ARTICLESPUBLISHED ONLINE: 10 JANUARY 2010 | DOI: 10.1038/NMAT2612

Ordered mesoporous α-MoO3 with iso-orientednanocrystalline walls for thin-film pseudocapacitorsTorsten Brezesinski1,2, JohnWang3*, Sarah H. Tolbert1† and Bruce Dunn3†

Capacitive energy storage is distinguished from other types of electrochemical energy storage by short charging times and theability to deliver significantly more power than batteries. A key limitation to this technology is its low energy density and for thisreason there is considerable interest in exploring pseudocapacitive materials where faradaic mechanisms offer increased levelsof energy storage. Here we show that the capacitive charge-storage properties of mesoporous films of iso-oriented α-MoO3

are superior to those of either mesoporous amorphous material or non-porous crystalline MoO3. Whereas both crystallineand amorphous mesoporous materials show redox pseudocapacitance, the iso-oriented layered crystalline domains enablelithium ions to be inserted into the van der Waals gaps of the α-MoO3. We propose that this extra contribution arises froman intercalation pseudocapacitance, which occurs on the same timescale as redox pseudocapacitance. The result is increasedcharge-storage capacity without compromising charge/discharge kinetics in mesoporous crystalline MoO3.

Transition metal oxides with layered crystalline structures arevery attractive materials for charge storage because of theirability to intercalate ions in a wide range of sites1,2. Both

faradaic and non-faradaic mechanisms can store charge. The latteris the dominant factor in electric double-layer capacitors, in whichcharge is stored electrostatically from the reversible adsorption ofions onto high-surface-area materials3,4. The other option is tostore charge in a faradaic process, which again can take place ona surface, as in pseudocapacitors, or in the bulk of the material, asin conventional batteries5–7. Here, we explore the electrochemistryof nanostructured versions of layered transitionmetal oxides with aspecific goal of understanding how to increase faradaic capacitance.

Capacitive charge storage offers a number of desirable propertieswhen compared with conventional batteries, including chargingwithin seconds, long-term cycling stability and the ability to deliverup to ten times more power than batteries. These features aredesirable for a range of applications8. Whereas charge storagethrough non-faradaic means is fairly straightforward, a variety ofdistinct chemical options are available for charge storage throughfaradaic or redox-type reactions. The following reactions arepossible when a material is immersed in an electrolyte. (1) Cationscan react with the electroactive material, which then undergoeschanges to form different compositions or phases4,5. This typeof chemistry is common for primary batteries. (2) Cations canintercalate into interlayer gaps (van der Waals gaps) of layered ortunnel materials9. This intercalation is accompanied by reductionof metal ions in the walls to maintain charge neutrality. (3) Finally,cations can electrochemically adsorb onto the surface of a materialthrough charge-transfer processes7. Chemically, this process is verysimilar to option (2), in that cation adsorption is associated withmetal reduction. Option (3) alleviates the need for long-rangediffusion of ions through the van der Waals gaps, however, sothis process is kinetically more facile and is referred to as redoxpseudocapacitance. Intercalation into van der Waals gaps is oftenquite slow. In some cases, however, diffusion into the van derWaalsgaps can be fairly facile, and in this case the intercalation can also be

1Department of Chemistry and Biochemistry, University of California at Los Angeles, California 90095-1569, USA, 2Institute of Physical Chemistry,Justus-Liebig-University Giessen, Heinrich-Buff-Ring 58, 35392 Giessen, Germany, 3Department of Materials Science and Engineering, University ofCalifornia at Los Angeles, California 90095-1595, USA. *Present address: HRL Laboratories, LLC, Malibu, California 90265, USA.†e-mail: tolbert@chem.ucla.edu; bdunn@ucla.edu.

considered capacitive (known as intercalation pseudocapacitance).According to Conway and co-workers, these electrochemicalintercalation processes are pseudocapacitive in nature becausethe cations are faradaically stored and phase transitions (thatis, structural rearrangements) do not occur7,9. Perhaps a morepractical definition is that the reactions should be kinetically facileand not diffusion controlled on the time scale of interest. Atpresent there is widespread interest in using pseudocapacitor-basedmaterials for electrochemical capacitors because the energy densityassociated with faradaic reactions is substantially larger, by at leastone order of magnitude, than electric double-layer capacitors.Here, our goal is to develop nanostructured architectures fortransition metal oxides that allow for both redox and intercalationpseudocapacitance. In so doing, we have the potential to optimizeboth energy density and power density in a singlematerial.

As electrochemically active materials approach nanoscaledimensions, pseudocapacitive effects become increasinglyimportant10,11. In a recent report, we quantitatively establishedthe dependence of pseudocapacitance on the particle size ofanatase TiO2 nanocrystals12. The results showed that for particlessmaller than 10 nm in diameter, more than half of the total storedcharge is capacitive. As anatase TiO2 has no van der Waals gap,all of its capacitance is believed to be redox pseudocapacitance.Other studies have shown that materials with layered crystallinestructures such as TiO2(B), TiS2 and H2Ti3O7 are able to ac-commodate significant amounts of Li+ through pseudocapacitiveprocesses6,9,13,14. Although no distinction was made between redoxand intercalation pseudocapacitance, we expect that there was somecontribution from the latter in these materials. One reason for thishypothesis is that a pseudocapacitive response was identified atsweep rates that were much slower than one would expect for redoxpseudocapacitance13,14.

One way to increase pseudocapacitance is to increase surfacearea, as described above for TiO2 nanocrystals12. This increasecan be maintained in thick-film format by creating a hierarchicalarchitecture where nanocrystal building blocks assemble into a

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NATURE MATERIALS DOI: 10.1038/NMAT2612 ARTICLESnanoporous network using polymer templating methods15. Suchmaterials show capacitive charge-storage levels comparable to thatobserved in nanoparticle films. To further increase capacitive chargestorage, one would like to use a material that has the potentialto show both redox and intercalation pseudocapacitance, ideallywith similarly facile kinetics. Here, we demonstrate that enhancedpseudocapacitance can be achieved using thin-film materials withinterconnected mesoporosity and highly oriented crystalline wallswith van der Waals gaps. These materials are based on the co-assembly of inorganic oligomers with organic templates using anevaporation-induced self-assembly process16–19. In particular, weuse cubic mesoporous molybdenum oxide (MoO3) thin films withiso-oriented nanocrystalline grains. These films were producedusing a poly(ethylene-co-butylene)-b-poly(ethylene oxide) blockcopolymer (also referred to as KLE), which has been shownto strongly facilitate the conversion of the initially amorphousframework into a highly crystallinematerial20–22.

Molybdite (α-MoO3) is an electroactive two-dimensional (2D)layered material that can accommodate up to 1.5 Li/Mo (ref. 23).It comprises alternately stacked layers held together by weak vander Waals forces along [010]. The interlayer gaps can be usedfor Li+ insertion; however, MoO3 has not proven to be of muchinterest for battery applications because of modest reaction kineticsand poor cycling behaviour2,24. We are unaware of any studiesof MoO3 as a capacitor material. Here, we characterize capacitivecharge storage in mesoporous α-MoO3 and demonstrate thatthese films show enhanced pseudocapacitance compared with bothmesoporous amorphous and non-templated crystallineMoO3. Ourresults indicate that the insertion of Li+ in the quasi-2D planes, thatis, intercalation pseudocapacitance, is responsible for the high levelof capacitive charge storage.

Nanoscale structureThe mesoporous MoO3 films used in this work were synthesizedusing an evaporation-induced self-assembly process. The synthesisof these materials was reported previously and so this sectionwill focus on extending the characterization of these materials25.In the synthesis, an ethanolic solution containing a sol–gelprecursor and structure-directing agent is dip-coated onto a polarsubstrate. On evaporation of the solvent, the system co-assemblesto form a mesostructured inorganic/organic composite. Withthermal treatment, this composite develops a unique pore–solidarchitecture as described below.

Figure 1a shows low- and high-magnification transmissionelectron microscope (TEM) images of a self-organized MoO3 thinfilm after removal of the KLE template and crystallization at 450 ◦C.A porous cubic architecture without major structural defects canbe observed with 13-nm-diameter pores and an average pitch of∼23 nm. Figure 1b is an electron diffraction pattern taken fromthe same film, which shows spots instead of Debye–Scherrer rings.This pattern is characteristic of a system composed of iso-orientedmolybdite grains25. The homogeneity of the α-MoO3 films is alsoobserved with atomic force microscopy (AFM, Fig. 1c), whichemphasizes both the flatness of the top surface and the fact that thepores at the surface are open. Finally, synchrotron-based grazing-incidence small-angle X-ray scattering (GISAXS, Fig. 1d) showsdistinct spots that can be indexed to a highly ordered face-centred-cubic pore architecture with (111) orientation, indicating that theperiodic structure observed on the film surface persists throughoutthe film. More detail on evolution of the GISAXS with thermaltreatment can be found in SupplementaryDiscussion S1a.

The oriented nature of the MoO3 domains seems to beimportant for electrochemical performance, and so we examinethe oriented crystallization in more detail in Fig. 2a–c usingwide-angle X-ray diffraction (WAXD). The crystallization of theamorphous framework occurs at around 400 ◦C. The atomic

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Figure 1 | Morphology of mesoporous α-MoO3 with highly orientedcrystalline walls. a, Low-magnification bright-field TEM micrograph.Inset: A higher magnification micrograph. b, Selected-area electrondiffraction pattern showing diffraction spots characteristic of iso-orientedmolybdite oriented along the (110) direction. c, Tapping-mode AFM imageof the top surface. The root-mean-square roughness of this sample is lessthan 1 nm. Inset: A 3D-AFM image at a tilt of 50◦; the contrast coversheight variations in the 1–5 nm range. d, GISAXS pattern obtained at anangle of incidence, β , of 0.2◦ on a face-centred-cubic mesoporous α-MoO3

film calcined at 400 ◦C in oxygen for 2 min. Scattering vector s componentsare given in units of 1 nm−1; s= 2/λsinθ .

structure of α-MoO3 can be described as layers of MoO6 octahedranormal to the b axis (Fig. 2e). X-ray photoelectron spectroscopy(XPS) obtained on crystalline samples shows nearly perfect 1:3Mo/O stoichiometry and the presence of only Mo6+ (Fig. 2d andSupplementary Discussion S1b). At 400 ◦C, Raman spectroscopyindicates the presence of very small amounts of the β-MoO3phase; however, only the α-MoO3 phase is observed for samplesheated above this temperature (Supplementary Discussion S1c).1D and 2D WAXD for samples heated to 400 and 450 ◦C(Fig. 2a–c) shows only the (0k0) peaks of α-MoO3 in the directionnormal to the substrate (JCPDS reference card no. 5-0508).Like the electron diffraction presented in Fig. 1b, these datathus indicate that the layers shown in Fig. 2e lie parallel to theplane of the substrate. Applying the Warren–Averbach approachto 1D diffraction patterns collected on samples heated above400 ◦C, an average nanocrystal size of 40–45 nm perpendicularto the substrate is obtained26. This implies the fairly remarkableconclusion that a single nanocrystal forms the wall over a distanceencompassing several mesopores. Taken together, these resultsshow that calcination at 450 ◦C is optimal for phase purityand grain orientation.

Electrochemical propertiesVoltammetry experiments carried out at various sweep rates wereused to examine the redox processes occurring in the sol–gel derivedfilms. The electrochemical Li+ insertion process is represented bythe following equation:

MoO3 (colourless)+xLi++xe−↔ LixMoO3 (grey–blue)

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ARTICLES NATURE MATERIALS DOI: 10.1038/NMAT2612

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Figure 2 | Structure and chemical characterization of mesoporous α-MoO3 films. a, 1D WAXD patterns obtained on films calcined at 350 (A), 400 (B)and 450 ◦C (C) in oxygen for 2 min. The line pattern shows reference card no. 5-0508 for orthorhombic MoO3 (molybdite) according to JCPDS.b,c, 2D WAXD patterns collected after treatment of films at 400 (b) and 450 ◦C (c). At 450 ◦C, some broadening of the arcs in 2D WAXD indicatesreduced crystallographic orientation. d, Typical XPS survey spectrum for films heated above 400 ◦C in oxygen. Inset: A high-resolution scan of the Mo 3dregion. e, Molybdite unit cells showing the layer-like arrangement of molybdenum (green) and oxygen atoms (red).

with the mole fraction of inserted lithium given by x =QM/mF .Here, Q is the stored charge, M is the molecular weight, mis the mass and F is the Faraday constant. Figure 3a showscyclic voltammograms for mesoporous amorphous and crystallinefilms as well as for a crystalline film produced without anypolymer template but under otherwise identical conditions. Wenote that although this non-templated film does not have regularmesoporosity, it is solution processed and thus has some smallfractional porosity. These three curves, each of which is reversible,indicate that theMoO3 films used in this work possess very differentinsertion/extraction properties. The mesoporous amorphous filmsreveal a featureless voltammetric response over a broad potentialrange. This response arises from the disorder of the amorphous porewall structure, which leads to numerous sites spread over a widevariation in energies27. In contrast, themesoporous crystalline filmsshow well-defined redox peaks. The cathodic peak at around 2.7 Vversus Li/Li+ can be attributed to an irreversible phase transition(α-MoO3 into Li0.25MoO3, indicated by 1 in Fig. 3a), which hasbeen shown to cause a pronounced expansion of the interlayerspacing23,24,28. The voltammetric response at lower potentials ischaracterized by two pairs of reversible redox peaks (indicated by2/2′ and 3/3′ for reduction and oxidation, respectively, in Fig. 3a).According to previous reports, these charge-insertion sites can bedescribed as intralayer (2/2′) and interlayer (3/3′), respectively23,24.The non-templated crystalline films, on the other hand, produceonly broad peaks. This type of response is characteristic of denseMoO3 materials and indicates sluggish ion insertion/extractionkinetics owing to slow diffusion of Li+ through the bulk of

the film. These results are consistent with previous studies onMoO3 films prepared by electrodeposition and heat treated atvarious temperatures29.

Figure 3a further shows that the amount of stored chargeis strongly dependent on film morphology. At a sweep rate of1mV s−1, the charge storage is calculated to be 605 and 330C g−1for mesoporous crystalline and amorphous films, respectively.As porosity (30%) and Brunauer–Emmet–Teller surface area(200m2 g−1) are roughly the same for both mesoporous films,this result suggests that for the same dynamic conditions, a muchgreater level of charge storage is achieved with the mesoporouscrystalline materials. By comparison, non-templated films showa charge storage of only 215C g−1 at the same sweep rate,which clearly demonstrates that the introduction of interconnectedporosity significantly enhances electrochemical properties as thecombination of electrolyte access and short diffusion path lengthsfor both electronic and Li+ transport provides a beneficialmicrostructure for Li+ insertion/extraction.

We are particularly interested in determining the chargingbehaviour of mesoporous materials at short times because a keyadvantage of electrochemical capacitors over batteries is the abilityto have faster charging and to deliver high power. As shown inFig. 3b, the total amount of stored charge is observed to be stronglytime dependent. Regardless of the charging time, however, the totalcharge stored in mesoporous crystalline films is significantly highercompared with mesoporous amorphous ones. The non-templatedcrystalline films store even less charge than the mesoporousamorphous material (data not shown).

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NATURE MATERIALS DOI: 10.1038/NMAT2612 ARTICLES

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Figure 3 | Electrochemical characterization of sol–gel derived MoO3 films. a, Cyclic voltammograms for different MoO3 systems at a sweep rate of1 mV s−1. A refers to the mesoporous amorphous film and C refers to the mesoporous crystalline films (that is, α-MoO3). N is a non-templated crystallinefilm. b, Kinetic behaviour of the various mesoporous films. The total amount of charge storage (gravimetrically normalized) as a function of charging timefor crystalline MoO3 is much greater than the corresponding mesoporous amorphous films. c, The capacitive contribution to the total charge stored isplotted as function of charging time for the same films as in b. The capacitive contribution for the mesoporous crystalline films is three times larger thanthat of the mesoporous amorphous material.

The total stored charge in Fig. 3b can be separated intothree components: the faradaic contribution from diffusion-controlled Li+ insertion5, the faradaic contribution from charge-transfer processes with surface atoms and, possibly, also withatoms located in the interlayer lattice planes (referred to aspseudocapacitance)7, and the non-faradaic contribution from thedouble-layer capacitance3,4. Following our previously reportedapproach, we are able to use the voltammetric sweep ratedependence to determine quantitatively the capacitive contributionto the current response (Supplementary Discussion S2a)12,15. Thecurrent response (i) at a fixed potential (V ) can be described as thecombination of two separate mechanisms, namely capacitive effects(k1v) and diffusion-controlled insertion (k2v1/2) according to30:

i(V )= k1v+k2v1/2

where v is the sweep rate. By determining both k1 and k2, it is thuspossible to distinguish the fraction of the current arising from Li+insertion and that from capacitive processes at specific potentials.

In Fig. 3c we have used this analysis to determine the fractionof the total stored charge from capacitive processes as a functionof charging time. The striking similarity between Fig. 3b andc indicates that much of the increase in charge storage in themesoporous crystalline material compared with the amorphoussample is related to capacitive charge storage. This capacitivecontribution occurs only over the first three to four minutes,however. Thereafter, any further increase in stored charge shown inFig. 3b arises fromdiffusion-controlled Li+ intercalation.

The second significant feature shown in Fig. 3c is the magnitudeof the capacitive charge storage for the different mesoporous MoO3films. Capacitive charge storage for the mesoporous crystalline film(450C g−1) is some three times larger than that of the correspondingmesoporous amorphous film (150C g−1). In separate voltammetryexperiments designed to determine the double-layer capacitanceby using a bulky electrolyte, measurements indicate that chargestorage from the electrical double layer is in the range of 10–20C g−1for both crystalline and amorphous samples (SupplementaryDiscussion S2b). This value is much smaller than that obtainedfor either type of mesoporous film. We thus conclude thatthe difference between the two mesoporous materials arisesfrom differences in pseudocapacitive processes. Although redoxpseudocapacitance is expected to occur in both amorphous andcrystalline samples, there is no structural reason why crystallinematerials should show substantially larger values. Accordingly,we contend that mesoporous α-MoO3 with iso-oriented layerednanocrystalline grains enables intercalation pseudocapacitance tooccur on the same timescale as redox pseudocapacitance. This

behaviour differs significantly from the pseudocapacitance reportedin previous studies on TiO2(B) and protonated titanates thatwere carried out at relatively slow sweep rates (<2mV s−1;refs 13, 14, 31). At charging times longer than four minutes, thediffusion-controlled Li+ insertion capacities are similar for bothtypes of film (Fig. 3b).

In Fig. 4a,b, the voltage profile for the capacitive current (shadedregion) is compared with the total measured current obtainedat a sweep rate of 0.1mV s−1 (charging time of 20,000 s) forthe mesoporous crystalline and amorphous films. The capacitivecontribution for the mesoporous crystalline films is significantlyhigher than that for the amorphous ones. For α-MoO3, thecapacitive contribution is 70%, whereas only about 35% of the totalstored charge is capacitive in the case of mesoporous amorphousfilms. As discussed above, we can attribute the capacitive currentto be primarily from pseudocapacitive contributions and not fromdouble-layer capacitance. The voltage profile in Fig. 4a also revealsa broad peak in the capacitive curve for the mesoporous α-MoO3samples. This is characteristic of a surface-confined charge-transferprocess, which again indicates that the stored charge is mainly pseu-docapacitive in nature. Similar behaviour has also been reported forother hostmaterials with layered crystalline structure13,31.

To explain our experimental results, it is important toconsider what there is about the structure of the templatedcrystalline films compared with the amorphous one that exertssuch a significant influence on charge storage. This considerationis closely related to our goal of designing a material whereintercalation pseudocapacitance could be used to augment themore conventional redox pseudocapacitance. Taking into accountthe Brunauer–Emmet–Teller surface area of 200m2 g−1, thesurface-normalized capacitance is calculated to be approximately150 µF cm−2 for mesoporous α-MoO3 films. This value confirmsthat the double-layer capacitance determined for the mesoporousα-MoO3, ∼20 µF cm−2, represents a relatively small contributionto the total charge storage7. Interestingly, 150 µF cm−2 of chargecorresponds approximately to 4.5 e− per surface molybdenumassuming only redox pseudocapacitance. This value is chemicallyunreasonable and further argues that surface pseudocapacitance isnot the onlymechanismworking in thesematerials. By comparison,mesoporous amorphous films show a surface-area normalizedcapacitance of 50 µF cm−2, which, after subtracting the double-layer capacitance, corresponds to a very reasonable 1.3 e− persurface molybdenum. This analysis suggests that the amorphousstructure can provide surface sites for redox pseudocapacitance,but does not possess sufficient order to enable Li+ access tothe region below the surface in the time frame of interest. Itwould seem, therefore, that the facile ability of Li+ to be insertedin the interlayer gap and carry out charge-transfer reactions,

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ARTICLES NATURE MATERIALS DOI: 10.1038/NMAT2612

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Figure 4 | Capacitive and diffusion-controlled contributions to chargestorage. Voltammetric responses for mesoporous MoO3 films at a sweeprate of 0.1 mV s−1. The capacitive contribution to the total current is shownby the shaded region. a,b, For mesoporous α-MoO3 (a), some 70% of thetotal charge is capacitive, whereas for the mesoporous amorphous MoO3

(b), only 35% is capacitive.

that is, intercalation pseudocapacitance, represents an extrapseudocapacitivemechanism that accounts for the enhanced chargestorage ofmesoporousα-MoO3 (SupplementaryDiscussion S2c).

Two factors are expected to contribute to the substantial levelof intercalation pseudocapacitance in mesoporous α-MoO3 films.The nanoscale dimensions of the material provide for very shortdiffusion path lengths5. Perhaps more important, however, is theporosity and the iso-oriented nature of the material. In almostall cathode materials, some volume expansion is required on Li+insertion, and this expansion usually involves an increase in thewidth of the van der Waals gaps in layered metal oxide systems.For thin-film materials, expansion in the plane of the film is notpossible because of substrate pinning, but expansion normal tothe plane is facile. With the iso-oriented nature of these materials,however, in-plane expansion is not needed to accommodate thestrain of Li+ insertion because all expansion occurs normal tothe substrate, a feature that should markedly reduce the stressassociated with Li+ insertion. Although it might seem that thisadvantage could be found equally well in an iso-oriented bulksystem, it is important to remember that ions enter the van derWaals gaps from the edges. Nanoscale porosity enables electrolyte toaccess the interior of the film and exposes the edges of the layers forfacile Li+ insertion. Moreover, previous experiments on templatedoxide systems have shown that the pore architecture can flex quiteeasily in response to a mechanical load, much like a nanoscaletruss, further facilitating this volume expansion32. This unique

combination of factors is probably responsible for the unusually fastintercalation pseudocapacitance observed in thesematerials.

The results presented here illustrate the benefits of both amesoporous morphology and the unique properties of crystallo-graphically oriented films. Redox pseudocapacitance benefits fromthe mesoporous morphology by enabling electrolyte access tothe electrochemical sites and short diffusion path lengths. Theiso-oriented crystalline domains and the flexible nanoporous archi-tecture, however, enable another capacitive contribution to arise,that of intercalation pseudocapacitance. This process is associatedwith the insertion of Li+ in the quasi-2D van der Waals gaps of theα-MoO3. We find that this intercalation pseudocapacitancemecha-nism occurs at rates comparable to traditional redox pseudocapaci-tance in porous MoO3 films because of the unique nanoscale archi-tecture. The net result is that mesoporous films with iso-orientedgrains lead to new pseudocapacitive materials that offer increasedenergy density while still maintaining their high-power-densityadvantages. The potential for practical advances in energy storageusing thesematerials and these general design rules is exciting.

MethodsSynthesis. In a water-free container, 0.08 g of H(CH2CH2CH2(CH)CH2CH3)89(OCH2CH2)79OH (also referred to as KLE) dissolved in 4ml of ethanol iscombined with 0.4 g of MoCl5. Once the solution is homogeneous, 0.2 g ofdouble-distilled H2O is added. Thin films are produced by means of dip-coatingon polar substrates. Optimal conditions include relative humidities below 25% andconstant withdrawal rates of 1–10mm s−1. The films need to be aged at 200 ◦C for12 h before template removal to prevent loss of mesoscale order over the course ofcrystallization. For best results, the films are calcined in oxygen using a 1 h ramp to450 ◦C followed by a 2min soak.

Characterization. TEM was carried out using a Zeiss EM 912� at an accelerationvoltage of 120 kV. Tapping-mode AFM images were taken with a multimodeAFM from Vecco Instruments using Olympus microcantilevers (resonancefrequency: 300 kHz, force constant: 42Nm−1). 1D WAXD measurements werecarried out on a D8 diffractometer from Bruker instruments (Cu Kα radiation)equipped with an energy dispersive solid-state detector. For 2D-WAXD, a set-upwith a 2D MAR image plate detector was used. Both the source (Nonius FR590 generator) and the detector were fixed. GISAXS data were collected at theGerman synchrotron radiation facility HASYLAB/DESY on beamline BW4 using aMAR CCD (charge-coupled device) area detector and a sample–detector distanceof 1,820mm. XPS spectra were acquired on a Physical Electronics ESCA 5600spectrometer with a monochromatic Al Kα X-ray source (power: 200 W/14 kV)and a multi-channel detector OmniIV. The electron takeoff angle θ to the samplesurface was adjusted to 45◦. The adventitious hydrocarbon C1s signal at 284.6 eVwas used as the energy reference to correct for charging. The film thickness wasdetermined with a Vecco Dektak 6M profilometer.

Electrochemical experiments were carried out in a three-electrode cellusing a PAR EG&G 273 Potentiostat. The working electrode was ITO glass onwhich the MoO3 films were deposited. Lithium wire was used as a referenceelectrode. A lithium foil several times the area of the working electrode was usedas the counter electrode. The electrolyte solution was 1.0M LiClO4 in propylenecarbonate. For the double-layer capacitance experiments, the electrolyte waschanged to 1.0M tetrabutylammonium perchlorate in propylene carbonate. Allmeasurements were carried out in an argon-filled glove box using cutoff voltagesat 3.5 and 1.5 V versus Li/Li+.

Received 11 March 2009; accepted 23 November 2009;published online 10 January 2010

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AcknowledgementsThe authors thank K.-I. Iimura, A. Reinecke, K. Brezesinski, J. Perlich and V. Augustynfor their assistance in materials preparation and measurements. This work was supportedby the Office of Naval Research (B.D. and S.H.T.), the National Science Foundationunder grant CHE-0527015 (S.H.T.) and by the Fonds der Chemischen Industrie(Liebig Fellowship, T.B.).

Author contributionsT.B. and J.W.: experimental work and data analysis. S.H.T. and B.D.: projectplanning and data analysis.

Additional informationThe authors declare no competing financial interests. Supplementary informationaccompanies this paper on www.nature.com/naturematerials. Reprints and permissionsinformation is available online at http://npg.nature.com/reprintsandpermissions.Correspondence and requests formaterials should be addressed to S.H.T. or B.D.

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