ 171 ] or when a carbon layer remains after pyrolysis of a polymer-sphere template under an inert atmosphere. [ 172 ] A carbon coating can also be added after the synthesis of a porous electrode. A carbon coating on macroporous SnP 2 O 7 with irregular macropores sta-bilized the material so that its discharge capacity after 20 cycles was three times higher (392 mAh/g) than for the uncoated material (123 mAh/g). [ 129 ]
The challenges associated with Li-alloy phases that undergo large volume changes during lithiation and delithiation have also been demonstrated with 3DOM Si systems. [ 164 ] Silicon has a theoretical capacity of 4200 mAh/g forming Li 4.4 Si, but its volume changes by more than 300% during cycling (Figure 12 D). Amorphous 3DOM Si prepared by CVD in a colloidal crystal template had a low conductivity ( ∼ 10 − 6 S/cm) and exhib-ited high capacity only at low rates. The conductivity could be increased ( ∼ 10 − 3 S/cm) by converting the amorphous Si in the skeletal walls to nanocrystalline Si by thermal annealing in an inert atmosphere, but this material lost capacity very quickly as the periodic macropores structure was destroyed during cycling. A thin carbon coating raised the conductivity of the material further but was insuffi cient to stabilize it against degradation. Improved properties were found by employing 3DOM C/Si composites, in which amorphous Si was coated onto a 3DOM C skeleton derived from a sucrose precursor. This skeleton pro-vided better conductivity ( ∼ 10 − 1 S/cm), greater mechanical sta-bility for the composite, and improved capacity retention.
Carbon supports or coatings help to maintain and even improve the conductivity of these active materials. Noncon-ductive phases may also be helpful in stabilizing active com-ponents with low cycle lives. In mesoporous SnO 2 templated from SBA-15 silica by melt-infi ltration of SnCl 2 ·2H 2 O, some residual SiO 2 left from incomplete etching helped to prop up the SnO 2 phase. [ 61 ] At 6-wt% residual SiO 2 , the mesoporous SnO 2 was structurally stable up to 700 ° C. Incorporation of 3.9–6.0 wt% SiO 2 improved cyclability, so that after 30 cycles a specifi c capacity of 600 mAh/g was still maintained at a rate of 50 mA/g. It was proposed that the residual silica acted as a bar-rier to mitigate aggregation of Sn clusters during cycling.
3.9. A Secondary Conductive Phase Improves Conductivity and High Rate Capacities of Active Phases
Several active electrode materials with high specifi c capacities or other useful properties for LIBs have low intrinsic electronic conductivities. As bulk particles, such materials cannot be used in practical LIBs, except possibly at very low current densities. To improve the electronic conductivity of an electrode, a con-ductive agent such as carbon black is normally used, which may occupy 10–20 wt% of the electrode without contributing to its capacity. For micrometer-sized particles of active mate-rial, carbon is mixed with the active material only at the mac-roscopic level. Whereas the carbon improves the conductivity of the electrode by forming a percolating network that connects the particles of active material, the carbon has no effect on the conductivity of each particle itself. In the case of nanoporous
electrodes, however, it becomes possible to distribute a con-ductive phase throughout a particle of active material. In some cases, the conductive agent may also help to maintain the mesostructure of the active materials. Many examples of such nanocomposites have been described in the recent literature. Several different allotropes of carbon are used as a conductive phase, including carbon nanotubes, graphene sheets, amor-phous carbon layers, or graphitic carbon. In addition, metals or conductive metal oxides have been incorporated in porous nanocomposites with active material to improve the overall con-ductivity and rate capabilities of the electrode.
Various approaches to incorporate metallic conductive phases in porous electrodes have been explored. One approach involved a vacuum deposition of a 5 nm-thick Cu or Sn fi lm on the sur-face of mesoporous anatase TiO 2 electrodes. [ 19 , 173 ] Less than 1 wt% metal was added, and additional conductive carbon and binder were used in the preparation of the fi nal electrode. The vacuum deposition process did not change the anatase struc-ture. Both Cu and Sn metal coatings mitigated polarization and improved the electrode capacity compared to uncoated TiO 2 electrodes. The extent of improvement was similar over all rates. The degree of metal penetration into material and the distribu-tion of the metal throughout the electrode were not discussed, but it is likely that most of the metal was deposited only on the top surface. Partial incorporation of Sn within a mesoporous host is facilitated by imbibition from a solution precursor, as was demonstrated for a mesoporous TiO 2 /Sn nanocomposite. [ 174 ]
A more even distribution of a conductive phase throughout a mesoporous electrode is possible with RuO 2 , which forms a self-connecting, percolating, and highly conductive network throughout the pores. [ 175 ] Cryogenic decomposition of RuO 4 produces a web of interconnected, ca. 4-nm-diameter RuO 2 crystallites in the void space, leaving suffi cient room for pen-etration of electrolyte. Alternatively, the conductive network can be formed by thermal decomposition of a RuCl 3 solution that fi lls a mesoporous host. The latter method was applied to prepare mesoporous anatase TiO 2 /RuO 2 composite anodes. [ 151 ] Without the RuO 2 phase, at low current rates, mesoporous titania has a similar specifi c capacity as titania nanoparticles (5 nm in size). However, at higher rates the mesoporous titania loses capacity more quickly because of its relatively low con-ductivity. This situation was improved by incorporating 5 wt% RuO 2 in mesoporous anatase spheres. The resulting composite exhibited excellent capacity retention even at rates as high as 30C (91 mAh/g). Unfortunately, because of its high cost, RuO 2 is most suited for fundamental studies.
Carbon provides a less expensive option for the formation of a composite containing a conductive phase. Among carbon additives, nanotubes and graphene exhibit the highest elec-tronic conductivities. Graphite and glassy-carbon-based conduc-tive carbons are characterized by their graphitic domains, which contain mostly sp 2 carbon atoms, and disordered domains, which contain sp 3 carbon atoms. In general, a carbon sample with higher graphitic content has a higher conductivity. The conductivity of carbon is strongly dependent on the precursor and the temperature at which it is fabricated. Soft carbons have higher conductivities than hard carbons. Soft carbons are graphitizable carbons derived from coal tar, petroleum pitch, and organic polymers. These precursors already contain large
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fractions of sp 2 carbon, which can be converted to graphite at temperatures below 2000 ° C. Hard carbons are non-graphitiz-able carbons derived from phenolic resins, carbohydrates, and biomass. These precursors usually have high contents of sp 3 carbon and oxygen and cannot be fully converted to graphite even at temperatures higher than 3000 ° C.
Carbon phases can be included directly in the synthesis of the active material. For example, single walled carbon nano-tubes (SWNTs) were added to the synthesis mixture of a vanadia aerogel to enhance its conductivity. [ 123 ] The SWNT formed ca. 10 nm diameter bundles intertwined between 20–30 nm vanadia ribbons. The two phases formed good electrical contact while keeping suffi cient room for electrolyte access. This confi gura-tion maintained high lithium capacity at high discharge rates due to the excellent contact between the active and the conduc-tive phase. Similarly, the inclusion of cut SWNTs in the reac-tion mixture for microemulsion-templated mesoporous TiO 2 , produced a composite in which the SWNTs helped to suppress polarization during high rate cycling. [ 176 ] As a consequence, a capacity of ∼ 70 mAh/g was achieved at the high rate of 40C.
Alternatively, carbon coatings can be formed on the surface of a porous electrode material when a polymeric template decom-poses during pyrolysis in an inert atmosphere. By this approach mesostructured Li 4 Ti 5 O 12 /carbon [ 177 ] and mesoporous TiO 2 /carbon nanocomposites [ 178 ] were obtained via carbonization of a block-copolymer template. In the Li 4 Ti 5 O 12 /carbon nanocom-posite, the conductive carbon phase coated the mesopore walls and provided the material with a capacity of 115 mAh/g at 10C rate and high capacity retention. [ 177 ] For the mesoporous TiO 2 /carbon nanocomposite, the built-in carbon phase helped to maintain electrical contact during the delithiation steps which were ineffi cient without the carbon phase. [ 178 ]
A carbon coating can also be introduced after the formation of a mesoporous electrode. In the case of fl ower-like β -In 2 S 3 microspheres with a defect spinel structure, the addition of a thin amorphous carbon layer by hydrothermal treatment in a glucose solution improved the cyclability of the electrode sig-nifi cantly. [ 179 ] Without the carbon coating, the specifi c capacity decreased from ∼ 1100 mAh/g to 65 mAh/g within only ten cycles. [ 180 ] With the coating, a specifi c capacity of 400 mAh/g could be maintained over 50 cycles at a rate of 100 mA/g. Comparable observations were made for mesoporous silicon prepared by a magnesiothermic reduction of SBA-15 mesopo-rous silica. [ 115 ] To improve the conductivity of the particles, they were coated with a thin carbon layer by CVD, using acetylene as the carbon source. Carbon was deposited on both external and internal surfaces but left suffi cient mesopore space to accom-modate electrolyte and provide room for particle expansion during alloying. The composite material maintained a specifi c capacity of 1500 mAh/g over 100 cycles at 0.05C. At the same slow rate, the specifi c capacity of mesoporous Si without the carbon dropped from ca. 2700 mAh/g in the fi rst cycle to less than 1200 mAh/g, whereas the specifi c capacity of non-porous Si nanoparticles dropped to less than 300 mAh/g after only 20 cycles. Notably, the carbon-coated mesoporous Si achieved an impressive capacity above 500 mAh/g at the fast rate of 15C.
The arrangement of an active material/carbon composite may be reversed by starting with a porous carbon host and infi l-trating the host with active material. In such a system, the ratio
of active material to carbon is typically lower than when carbon is added to the active component. For example, CMK-3 carbon was impregnated with ∼ 37 wt% NiO nanoparticles. [ 44 ] Com-pared to the pure CMK-3 host, the nanocomposite had a slightly lower charge capacity in the fi rst cycle. On the other hand, the irreversible capacity loss in the fi rst cycle was much lower for the composite electrode, possibly due to NiO taking up the mesopore space that would otherwise be occupied by polymeric electrolyte decomposition products. At high rates (3200 mA/g) nanocomposite showed a stable charge capacity of 413 mAh/g, more than twice as much as that of CMK-3. After compensa-tion for the carbon content, an estimated charge capacity of 808 mAh/g was attributed to the occluded NiO nanoparticles, i.e., nearly complete utilization of NiO even at this high charge rate. A similar approach was used for monolithic carbon mono-liths, whose coral-like macropores and disordered mesopores were infi ltrated with LiFePO 4 precursor solution. [ 181 ] After pyrolysis in a reducing atmosphere to form carbon/LiFePO 4 composites, suffi cient porosity was maintained in the com-posite to permit infi ltration by an electrolyte. However, polari-zation of the electrode due to diffusion resistance was apparent even at a slow discharge rate of 0.1C, perhaps because little conductive carbon was present between the LiFePO 4 and the electrolyte interface.
Much less of a polarization effect was observed when the carbon phase was intimately mixed with LiFePO 4 by a more direct synthesis of a 3DOM LiFePO 4 /C composite using poly(phenol formaldehyde) (PF) as the carbon precursor. [ 182 ] Even though a phenolic resin is generally considered a hard carbon precursor, the graphitic domains in the composite were quite large ( ∼ 5 nm), and the sample was able to support current densities as high as 2720 mA g − 1 . The large graphitic domains were attributed to the presence of Fe 2 + in the precursor mix-ture, which can act as a catalyst for the formation of graphitic carbon. Because the carbon was formed in-situ under mild con-ditions (600 ° C under an N 2 atmosphere), no impurity phases such as Fe 3 P, Fe 4 C 3 were detected in the composite. Carbon also helped to maintain the mesostructure of the composite, which had collapsed in a control sample that was prepared without the PF sol.
3.9.1. Composites of Nanoporous Carbon with Sulfur
Sulfur has several attractive properties as a potential electrode material for rechargeable lithium batteries, including a very high theoretical capacity (1675 mAh/g) and a low price. A cell with a sulfur cathode and a lithium metal anode can provide a theoretical specifi c energy density as high as 2600 mWh/g at full discharge with an average voltage of 2.2 V. However, its bulk electrical conductivity is extremely low ( ∼ 10 − 30 S/cm at room temperature) and its ionic conductivity is also low. [ 183 ] Another shortcoming of sulfur-based rechargeable lithium batteries is their fast capacity fading upon cycling due to dissolution and shuttling of polysulfi de intermediates (Li 2 S n 2 − , n > 4) between the cathode and the anode. The last process also leads to the corrosion of lithium anodes. Recently the via-bility of sulfur-based electrodes for lithium batteries has been improved by forming composites of sulfur with nanoporous carbon.
It is diffi cult to produce nanoscale carbon/sulfur compos-ites by physical mixing. Instead, to achieve a homogeneous distribution of sulfur in a carbon matrix, a melt-diffusion method has been widely employed, in which the carbon/sulfur mixture is treated above the melting point of sulfur, usu-ally above 150 ° C, to allow diffusion of molten sulfur, or even vaporization of sulfur, into nanopores by capillary forces. [ 6 , 184 ] Nanoporous carbons with high surface areas help to ensure intimate contact between carbon and sulfur or other insoluble products (Li 2 S 2 , Li 2 S), and hinder the diffusion of polysulfi des away from the cathode. [ 6 , 185 , 186 ] Highly porous carbon was pro-duced by the pyrolysis of polyacrylonitrile, and its composite with 57 wt% sulfur showed an initial discharge capacity of 1155 mAh/g and a stable capacity of 745 mAh/g after 84 cycles at the current density of 40 mA/g. [ 187 ] In contrast, a composite with 75 wt% sulfur could not be charged due to the formation of poorly conductive and insoluble Li 2 S. A composite cathode based on disordered mesoporous carbon showed an initial capacity of 1300 mAh/g and 600 mAh/g after 40 cycles in an ionic liquid electrolyte. [ 185 ]
Better engineering of the mesoporous carbon structure can lead to improved performance. Mesoporous carbon CMK-3 has cylindrical pores in a hexagonal array [ 188 ] and can have a theo-retical sulfur loading up to 79 wt%. A CMK-3/S composite with 70 wt% sulfur was prepared by treating mesoporous carbon and sulfur at 155 ° C, intentionally leaving some space for the volume expansion from S 8 to Li 2 S. [ 6 ] The entrapment effect of the cylindrical channels helped to prevent loss of sulfur during electrochemical cycling. The composite delivered a discharge capacity of 1005 mAh/g at the fi rst cycle and around 800 mAh after 20 cycles. A polyethylene glycol coating on the outside of the CMK-3/S particles further increased the initial capacity to 1320 mAh/g with a stable capacity of 1100 mAh/g.
A scheme for the electrochemical reactions in the pores has been proposed ( Figure 13 ). Polysulfi des are formed from the
Figure 13 . Proposed scheme for the electrochemical reaction processes occurring in a cathode composed of sulfur in nanopores of a highly porous carbon (HPC). Reproduced with permission. [ 187 ] Copyright 2009, American Chemical Society.
Adv. Energy Mater. 2012,DOI: 10.1002/aenm.201200320
sulfur embedded in the pores and remain there during cycling, so that mesopores act as reservoirs for polysulfi des. High capacity retention of C/S composite cathodes at high sulfur load-ings within mesoporous carbons demonstrates the advantage of a mesoporous structure to mitigate the polysulfi de shuttle. To optimize the structure of mesoporous carbon for C/S composite cathodes, Li et al. synthesized mesoporous carbons with varying pore sizes and pore volumes and found that all the compos-ites exhibited similar performance at complete fi lling of pores and improved performance at partial fi lling, concluding that carbon hosts with higher pore volumes should be considered in order to achieve a practical amount of sulfur loading. [ 189 ] If a pore is fully occupied by sulfur, the lithium insertion capacity is reduced correspondingly due to the structural confi nement. A C/S composite with 50 wt% sulfur prepared from mesopo-rous carbon and coated with a conductive polymer showed a discharge capacity of ∼ 1400 mAh/g in the fi rst cycle and 750 mAh/g after 150 cycles. [ 189 ] Mesoporous carbon spheres were also studied as the conductive host, and a reversible capacity of about 800 mAh/g after 25 cycles was achieved with a gel elec-trolyte. [ 190 ] The addition of mesoporous silica into mesoporous C/S composites can stabilize the cycling performance because extra mesopores act as a reservoir for polysulfi des and prevent loss of active species during cycling. [ 191 ] The observed high cou-lombic effi ciency suggested that the polysulfi de shuttle effect was almost completely eliminated.
The distribution of sulfur in hierarchically porous carbon can be revealed by N 2 sorption analysis. Liang et al. synthesized hierarchically structured C/S composites using bimodal porous carbon with a uniform distribution of surfactant-templated mesopores and micropores that were created by KOH activa-tion. [ 186 ] Sulfur was dissolved in CS 2 and then loaded into the carbon host after solvent evaporation. A pore size distribution analysis of the composite with different sulfur loadings showed preferential fi lling of sulfur into the micropores fi rst, accompa-nied by a decrease in the average mesopore diameter. The com-posite exhibited an initial discharge capacity of 1585 mAh/g at 11.7 wt% sulfur loading which decreased to 818 mAh/g at 51.5 wt% sulfur loading. A similar trend of pore fi lling was observed in C/S composites based on bimodal mesoporous carbon with hexagonal mesopore channels (4–8 nm) and small intra-wall mesopores (1–3 nm) for C/S composites. [ 192 , 193 ] At 60 wt% sulfur loading, these composites delivered a reversible discharge capacity of about 800 mAh/g at 0.1 C [ 192 ] and about 500 mAh/g at 1C. [ 193 ]
3.9.2. Porous Current Collectors
In the above examples, the composite electrodes were typically supported on an external current collector for use in a cell and electrochemical testing. It is also feasible to build an electrode around a porous current collector to improve the electrical contact between the current collector and the active material. In one case, a 50- μ m-thick macroporous copper backbone was fabricated by tape casting of a slurry that contained CuO par-ticles ( < 5 μ m), a pore former, and a binder. The material was then sintered and fi nally heated in a reducing gas. The structure contained ca. 60% porosity with micron-sized macropores. [ 194 ] This current collector was then used as a support for a silicon
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layer [ 194 ] or a SnO 2 layer. [ 195 ] In both cases, it improved the adhesion of the electrode material compared to typically used copper fi lms and required no conductive additive nor binder. It also decreased the electrical resistance and improved capacity retention during cycling. [ 194 ]
Another system for which excellent rate capabilities were demonstrated was built around a 3DOM Ni current collector. [ 73 ] This was prepared by electrodeposition of Ni within a colloidal crystal of micrometer-sized polymer spheres. This support was then coated with various active cathode materials, including manganese oxide for LIBs and nickel oxyhydroxide for nickel-metal hydride batteries. The excellent contact between the cur-rent collector, the active phase, and the electrolyte that fi lled the remaining pore space in this tricontinuous structure resulted in excellent charge and discharge rates, allowing the LIB to be charged to 90% capacity in 2 minutes. The cycle life perform-ance of this system also remained very good.
Small portable device applications rely on battery systems that provide high capacities per unit of volume. In autonomous electronic devices with very small footprints that require inte-grated batteries, another parameter becomes important: the energy density or capacity per unit of footprint area. To be practical in a small sensor based on microelectromechanical systems (MEMS), for example, a battery should provide ∼ 1 J of energy per mm 2 or have areal capacities of ∼ 0.05–0.10 mAh/mm 2 . [ 196 , 197 ] The energy per unit area of existing two-dimen-sional thin fi lm lithium batteries has been reported to be in the range from ∼ 0.003–0.02 J/mm 2 . [ 196 , 198 ] These values are still too low to supply suffi cient energy to a device with a mm × mm footprint for a reasonable period of time. To increase areal capacities, three-dimensionally (3D) nanostructured or micro-structured, interpenetrating battery architectures are now being considered. [ 151 , 196 , 197 , 199–208 ] Such architectures can signifi cantly enhance the interfacial areas between electrodes and the sepa-rator and extend the battery volume into the third dimension, while maintaining short diffusion distances for charge carriers between opposite electrodes. The major approaches towards 3D nanostructured batteries are based on lithographically pat-terned microchannel plate electrodes or arrays of cylindrical electrodes, mesoporous sol-gel electrodes, and macroporous inverse opal electrodes. In each case, the patterned or porous electrode is coated with an electrolyte/separator layer and fi lled in with material for the opposite electrode. Alternatively, col-loidal self-assembly approaches towards rechargeable lihium-ion batteries are being investigated. [ 209 ]
Simulations predict that interpenetrating electrode structures outperform traditional planar designs at intermediate and high discharge rates. [ 207 ] The extension of a nanostructured battery into the third dimension effectively increases the geometric energy density. [ 196 ] A lithographically patterned, corrugated 3D microbattery with 1–2 μ m feature sizes, for example, has been predicted to provide a geometric energy density of 0.2 J/mm 2 . [ 204 ] Another 3D microbattery was built around silicon microchannel-plates containing hole arrays of 500- μ m-long and
< 50- μ m-diameter channels. [ 203 ] It provided an energy density of 0.38 J/mm 2 and a geometric capacity of 0.02 mAh/mm 2 , a value that is 20–30 times higher than that of an equivalent 2D thin-fi lm cell. [ 202 ] A thin-fi lm Ni current collector was fi rst laid down by electroless deposition, followed by electrodeposition of a 1- μ m thick MoS 2 cathode, addition of a solid polymer electro-lyte or hybrid polymer electrolyte layer, and lithiated graphite particles as the anode and the anode current collector. Another approach toward 3D batteries at this length scale used a con-formal MnO 2 fi lm on carbon foam with pore dimensions of several hundred micrometers. [ 208 ]
Instead of building a 3D battery around an electrode, as in the above examples, one might consider a solid electrolyte structure as a starting point. Following this approach, an all-solid-state battery was fabricated, which was composed of a composite cathode of LiMn 2 O 4 deposited on a 3DOM Li 1.5 Al 0.5 Ti 1.5 (PO 4 ) 3 (LATP) solid electrolyte, a similar composite anode of Li 4 Mn 5 O 12 on 3DOM LATP and a solid polymer electrolyte as a separator between the two composite electrodes. [ 84 ] However, in this con-fi guration the battery is not a truly interpenetrating network of both electrodes and the electrolyte. Thus, an alternative con-fi guration employed a Li 0.35 La 0.55 TiO 3 ceramic with honeycomb structure as a solid electrolyte with a LiCoO 2 cathode on one side and a Li 4 Mn 5 O 12 anode on the other side to form an all-solid-state 3D battery. [ 210 ] The membrane had arrays of 180 μ m holes on both sides that formed an interdigitated pattern. These holes were fi lled with particles of active material by vacuum infi ltration of a suspension, followed by calcination (800 ° C for LiCoO 2 and 700 ° C for Li 4 Mn 5 O 12 ). A thin gold fi lm was sputtered onto the electrode particles to reduce the electrical resistance between the particles and the current collector. The complete battery was operated at 1.1 V, but had a low discharge capacity of only 0.073 μ Ah/mm 2 at 30 ° C.
In principle, higher geometric energy densities should be possible by using electrodes with smaller pore sizes. The syn-thetic challenges involved in assembling an interpenetrating battery structure built around a mesoporous electrode are sig-nifi cant but not insurmountable. An interpenetrating structure may be based on sol-gel derived aerogels or ambigels, including carbon aerogels, porous V 2 O 5 , and porous MnO 2 . [ 206 ] This approach was conceptually illustrated for a 3D interpenetrating structure formed around a mesoporous MnO 2 ambigel whose interior walls were coated with an ultrathin poly(phenylene oxide) (PPO) separating layer. The pores were then infi ltrated with 1–2 nm RuO 2 nanoparticles. In this approach it is impor-tant that the electropolymerization of PPO is a self-limiting process. It proceeds only on a conducting substrate. As soon as the coating covers the surface of the porous electrode surface, the electrooxidation reaction stops. As a result, the coatings are ultrathin ( < 15–20 nm), conformal to the electrode surface, and effectively pinhole free. [ 199 , 200 , 211 ] The ruthenia component was also strategically chosen, as it could be incorporated in the PPO-coated MnO 2 electrode under mild conditions that did not destroy the PPO layer. The Li-ion uptake of ruthenia depends on its degree of crystallinity and is highest ( ∼ 260 mAh/g) for amorphous RuO 2 that contains a small fraction of more con-ductive, crystalline RuO 2 in the rutile phase. [ 212 ]
The slightly larger macropore size regime with pores in the 200–2000 nm diameter range can simplify the infi ltration of
multiple components compared to mesopores, while still prom-ising larger geometric capacities than micromachined features. For the macropore size regime, a proof of concept has been shown for 3DOM electrodes prepared by colloidal crystal tem-plating, where critical feature sizes are ca. 0.2–1 μ m for pores, 20–50 nm for pore walls, and interfacial areas can be much greater than in lithographically patterned arrays. [ 199 , 200 ] The nanoscale interpenetrating battery was built around a mono-lithic 3DOM carbon anode on a conductive substrate that acted as a current collector. [ 213 ] The voids in 3DOM carbon are large enough to permit functionalization of the large internal carbon surface with multiple components, including electrolyte layers and cathode material. The entire surface of the macroporous anode (except for the point of contact with the current collector) was coated with a thin, conformal layer of polymer (PPO) that acted as the separator and electrolyte membrane. The remaining volume was infi ltrated with a vanadia ambigel as a cathode material in a procedure that was mild enough to pre-serve the polymeric separator ( Figure 14 ). A current collector was attached to the cathode surface, which extended slightly beyond the porous area. After lithiation, an open circuit voltage of 2.76 V was measured, and the cell was cycled multiple times between charged and discharged states.
Although this approach shows promise and simplifi es the incorporation of multiple components in macropores compared to mesopores, several challenges remain before making such interpenetrating cells more practical. The 3DOM cell behaved more like a pseudocapacitor than a true battery. While this behavior confi rmed that no hard shorts were present between opposing electrodes, charge and discharge rates were very slow, especially in early prototypes. [ 199 ] In subsequent work, the cycling current of the cell was raised from 1 μ A to 20 μ A, the capacity on the fi rst discharge increased from 70 μ Ah/g to 16 mAh/g, and the reversible discharge capacity from 0.7 μ Ah/g to 0.35 mAh/g. [ 200 ] Although these specifi c capacities are still small, the geometric capacities and energy densities are impres-sive during early cycles: 0.1 mAh/mm 2 (0.1 J/mm 2 ) during the fi rst cycle and > 0.004 mAh/mm 2 ( > 0.04 J/mm 2 ) up to the tenth cycle. It was possible to briefl y drive an LCD display from the 3DOM cell.
Several reasons for the limited performance of these inter-penetrating electrochemical cells could be identifi ed, including
Figure 14 . Schematic diagrams (top) and SEM images (bottom) of the 3D iC (dark grey). 3DOM C coated with PPO (white). Coated monolith infi ltrateafter 100 cycles. Assembled cell with current collectors.
Adv. Energy Mater. 2012,DOI: 10.1002/aenm.201200320
incomplete penetration of cathode material into the macro-scopic 3DOM anode monoliths, limited connectivity through ca. 100 nm-windows between adjacent macropores, and a relatively high resistance of the vanadia cathode material, which demanded very low lithiation and charging rates. [ 199 , 200 ] In addition, it would be desirable to improve the mobility of Li-ions through the separator membrane and enhance the con-ductivity of the carbon-based anode. This last requirement can be achieved by building the interpenetrating battery around a macroporous carbon foam that is embedded within a highly conductive carbon-fi ber paper. The advantage of such a com-posite electrode was recently demonstrated for a resorcinol-formaldehyde-derived carbon foam (50–300 nm interconnected macropores with ca. 20 nm thick walls) supported on carbon paper. [ 7 ] The carbon paper increased the overall electronic con-ductivity of the porous electrode by about two orders of magni-tude to values between 25–220 S/cm, depending on the type of carbon paper used.
These early demonstrations of 3D-interpenetrating batteries show that synthetic methods for fabricating such complex structures exist. However, the procedures are complex and not yet scalable. Further development is needed in this area. Ideally, one would like to employ a self-assembly approach. [ 209 ] First steps towards this concept were demonstrated with a prototype, self-organized electrochemical cell using LiCoO 2 and graphite colloids as the active electrode materials. [ 209 ] This requires suf-fi cient short-range repulsion between the opposite electrode components in an appropriate liquid medium. Given the rapid development of porous and nanostructured materials synthesis, more progress in this area is expected over the coming years.
5. Conclusions and Future Directions
The tremendous recent progress in synthesizing porous, nanos-tructured materials with a high degree of control over pore architecture has benefi ted numerous fi elds, including those related to electrical energy storage. Among electrical energy storage devices, LIBs continue to be the main power sources for portable devices, now expanding into transportation and sta-tionary energy storage sectors. As the studies highlighted in this review demonstrate, porosity and nanostructure can address
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nterpenetrating electrochemical cell components. From left to right: 3DOM d with vanadia (middle grey); the SEM image shows a cut through the cell
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some of the technical needs and provide signifi cant advantages for LIBs, if applied in the proper context. Depending on the application though, different technical requirements must be met. For transportation applications, rate capability is the more important factor, whereas energy density is of highest priority for stationary energy storage.
Are porous electrodes really better for LIBs? At slow rates of charge and discharge, porous electrode materials do not nec-essarily provide improved capacities over traditional materials with micrometer-sized particles. For those applications, the additional effort and cost associated with controlling the par-ticle architecture on a nanometer length scale may therefore not be justifi ed. However, high rate applications benefi t from nanoporous electrode materials that provide short diffusion paths for charge carriers in bicontinuous electrode and electro-lyte phases. A large interfacial area between these phases facili-tates charge transfer, potentially reducing internal resistances and further improving rate capabilities. The high-rate advan-tages may benefi t EV and HEV applications, as well as applica-tions requiring pulses of high power, such as those in certain communication devices.
Volumetric energy density is an important consideration in small portable devices. Non-utilized porosity increases the amount of space required for a given amount of stored energy. This limitation of porous electrodes in volumetric energy densi-ties is only overcome if the presence of pores enables the use of novel electrode materials with signifi cantly higher bulk energy densities than those currently employed.
For transportation applications, in which space for the power sources is not as limited as in portable devices, safe, inexpen-sive, and environmentally benign electrode materials with good rate capability are urgently needed. On the cathode side, spinel-type LiMn 2 O 4 and olivine-type LiFePO 4 satisfy the fi rst three criteria and have been actively investigated. [ 2 ] The rate capa-bility of these electrodes can be signifi cantly improved through nanostructuring and introducing nanopores into the materials. The challenges for synthesizing nanoporous structures of these materials are their fast crystallization rates at moderate tem-peratures, causing collapse of the porous structure upon heat treatment. The nanocasting methods discussed in this review and multi-step heat treatment processes have addressed these issues. However, further progress in materials synthesis is needed because preparations of these electrode materials with complex composition usually involve multiple precursors, requiring a wise selection of starting materials to form the target phase rather than mixtures of metal oxides or metal salts. On the anode side, carbon continues to be the most prominent elec-trode material due to its low cost and reasonably high capacities. Yet, the rate performance of nanostructured and nanoporous carbon materials is limited by the formation of an SEI layer on the large surface, slowing down the diffusion of lithium ions. Other anodes, such as Li 4 Ti 5 O 12 and TiO 2 with higher voltages that exceed the reduction potentials of most aprotic electrolytes, offer improved safety and longer cycle lives and may compete with graphite in the future, particularly when their lower capaci-ties are not really a huge disadvantage for transportation appli-cations. Some of the work on nanoporous LTO and TiO 2 anodes reviewed in this article showed very promising results with very good cycle lives and reasonable capacities.
In addition, LIB system designs with porous electrode mate-rials permit the consideration of novel electrode materials with desirable properties but which are unfeasible when used in bulk form, either because of large volume changes during elec-trochemical cycling or low intrinsic conductivities. Such mate-rials may particularly benefi t stationary storage applications that require high specifi c energy density cells. On the cathode side, research focuses on increasing the redox potential and the capacity of the electrode materials. One approach, doping with fl uorine, generally increases the voltage of an electrode material but it also increases the band gap, thereby lowering the elec-tronic conductivity. [ 214 ] Ionic conductivities of these materials are often also low. Nanoporous materials with their short ionic diffusion paths can address the issue of low ionic conductivity, while adding conductive agents in a nanoporous composite could overcome the limitations of low electronic conductivity. Improvements in the capacity of cathode materials have been achieved by mixing metal oxides of Ni, Co, and Mn. [ 2 ] Another way to improve the capacity is to use electrode materials that are able to insert more than one Li per 3d metal, such as Li 2 MnSiO 4 or some conversion cathodes (CoO, NiP 2 ). [ 215 ] However, the insertion/extraction of the second Li into LiMnSiO 4 is poorly reversible and could be improved by fabricating nanostructured porous electrodes. Nanocomposites with well-dispersed con-ductive phases facilitate charge transport through highly resis-tive phases. On the anode side, alloyed anode materials with high theoretical capacities are the most promising candidates if their volume expansion problems during cycling are properly addressed. Porous materials can effectively act as buffer spaces against volume changes, thereby minimizing cracking of the electrode. Properly constructed nanocomposites can also trap active phases that might otherwise dissolve in an electrolyte, e.g., in the case of sulfur-based electrodes. Advances with such materials promise to provide the extended capacities desired for small portable electronic devices. Finally, porous materials provide a platform for 3D-interpenetrating battery architectures needed for small devices with limited available “real estate” [ 205 ] for an integrated battery.
For the choice of pore sizes, it appears that materials with medium-sized to large mesopores provide the greatest advan-tages in terms of rate capabilities and effective use of available volume. The best packing densities are obtained when uni-formly-sized spherical particles with external dimensions of a few micrometers and internal mesopores are employed. These materials exceed the packing density of nanoparticles while pro-viding the same short diffusion paths. Macroporous solids and materials with hierarchical pore structure become most impor-tant if highly viscous electrolytes must be used or if the internal pore volume needs to be utilized to accommodate other solid phases, as in the case of 3D interpenetrating batteries. With regards to pore morphologies, bicontinuous structures are most desirable as they minimize dead ends and enable effi cient transport of charge carriers through both the pore space (Li ions, counterions in electrolyte) and the solid space (electrons and Li ions). However, more theoretical and systematic experi-mental studies are needed to identify specifi c pore geometries for optimal performance.
The biggest hurdles for porous electrodes to succeed in commercial LIB systems are the relatively complex synthesis
processes and the associated costs. However, with suffi cient incentives related to systems performance, such hurdles are not unsurmountable. Many of the synthetic methods used to pre-pare porous electrodes (hydrothermal and solvothermal reac-tions, templating, self-assembly) are already applied in com-mercial materials on large scales, zeolites and metal-organic frameworks (MOFs) being representative examples. With the prospect of limited availability of some existing LIB electrode materials in the future, alternate compositions —now feasible through nanostructure and nanocomposite formation—provide resources to secure the future of LIBs on a large scale. It should not be long before we see porous electrode materials incorpo-rated in commercial LIBs.
Acknowledgements For parts of the work highlighted in this review, we thank the following sources for funding: the National Science Foundation (DMR-0704312), the University of Minnesota Initiative for Renewable Energy and the Environment, the Offi ce of Naval Research (Grant N00014-07-1-0608), and the University of Minnesota Characterization Facility, which receives partial support from the NSF through the MRSEC, ERC, MRI, and NNIN programs.
Received: May 2, 2012 Revised: June 14, 2012
Published online:
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