Acs_ar_ar-2012-00176y 1.9Lithium Insertion in Nanostructured TiO2(B) ANTHONY G. DYLLA, GRAEME HENKELMAN, AND KEITH J. STEVENSON* Department of Chemistry & Biochemistry, The University of Texas at Austin, Austin, Texas 78712, United States RECEIVED ON JUNE 12, 2012 to become feasible alternatives to current technology, but only if scientists can develop energy storage materialsthat offer high capacity and high rate capabilities. Chemists havestudied anatase, rutile, brookite and TiO2(B) (bronze) in bothbulk and nanostructured forms as potential Li-ion batteryanodes. In most cases, the specific capacity and rate of lithiationand delithiation increases as the materials are nanostructured.
Scientists have explained these enhancements in terms of highersurface areas, shorter Liþ diffusion paths and different surfaceenergies for nanostructured materials allowing for more facilelithiation and delithiation. Of the most studied polymorphs,nanostructured TiO2(B) has the highest capacity with promising high rate capabilities. TiO2(B) is able to accommodate 1 Liþ per Ti,giving a capacity of 335 mAh/g for nanotubular and nanoparticulate TiO2(B). The TiO2(B) polymorph, discovered in 1980 by Marchand andco-workers, has been the focus of many recent studies regarding high power and high capacity anode materials with potential applicationsfor electric vehicles and grid storage. This is due to the material's stability over multiple cycles, safer lithiation potential relative to graphite,reasonable capacity, high rate capability, nontoxicity, and low cost (Bruce, P. G.; Scrosati, B.; Tarascon, J.-M. Nanomaterials for RechargeableLithium Batteries. Angew. Chem., Int. Ed. 2008, 47, 2930"2946). One of the most interesting properties of TiO2(B) is that both bulk andnanostructured forms lithiate and delithiate through a surface redox or pseudocapacitive charging mechanism, giving rise to stable high ratecharge/discharge capabilities in the case of nanostructured TiO2(B). When other polymorphs of TiO2 are nanostructured, they still mainlyintercalate lithium through a bulk diffusion-controlled mechanism. TiO2(B) has a unique open crystal structure and low energy Liþ pathwaysfrom surface to subsurface sites, which many chemists believe to contribute to the pseudocapacitive charging.
Several disadvantages exist as well. TiO2(B), and titania in general, suffers from poor electronic and ionic conductivity.
Nanostructured TiO2(B) also exhibits significant irreversible capacity loss (ICL) upon first discharge (lithiation). NanostructuringTiO2(B) can help alleviate problems with poor ionic conductivity by shortening lithium diffusion pathways. Unfortunately, this alsoincreases the likelihood of severe first discharge ICL due to reactive Ti"OH and Ti"O surface sites that can cause unwantedelectrolyte degradation and irreversible trapping of Liþ. Nanostructuring also results in lowered volumetric energy density, whichcould be a considerable problem for mobile applications. We will also discuss these problems and proposed solutions.
Scientists have synthesized TiO2(B) in a variety of nanostructures including nanowires, nanotubes, nanoparticles, mesoporous- ordered nanostructures, and nanosheets. Many of these structures exhibit enhanced Liþ diffusion kinetics and increased specificcapacities compared to bulk material, and thus warrant investigation on how nanostructuring influences lithiation behavior. ThisAccount will focus on these influences from both experimental and theoretical perspectives. We will discuss the surface chargingmechanism that gives rise to the increased lithiation and delithiation kinetics for TiO2(B), along with the influence of dimensionalconfinement of the nanoarchitectures, and how nanostructuring can change the lithiation mechanism considerably.
development of energy storage materials that offer high In order for electric vehicles and grid storage devices to capacity and high rate capabilities is needed.2,3 Pseudoca- become viable alternatives to current technology, further pacitors can be considered hybrids of traditional batteries 1104 ' ACCOUNTS OF CHEMICAL RESEARCH ' 1104–1112 ' 2013 ' Vol. 46, No. 5 Published on the Web 02/20/2013 www.pubs.acs.org/accounts & 2013 American Chemical Society Lithium Insertion in TiO2(B) Architectures Dylla et al.
(high energy density but typically poor power output) anddouble layer capacitors (high power output during shortbursts but low energy density). The pseudocapacitive energystorage mechanism is different from batteries in that surfaceredox properties dominate the charge-transfer processesrather than normal Faradaic diffusion-controlled insertionprocesses.4 Metal oxides such as RuO2 and IrO2 act as pseudo-capacitors offering exceptional power, fast charging, and long-term stability while also affording some of the advantages oftraditional secondary batteries such as reasonable storagecapacity.4,5 Pseudocapacitors with moderate energy densityat high charge rates over many cycles could find applications inhybrid-electric or electric vehicles. The pseudocapacitor could FIGURE 1. Scheme showing idealized voltage and differential capacity be employed when fast power delivery during acceleration is (dC/dV) profiles for three basic charge storage mechanisms along with acontinuum of their energy storage properties. Pseudocapacitors are required. Efficient use of renewable energies through load attractive because they take advantage of the attractive properties of leveling will also require the ability to store and deliver charge both batteries (charge storage) and capacitors (high rate delivery or rapidly.3 For these applications, RuO 2 and IrO2 would be cost prohibitive making, TiO2 a relatively inexpensive option. By Equation 2 describes current due to normal diffusion con- nanostructuring certain electroactive materials, this surface trolled Faradaic Liþ insertion processes where the current is charge-transfer process (pseudocapacitive effect) becomes linear with the square root of scan rate. In the case of the dominant storage mechanism and can offer 10"100 times pseudocapacitance, the extent of charge (Δq) is dependent the capacitance of a traditional carbon-based double layercapacitor.4 As mentioned in the Conspectus, TiO upon the change in voltage (ΔV), and thus, the total charge anode material in that both bulk and nanostructured forms passed, d(Δq)/d(ΔV), is the equivalent of capacitance and lithiate/delithiate through this pseudocapacitive mechanism, gives rise to the sloping galvanostatic (I"V) profiles often making it an attractive material for the applications noted exhibited by those materials.4 Figure 1 shows a general schematic of the three basic charge storage mechanisms in Electrochemical methods have been used to measure relation to galvanostatic and differential capacity (dC/dV) Liþ diffusion kinetics as well as the relative contributions of plots. Liþ insertion via pseudocapacitance is similar to a both capacitive (including double layer capacitance and battery in that a redox reaction is required for charge storage pseudocapacitance) and diffusion controlled insertion pro- but dissimilar in that the process is not diffusion limited.
cesses to the overal capacity of nanostructured anatase Wang and co-workers have systematically studied how materials.6,7 More specifical y, cyclic voltammetry (CV) can be nanoparticle size affects the relative contributions of pseu- used to determine the general relationship between current docapacitance and bulk diffusion controlled intercalation of and scan rate by eqs 1 and 2 below: Liþ into nanocrystalline anatase.9 By combining eqs 1 and 2 above and simplifying to eq 3 below: i(V) ¼ k1v þ k2v1=2 They were able to determine relative contributions of pseudocapacitance and diffusion controlled intercalation where CΦ is surface capacitance, ν is the scan rate, n is the at a given potential using scan rate studies. Figure 2 number of electrons, F is the Faraday constant, A is the shows that as the particle size decreased, the contribution electrode area, c is the concentration of Liþ, D is the diffusion of pseudocapacitance increased as well as slightly in- coefficient, R is the transfer coefficient, R is the gas constant, creasing the overall storage capacity of the material. The and T is the temperature. Equation 1 describes capacitive galvanostatic plot shown in Figure 2 also shows the more current due to surface confined redox processes (pseudo- capacitive-like sloping behavior indicative of this surface capacitance) where the current is linear with scan rate.
Vol. 46, No. 5 ' 2013 ' 1104–1112 ' ACCOUNTS OF CHEMICAL RESEARCH ' 1105 Lithium Insertion in TiO2(B) Architectures Dylla et al.
FIGURE 2. Contribution of diffusion controlled capacity and surface capacitance (pseudocapacitance) as a function of anatase nanoparticle size (left)and galvanostatic discharge curves for the three sizes studied (right). Reprinted with permission from ref 9. Copyright 2007 American ChemicalSociety.
2. The TiO2(B) Structure In subsequent cycles, LixTiO2(B) is reversibly lithiated/ The advantages of using TiO delithiated as shown in eq 5.
2 as an anode in rechargeable lithium ion batteries lie in its characteristic safety and stabi- LixTiO2(B)TLix " yTiO2(B) þ yLiþ þ ye " lity. Graphite is the most widely used anode material inrechargeable lithium ion batteries due in part to its low The unit cell contains 8 Ti sites and 10 Liþ sites, giving a lithiation potential (∼0.1 V vs Li/Liþ) that allows for a large theoretical capacity of 1.25 Liþ/Ti (∼420 mAh/g), though voltage difference between cathode an anode and reason- calculations suggest that because of Liþ"Liþ repulsions ably high capacity. The fact that graphite lithiates at a only 8 Liþ sites can be filled giving a capacity of 1.00 Liþ/ potential near that of the Li/Liþ couple poses a problem in Ti (335 mAh/g).11 Liþ can bind to three unique sites within that the lithium electroplating can cause short circuit and the crystal: four A1 and four A2 sites sit near equatorial thermal runaway conditions resulting in combustion of and axial oxygens in the titania octahera, respectively, organic electrolyte and catastrophic battery failure. Choos- and two C sites lie in the open channel along the b-axis.
ing an anode with a higher lithiation potential such as TiO2( Figure 3 shows the crystal structure along with labeled Liþ ∼1.6 V vs Li/Liþ) greatly reduces the chance of this type of battery failure. Among the common polymorphs, TiO site occupations.
has attracted recent attention due mainly to high energy Much of the interest in TiO2(B) as an anode material lies in density, but also because of the ability to nanostructure this its unique crystal structure. Compared to other titania poly- polymorph into several distinct architectures which provides morphs, TiO2(B) has the lowest density, and the perovskite- opportunities to systematically study the charge storage like layered structure along with the open channel parallel to the b-axis suggests that fast Liþ diffusion should be possible. While high rate capabilities have been observed 2(B) was first synthesized in 1980 by Marchand and co-workers from the layered titanate K for nanostructured TiO 2(B), the mechanism responsible is still debated.11 7Liþ spin-alignment echo nuclear magnetic 2Ti4O9 via acid washing and finally dehy- drated to the layered TiO resonance (SAE-NMR) correlation spectroscopy studies have 2(B) structure.10 TiO2(B) has a monoclinic C2/m structure comprised of edge- and corner- shown Liþ self-diffusivity to be quite slow even in nanowire 6 octahedra with an open channel parallel to 2(B) suggesting that another mechanism is at the b-axis that sits between axial oxygens. Initially TiO play besides facile Liþ diffusion through the b-axis channel in lithiates as shown in eq 4.
the bulk of the crystal.12,13 High surface areas and uniquesurface energetics have been largely invoked to describe the xLi þ TiO2(B) þ xe "TLixTiO2(B) high capacity and rates for nanostructured TiO2(B).1 1106 ' ACCOUNTS OF CHEMICAL RESEARCH ' 1104–1112 ' 2013 ' Vol. 46, No. 5 Lithium Insertion in TiO2(B) Architectures Dylla et al.
FIGURE 3. Unit cell of TiO2(B) with idealized Liþ insertion sites.
3. Lithiation of TiO2(B) Bulk TiO2(B). Bulk TiO2(B) lithiates to x = 0.7"0.8 (LixTiO2(B)) upon first discharge (lithiation) with stable slowrate cycling in the range of x = 0.3"0.6.14"17 The galvano-static charging/discharging profiles of bulk TiO2(B) show twodistinct plateaus between 1.4 and 1.6 V vs Li/Liþ uponreduction (intercalation) and similar peaks at slightly higher FIGURE 4. Typical cyclic voltammogram of Liþ insertion and potentials (ΔV ≈ 100 mV) upon oxidation (deintercalation).
deinsertion into bulk TiO2(B) along with an inset of the peak current Zukalova and co-workers were the first to determine the for both oxidation and reduction as a function of scan rate that shows pseudocapacitive Liþ insertion behavior of TiO a linear relationship associated with surface charging mechanism.
Reprinted with permission from ref 18. Copyright 2005 American rate dependent studies on bulk TiO2(B) showing a linear rela- Chemical Society.
tionship between scan rate and peak current (see Figure 4).18 Nanowires and Nanotubes. Much of the initial research bulk TiO2(B). The main difference between the 1-D nano- on the lithiation of nanostructured TiO2(B) comes from Bruce architectures is that the nanotubes have a more sloped and co-workers on TiO2(B) nanowires and nanotubes.16,19,20 profile and increased first cycle ICL relative to the nanowires.
Armstrong et al. used a simplified hydrothermal reaction Sloping profiles in galvanostatic charge/discharge cycles are involving anatase powders in KOH solutions to produce the often observed for nanostructured Liþ intercalation materi- potassium titanate precursor followed by ion exchange in als. This behavior was explained in terms of the curvature of acid to form the hydrogen titanate that was then dehydrated the nanotube walls leading to strain-induced surface free to form TiO2(B) nanowires and nanotubes.19,21 They were energy changes. As described above, the pseudocapacitive able to produce either nanotubes or nanowires by control- charging mechanism requires that the degree of lithiation be ling the time and temperature of the final dehydration step.
dependent on the potential. In this mechanism, the materi- Shorter times at lower temperatures produced nanotubes al's surface acts as a solid solution host to Liþ insertion and with an external diameter of 10"20 nm and an internal the free energy of the surface changes incrementally as a diameter of 5"8 nm with micrometer lengths, while longer function of Liþ concentration. This is in contrast to nanowires times at higher temperatures produced nanowires with 20" that have relatively large widths and less surface strain and act 40 nm diameters with lengths of several micrometers. Other more like bulk materials. The higher degree of ICL for TiO2(B) methods for creating nanowires and nanotubes have been nanotubes relative to nanowires was also explained in terms of reported including low temperature22 and microwave23 meth- higher reactivity toward electrolyte decomposition on the ods. In their study of low temperature synthesis methods for strained nanotube surface. The first discharge capacity of TiO2(B) producing TiO2(B) nanowires, Daoud and Pang found that the nanowires and nanotubes at a charge rate of 50 mA/g synthesis of nanotube, nanoribbon, or nanowire structures was (0.15 C) was 305 (0.91 Liþ/Ti) and 338 (1.01 Liþ/Ti) mAh/g, highly dependent upon the temperature of the reaction.22 This respectively. The capacity dropped to 200 and 230 mAh/g, suggests that the collapse and dehydration of the layered respectively, for the second cycle discharge.
hydrogen titanate intermediate plays a significant role in the In terms of specific site filling and lithiation mechanisms, resultant TiO2(B) nanostructure formed.
several experimental studies have suggested that facile The galvanostatic charge/discharge cycles of nanowires lithiation occurs radially into the walls of the nanotubes and nanotubes typically exhibit redox plateau positions, or and perpendicular to the flat surface of the nanowires rather redox peaks in differential capacity plots (dC/dV), similar to than down the long axis of the 1-D structures that is also Vol. 46, No. 5 ' 2013 ' 1104–1112 ' ACCOUNTS OF CHEMICAL RESEARCH ' 1107 Lithium Insertion in TiO2(B) Architectures Dylla et al.
parallel to the b-axis or C site tunnels.1,24 Brutti and co-workers used combined FTIR, XPS, and electrochemicalstudies to investigate the increased ICL observed for nano-tube geometries.25 Earlier work had suggested that eitherelectrolyte decomposition or high-energy surface sites act-ing as irreversible Liþ traps during the first lithiation cyclewere the cause of the ICL.24 The latter was ruled out as themain contributor because surface treatments of TiO2(B)nanotubes with Li(CH3CO2) dramatically mitigated the ICL.
The lithium acetate reacts with surface Ti"OH to form aTi"O"Li tube surface that is less reactive toward electrolytedecomposition.
Nanoparticles. The synthesis of TiO2(B) nanoparticles was first described by Kobayashi and co-workers.26 Theyfirst synthesized a Ti-glycolate complex that was then re-acted with H2SO4 to form the hydrogen titanate precursor.
Finally, the solution of hydrogen titanate was hydrother-mally reacted to form nanoparticles of TiO2(B) in the 3"6 nmsize range.
Galvanostatic charging curves of phase-pure TiO2(B) nano- particles have only been recently reported by Ren and FIGURE 5. Gravimetric capacity (a) and relative volumetric charge Bruce.27 The redox peak positions reported in the dC/dV retention (b) as a function of charge rate for various TiO2 nanostructures.
Reproduced from ref 27. Copyright 2012 John Wiley & Sons.
plots were similar to those of nanowire and nanotubegeometries, implying a similar Liþ insertion mechanism.
capability. Porosity and high surface area has been added The nanoparticle dC/dV plot does differ from both the to anatase through sol"gel routes involving structure direct- nanowire and bulk architectures in that a large capacitive- ing diblock copolymers,28 by templating of preformed ana- like region was observed in the 1.4"1.0 V range. This tase nanocrystals,28 and by reactive ballistic deposition.29 capacitive profile is similar to that observed for nanotubes This increased porosity can enhance both overall storage which may suggest that curvature and/or strain near the capacity and high rate capacity due to improved Li-ion surface of these nanomaterials plays an important role in coupled electron transfer kinetics, faster diffusion of Liþ the lithiation mechanism. The first discharge capacity for through the interconnected TiO2 network, and increased TiO2(B) nanoparticles was reported as 322 mAh/g (0.96 Liþ/Ti) electrode/electrolyte contact areas.
at 50 mA/g (0.15 C) charge rate. This discharge capacity is Interestingly, when using the sol"gel method for synthe- close to the theoretical value, but some ICL is observed upon sizing mesoporous TiO2, a second TiO2 phase is often first charge pointing to redox behavior associated with introduced as a minor component. The hydrolysis condi- electrolyte degradation on the first cycle. Upon second tions required for the formation of Ti"O"Ti oligomeric discharge at the same rate, the capacity falls to 251 mAh/g networks are difficult to control and can result in multiphase (0.75 Liþ/Ti) with good retention upon charge (96%). The materials. Kavan and co-workers noted the impurity during nanoparticulate TiO2(B) outperforms nanotubes, nanowires, electrochemical experiments showing unique Liþ insertion and bulk TiO2(B) in both fast and slow charge rate capacity as at 1.6 V.30 This phase was first believed to be an amorphous can be seen in Figure 5 where several nanoarchitectures of surface anatase species, but was later confirmed to be a TiO2(B) along with other titania polymorphs are compared.
TiO2(B) impurity.18 In the case of reactive ballistic deposition The authors attributed the enhanced capacity and kinetics of of titanium in an oxygen environment, Lin and co-workers the nanoparticles to increased surface structural distortions observed lithium insertion peaks in the cyclic voltammetry making Liþ transport more facile.
similar in position to those found in TiO2(B) and determined Mesoporous TiO2(B). As mentioned in the previous sec- that a mixture of surface pseudocapacitance and diffusion tions, high surface area, nanostructured TiO2(B) with strained controlled Liþ insertion was responsible for the charging surfaces appear to enhance both the capacity and rate behavior.29 These examples point to the complex conditions 1108 ' ACCOUNTS OF CHEMICAL RESEARCH ' 1104–1112 ' 2013 ' Vol. 46, No. 5 Lithium Insertion in TiO2(B) Architectures Dylla et al.
that are required to control titania polymorphism at the followed by dehydration at 350 !C.35 They argued that the increased capacity of the layer-structured derived TiO2(B) re- Proch!azka et al. reported on a dip-coating method to lative to conventional TiO2(B) was due to the increased surface produce mesoporous TiO2(B) on conductive substrates and area. Stable 253 mAh/g capacity was achieved at 100 mA/g showed fast lithiation kinetics of the thin films for smart (0.3 C). Liu and co-workers used a modified version of the Xiang electrochromic window applications.31 Some anatase con- synthesis to create a porous nanosheet morphology.36 By taminants were observed that were likely due to the calcina- adding ammonia to TiCl4 and ethylene glycol in the during tion procedure required to convert the amorphous titania to solvothermal condensation, they found that a hierarchical TiO2(B). Liu and co-workers recently reported on the lithia- pore structure was added to the nanosheets. The cyclic vol- tion of mesoporous TiO2(B) microspheres.32 The advantage tammetry studies showed redox peaks similar to other reports of a mesoporous microsphere geometry lies in the ability to for nanostructured TiO2(B) along with a smal amount of anatase increase volumetric energy density due to efficient micro- byproduct. The first discharge capacity at 34 mA/g (0.1 C) sphere packing while maintaining the nanostructure needed charge rate was 332 mAh/g with significant ICL of 25%. The for fast surface lithiation kinetics. The CVs showed broad material showed good high rate capability with highly stable capacitive-like profiles with redox peaks associated with both 220 mAh/g capacity at a 10 C charge rate over 200 cycles.
TiO2(B) and anatase. The 0.1 C rate capacity was steady at They attributed the high rate capability to both the ultrathin 250 mAh/g after a first cycle capacity of 310 mAh/g, and geometry and the pore structure that allowed facile access the material exhibited highly stable charging capacity of between layers that may otherwise stack on top of one 160 mAh/g at 10 C. Recently, Dylla and co-workers reported another and therefore be less accessible to electrolyte and on the influence of mesoporosity on the rate capability of would hinder Liþ diffusion. Most recently, Dyl a and co-workers mesoporously ordered 4"5 nm TiO2(B) nanoparticles.33 Using have reported a combined experimental/theoretical study on TiO2(B) nanoparticles as building blocks and a common di- the lithiation of TiO2(B) nanoparticles and nanosheets. The block copolymer (P123) as a structure directing agent, they experimental dC/dV plots and simulations support significantly were able to control the phase purity of the resultant mesopor- different lithiation/delithiation behavior for nanosheets com- ous thin film structure. They found that the open pore structure pared to nanoparticles due to the elongated geometry of the enhanced the high rate charging capacity significantly over nanosheet crystal structure along the a-axis.37 nontemplated composites of TiO2(B) nanoparticles and showedthrough scan rate dependent studies that the mesoporous 4. Theoretical Studies structures lithiated through a pseudocapacitive mechanism.
Panduwinata and Gale calculated from first principles the Similar to the other mesoporous studies, large capacitive-like thermodynamics of low concentration Liþ site occupancy in CVs were observed with the double peak behavior consistent TiO2(B).38 They found the A2 site to be slightly more favor- with TiO2(B) lithium insertion and 170 mAh/g capacity was able than the C site due to the increased interactions with achieved at a rate of 2.3 C.
neighboring oxygens, five versus two interactions, respec- Nanosheets. The previous sections on various forms of tively. Finally, the A1 site was found to be the least favorable TiO2(B) have shown that the kinetics and capacity for lithia- due to increased Liþ"Ti repulsions and lattice strain induced tion of TiO2(B) are influenced by both size and architectural by the cation. They also pointed out that there was con- control. Ideally then, a flat thin surface of TiO2(B) should siderably less lattice expansion within the TiO2(B) structure provide superior fast lithiation behavior since the entirety of compared to anatase or rutile which undergoes significant the material would be considered surface. Xiang and co- anisotropic expansions upon Liþ insertion. Liþ diffusion in workers reported the first synthesis of ultrathin TiO2(B) the dilute limit was also explored, and diffusion from C sites nanosheets. Using a simple hydrothermal reaction between to A2 sites was found to be lowest in energy followed by A2 ethylene glycol and TiCl3 they were able to produce high to C and A1 to C. These three lowest energy diffusion barriers purity, atomically thin sheets with nanocrystalline domains all occur in either the a- or c-axis direction which are known throughout the 200"300 nm continuous sheets.34 to be parallel to radial diffusion in TiO2(B) nanotubes and Jang and co-workers were able to produce nanosheet- nanowires. These calculations correlate well with the experi- derived TiO2(B) that showed promising fast charge rate capa- mental results showing increased rate capability for the city by condensing a layer-structured titanate (H2Ti4O9) into a nanotube geometry that allows for radial diffusion through tunnel-structured titanate (H2Ti8O17) by heating to 200 !C the a- and c-axes. Contrary to these results, Arrouvel and Vol. 46, No. 5 ' 2013 ' 1104–1112 ' ACCOUNTS OF CHEMICAL RESEARCH ' 1109 Lithium Insertion in TiO2(B) Architectures Dylla et al.
co-workers used DFT to calculate an off-centered C site asbeing most stable at low Liþ concentrations followed by A2and A1 site occupation.39 They also calculated diffusionalong the b-axis to be most favorable although they onlyconsidered C-to-C site diffusion, which again is in contra-diction to the calculations of Panduwinata and Gale.
In a combined DFT and experimental study, Islam and co- workers found C-site binding to be preferential followed byA2 and A1. However, speculation of anatase impuritiesmixed with the TiO2(B) used in this study may have led toan incorrect interpretation of the data.40 They also arguedthat the curved surfaces of these nanostructures would be ahindrance to diffusion into the interior of the structure due to FIGURE 6. Differential capacity plots of TiO2(B) nanoparticles (NP-top) compression in the a"b plane as Li+ diffused deeper and nanosheets (NS-bottom) at 50 mA/g charge rate along with anoverlay of DFTþU calculated lithiation/delithiation potentials and into radius of the 1-D structures. Using DFT calculations, relative Liþ concentrations (height) for the lithiation/delithiation Koudriachova showed A1 site absorption on (001) surfaces process. The histogram colors correlate to the intercalation sites being to be favorable and connected to subsurface A2 sites via filled: blue = A1, green = A2, and yellow = C. Reprinted with permission low energy diffusion pathways in the c-axis direction until from ref 37. Copyright 2012 American Chemical Society.
all of the A2 sites had been filled, leading to 0.5 Liþ/Ti.41 lithiation, all of the A2 sites are filled, followed by alternating This low energy diffusion pathway was thought to contri- A1 sites until x = 0.75. Similar to the work by Dalton and co- bute to the pseudocapacitive behavior of TiO2(B) in nano- workers,42 the symmetric C site was found to be unstable wires and nanotubes which have a stretched surface due to Liþ"Liþ repulsions from neighboring A2 sites and a allowing for minimization of Liþ"Ti interactions from the new site was found shifted in the b-axis direction though was A1 surface sites to the A2 subsurface sites. Recently, Dalton found to be filled at extremely low voltages out of the range and co-workers used DFT in combination with Monte Carlo typically studied experimentally. In contrast, TiO2 na- simulations to predict Liþ site occupancies as well as voltage nosheets were calculated to favorably lithiate at the C sites curves and phase diagrams.42 Their calculations showed A1 at low concentrations followed by A2 sites and A1 sites. The to be the most stable site at low Liþ concentrations followed difference in calculated site filling between the 2-D and 3-D by A2. The C site was shifted in the b-axis direction but was nanostructures was attributed to the relaxed calculated still highly unfavorable compared to A1 and A2 site occu- TiO2(B) nanosheet architecture that allows for decreased pancy. Upon filling to x = 0.25, A1 sites are filled on different Liþ"Liþ repulsions at A2 and C sites allowing for favorable planes in the c-axis direction and all of A1 sites are fil ed at x = C site binding and a shifting in the A2 site away from the ideal 0.5. At x = 0.75, an inversion occurs that causes A2 and C sites axial position. Figure 6 shows a comparison of dC/dV plots to be fil ed in the (001) plane. Final y, at x = 1.0, all of the A2 and along with histogram overlays of the DFTþU derived Liþ site C sites are filled.
occupancy voltages and relative concentrations.
In the previously mentioned theoretical studies, the en- ergy penalty (þU) was not included in the calculations.
5. Challenges and Future Prospects Morgan and Watson recently showed that DFT alone is While the high surface energy architectures such as nano- inadequate to describe Liþ binding in anatase because of tubes and nanoparticles provide facile lithiation kinetics, the delocalization of 3d electrons over all Ti centers.43 Only they are often coupled with ICL associated with electrolyte by adding the energy penalty that forces 3d electrons to degradation occurring on those same high energy sites. The localize at Ti centers does the theory predict the expected Liþ nanosheet geometry may be the best option to take advan- site occupation behavior. Most, recently Dylla and co-work- tage of both the high capacity and rates offered by these ers showed in a combined experimental and DFTþU study nanoarchitectures while avoiding significant ICL as the sur- that site filling preferences change as a function of face strain effect should be lessened for the flexible sheets.
nanoarchitecture.37 For 3-D structures such as bulk and Furthermore, any nanoarchitecture used in a battery where nanoparticles, the preference for filling follows A2 > C > A1 volumetric power or energy density is of concern; that is, at the dilute limit with A2 and C being close in energy. Upon mobile and vehicle applications will require a hierarchical 1110 ' ACCOUNTS OF CHEMICAL RESEARCH ' 1104–1112 ' 2013 ' Vol. 46, No. 5 Lithium Insertion in TiO2(B) Architectures Dylla et al.
assembly of nanomaterials into larger micrometer-sized 5 Conway, B. E. Transition from supercapacitor to battery behavior in electrochemical structures in order to increase tap and volumetric energy energy-storage. J. Electrochem. Soc. 1991, 138, 1539–1548.
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as shown above by Brutti and co-workers in the case of 9 Wang, J.; Polleux, J.; Lim, J.; Dunn, B. Pseudocapacitive Contributions to Electrochemical Energy Storage in TiO2 (Anatase) Nanoparticles. J. Phys. Chem. C 2007, 111, 14925– TiO2(B) nanowires.25 Throughout this Account, we have shown that TiO 10 Marchand, R.; Brohan, L.; Tournoux, M. TiO2(B) A new form of titanium-dioxide and the potassium octatitanate K2Ti8O17. Mater. Res. Bull. 1980, 15, 1129–1133.
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understanding of why nanostructuring is so influential of the Electronic and local structural changes with lithium-ion insertion in TiO2-B: X-ray absorption spectroscopy study. J. Mater. Chem. 2011, 21, 15369.
high rate capability and high capacity is still an open ques- 12 Wilkening, M.; Heine, J.; Lyness, C.; Armstrong, A.; Bruce, P. Li diffusion properties of tion. It appears that more surface analytical studies and mixed conducting TiO2-B nanowires. Phys. Rev. B 2009, 80, 064302.
13 Wilkening, M.; Lyness, C.; Armstrong, A. R.; Bruce, P. G. Diffusion in Confined Dimensions: calculations of surface Liþ binding on various TiO2(B) sur- Liþ Transport in Mixed Conducting TiO2"B Nanowires. J. Phys. Chem. C 2009, 113, faces are warranted as the majority of the charging behavior 14 Zachauchristiansen, B.; West, K.; Jacobsen, T.; Atlung, S. Lithium insertion in different TiO2 seems to be controlled by surface interactions rather than modifications. Solid State Ionics 1988, 28, 1176–1182.
15 Zachauchristiansen, B.; West, K.; Jacobsen, T.; Skaarup, S. Lithium insertion in isomorphous MO2(B) structures. Solid State Ionics 1992, 53, 364–369.
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BIOGRAPHICAL INFORMATION 17 Inaba, M.; Oba, Y.; Niina, F.; Murota, Y.; Ogino, Y.; Tasaka, A.; Hirota, K. TiO2(B) as a promising high potential negative electrode for large-size lithium-ion batteries. J. Power Anthony Dylla received his Ph.D. from the University of Maryland Sources 2009, 189, 580–584.
in 2009 on the synthesis, characterization, and catalytic properties 18 Zukalova, M.; Kalbac, M.; Kavan, L.; Exnar, I.; Graetzel, M. Pseudocapacitive lithium storage of bimetallic nanoparticles. He is currently a postdoctoral fellow at in TiO2(B). Chem. Mater. 2005, 17, 1248–1255.
the University of Texas at Austin. His research interests are in 19 Armstrong, A. R.; Armstrong, G.; Canales, J.; Bruce, P. G. TiO2-B Nanowires. Angew.
Chem., Int. Ed. 2004, 43, 2286–2288.
nanostructured metal oxides and the use of in situ vibrational 20 Armstrong, A.; Armstrong, G.; Canales, J.; Bruce, P. TiO spectroscopy to study materials relevant to energy storage.
2-B nanowires as negative electrodes for rechargeable lithium batteries. J. Power Sources 2005, 146, Graeme Henkelman received his Ph.D. from the University of 21 Armstrong, G.; Armstrong, A. R.; Canales, J.; Bruce, P. G. Nanotubes with the TiO Washington in 2001. He was a postdoctorate fellow at Los Alamos structure. Chem. Commun. 2005, 2454.
National Laboratory from 2002 to 2004 and is now associate 22 Daoud, W. A.; Pang, G. K. H. Direct Synthesis of Nanowires with Anatase and TiO2-B professor at the University of Texas at Austin. His research interests Structures at near Ambient Conditions. J. Phys. Chem. B 2010, 1–5.
include DFT and Monte Carlo studies of catalytic and energy 23 Qiao, Y.; Hu, X.; Huang, Y. Microwave-induced solid-state synthesis of TiO2(B) nanobelts with enhanced lithium-storage properties. J. Nanopart. Res. 2012, 14, 1–7.
storage related materials.
24 Armstrong, G.; Armstrong, A. R.; Canales, J.; Bruce, P. G. TiO2(B) Nanotubes as Keith Stevenson received his Ph.D. from the University of Utah Negative Electrodes for Rechargeable Lithium Batteries. Electrochem. Solid St. 2006, 9, in 1997. He was a postdoctorate fellow at Northwestern University 25 Brutti, S.; Gentili, V.; Menard, H.; Scrosati, B.; Bruce, P. G. TiO from 1997 to 2000 and is now full professor at the University of 2-(B) Nanotubes as Anodes for Lithium Batteries: Origin and Mitigation of Irreversible Capacity. Adv. Energy Mater.
Texas at Austin. His research interests include nanostructured 2012, 2, 322–327.
materials for energy storage, electroanalytical and surface chem- 26 Kobayashi, M.; Petrykin, V. V.; Kakihana, M. One-step synthesis of TiO2(B) nanoparticles from a 'Water-Soluble titanium complex. Chem. Mater. 2007, 19, 5373 istry, electrochemical sensors, and development of high-resolution 27 Ren, Y.; Bruce, P. G. Nanoparticulate TiO analytical tools and methods.
2(B): An Anode for Lithium-Ion Batteries. Angew.
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28 Brezesinski, T.; Wang, J.; Polleux, J.; Dunn, B.; Tolbert, S. H. Templated Nanocrystal-Based Porous TiO2 Films for Next-Generation Electrochemical Capacitors. J. Am. Chem. Soc.
2009, 131, 1802–1809.
*To whom correspondence should be addressed. E-mail: [email protected]
The authors declare no competing financial interest.
29 Lin, Y. M.; Abel, P. R.; Flaherty, D. W.; Wu, J.; Stevenson, K. J.; Heller, A.; Mullins, C. B. Morphology Dependence of the Lithium Storage Capability and Rate Performance of Amorphous TiO2 Electrodes. J. Phys. Chem. C 2011, 115, 30 Kavan, L.; Rathousk!y, J.; Gr€atzel, M.; Shklover, V.; Zukal, A. Surfactant-Templated TiO 1 Bruce, P. G.; Scrosati, B.; Tarascon, J.-M. Nanomaterials for Rechargeable Lithium (Anatase): Characteristic Features of Lithium Insertion Electrochemistry in Organized Batteries. Angew. Chem., Int. Ed. 2008, 47, 2930–2946.
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Book us supreme 11-1639 appendix complete
OF THIS UNITED STATES OF AMERICA James Frank Osterbur The UNITED STATES OF AMERICA and these defendants: the president Barack Obama US attorney general Eric H. Holder jr. US solicitor general Neal K. Katyal the internal revenue service Federal Bureau of Investigation On petition for a writ of Certiorari to this United States court of appeals, 7th circuit Chicago, IL
MSDS#: KIP070102-PPC Material Safety Data Sheet Page 1 of 1 CRYSTAL NAILS MASTER POWDER CLEAR Section 1 – Identification of the Substance/Preparation and of the Company/Undertakin Material/Product Name: Master Powder Clear MSDS Initial Approval Date: Chemical Name: N/A MSDS Prepared by: BSQ