M Lalanne

A Barnab

F Mathieu

Ph Tailhades

Synthesis and Thermostructural Studies of a CuFe 1-x Cr x O 2 Delafossite Solid Solution with 0 e x e 1

In this work, different CuFe 1-x Cr x O 2 compositions with 0 e x e 1 were prepared by a standard solid-state reaction. These oxides crystallize with the delafossite structure. The phase stability and thermal behavior of the complete CuFe 1-x Cr x O 2 solid solution was studied by thermogravimetric analysis and high-temperature X-ray diffraction experiments under an air atmosphere up to 1000 °C. For x = 0, CuFeO 2 is oxidized into the spinel (CuFe 2 O 4 ) and copper monoxide (CuO) phases, whereas for x = 1, CuCrO 2 is thermally stable. For all of the intermediate compositions (0 < x < 1), complex oxidation, reduction, and phase transitions between delafossite and spinel have been observed. chromium tends to stabilize the stoichiometric delafossite phase, while iron favors the delafossite-to-spinel phase transition.

Introduction

AMO 2 delafossite compounds derived from the CuFeO 2 mineral are quite interesting materials because of their ability to be stabilized with a large number of A and M cations and within a wide range of off-stoichiometry values, leading to different physical properties. 1,2 In particular, delafossite compounds with M = {Al, Ga} have attracted much attention as p-type transparent conducting oxides (TCOs) in the past decade. 3 Nowadays, studies on delafossite oxides are interesting not only for the TCO properties but also for applications such as in catalysis, 4 magnetics, 5 and possibly thermoelectrics. 6,7 The delafossite structure can accommodate monovalent noble metal cations A + = {Cu + , Ag + , Pd + , Pt + } and trivalent cations M 3+ with ionic radii varying from 0.535 A (octahedral Al 3+ ion) up to 1.032 A ˚(octahedral La 3+ ion) according to Shannon's table of effective ionic radii. 8 The AMO 2 delafossite structure can be described as a stacking of MO 2 layers made of edge-sharing MO 6 octahedra that are connected by planes of A + cations arranged as a triangular network (Figure 1). Each A + cation is linearly coordinated by two O -anions belonging to the upper and lower MO 2 layers. The oxygen layers can be stacked in different ways along the c axis in order to form two polytypes of the delafossite structure: the hexagonal 2H (space group P6 3 /mmc) and the rhombohedral 3R (space group R3m) polytypes.

The synthesis 1,9-12 and thermodynamic, 13,14 electric, 15 optoelectric, 16 thermoelectric, 17 lithium intercalation, 18 and magnetic properties 19 of CuMO 2 with M = Fe have been intensively studied so far: CuFeO 2 delafossite is a well-known p-type semiconductor with an antiferromagnetic behavior *To whom correspondence should be addressed. E-mail: barnabe@ chimie.ups-tlse.fr.

(1) Shannon, R. D.; Rogers, D. B.; Prewitt, C. T. Inorg. Chem. 1971, 10, 713-718.

(2) Doumerc, J. P.; Ammar, A.; Wichainchai, A.; Pouchard, M.; Hagenmuller, P. J. Phys. Chem. Solids 1987, 48, 37-43.

(3) Kawazoe, H.; Yasukawa, M.; Hyodo, H.; Kurita, M.; Yanagi, H.; Hosono, H. Nature 1997, 389, 939-942.

(4) Monnier, J. R.; Apai, G. R.; Hanrahan, M. J. U.S. Patent 4,748,144, 1988.

(5) El Ataoui, K.; Doumerc, J. P.; Ammar, A.; Grenier, J. C. below 13 K and frustrated magnetic interactions in triangular lattices. 20 However, n-type single crystals have also been grown and characterized. 21 The conductivity of p-type Cu-FeO 2 delafossite at room temperature (σ RT = 2 S/cm 15,16 ) is the largest among the CuMO 2 delafossite series when an offstoichiometric CuFeO 2+δ delafossite phase is formed. In delafossite compounds, the Cu/M ratio can have an influence on the electrical properties, 22 but the transport properties are mainly governed by the copper mixed valency Cu I /Cu II . 23,24 In CuFeO 2+δ , the Cu I /Cu II ratio is directly controlled by the oxygen nonstoichiometry value δ according to (Cu + 1-2δ Cu 2+ 2δ ) FeO 2+δ . However, for TCO applications, the optical transmittance of CuFeO 2 delafossite thin films is limited in the visible region because of its narrow band gap of about 2 eV.

CuMO 2 with M = Cr has also been studied for its thermodynamic, 13 magnetic, 25 optoelectronic, [26][27][28][29][30] and thermoelectric 31,32 properties, catalytic activity, 4,33 and lithium insertion: 34 CuCrO 2 delafossite is also a p-type semiconductor but with a lower intrinsic electrical conductivity (σ RT = 3.5 Â 10 -5 S/cm 27 ) because of difficulties in inducing an oxygen nonstoichiometric phase in this system. However, CuCrO 2 is a more promising TCO than CuFeO 2 because of good optical transparency in the visible range with a band gap of about 3.1 eV and with an electrical conductivity that can be drastically improved up to σ RT = 220 S/cm with appropriate doping like in a CuCr 0.95 Mg 0.05 O 2 thin film. 35 In this case, the copper mixed valency is directly controlled by the doping stoichiometry x according to (Cu + 1-x Cu 2+

x )-(Cr 3+ 1-x Mg 2+

x )O 2 . One can note that the mineral name of CuCrO 2 is Mac Connellite. It is worth mentioning that CuFeO 2 and CuCrO 2 are isostructural compounds. Henceforth, the general term delafossite will be used to refer to the compound studied in this work.

The present study deals with a CuFe 1-x Cr x O 2 (0 e x e 1) solid solution, which can combine the high intrinsic electrical conductivity of CuFeO 2 because of its ability to obtain offstoichiometric compounds and the good optical properties of CuCrO 2 in order to obtain new TCO materials. The purpose of this work is to synthesize the complete solid solution CuFe 1-x Cr x O 2 with 0 e x e 1 and analyze the effect of oxygen nonstoichiometry on the stability of the delafossite phase using thermostructural analyses.

Experimental Section

Polycrystalline samples of CuFe 1-x Cr x O 2 with 0 e x e 1 were prepared by a standard solid-state reaction. Cu 2 O, Fe 2 O 3 , and Cr 2 O 3 commercial powders are mixed in stoichiometric quantities. The obtained mixtures have been treated in a nitrogen atmosphere between 900 and 1000 °C for 30 h with intermittent grindings.

Powder X-ray diffraction (XRD) measurements for phase analysis and structural refinement were performed at room temperature using a Bruker AXS D4 diffractometer. Struc- tural refinements were made using the Rietveld analysis implemented in the FullProf-WinPlotR program. All samples were also analyzed with respect to temperature up to 1000 °C in air using a Bruker AXS D8 diffractometer equipped with a MRI Radiation high-temperature chamber. In all cases, copper radiations were used as the X-ray source (λ Cu KR 1 = 1.5405 A ˚and λ Cu KR 2 = 1.5445 A ˚) and filtered with a nickel filter.

The specimens for transmission electron microscopy (TEM) were prepared by crushing polycrystalline powders and mounting the crystal fragments onto a lacey-carbon copper grid. TEM analysis was performed with a Jeol 2100F field emission gun electron microscope operating at 200 kV equipped with a side-entry double-tilt specimen holder. The morphology of the CuMO 2 crystals was observed using a Jeol JSM-6400 scanning electron microscope. Energydispersive spectroscopy (EDS) was systematically performed on numerous grains.

The specific surface area of the powder was measured using a 70/30 mixture of helium and nitrogen adsorption by the Brunauer-Emmett-Teller (BET) method. The apparatus used was a Micromeritics Flowsorb II 2300 device. The powders were degassed in a nitrogen atmosphere at 250 °C for 30 min to decontaminate the surface before the measurements.

Thermogravimetric analysis (TGA) of samples was recorded with a TAG 1600 Setaram symmetric thermobalance. The samples were heated at 5 °C/min up to 1000 °C in air.

Results and Discussion

CuFe 1-x Cr x O 2 Solid Solution. After 30 h of thermal treatment at 900 and 1000 °C under a nitrogen atmosphere with intermittent grindings, all of the synthesized samples of the CuFe 1-x Cr x O 2 series with 0 e x e 1 exhibited a single delafossite phase. The complete series of solid solutions CuFe 1-x Cr x O 2 with x ranging from 0 to 1 was thus prepared for the first time. For intermediate composition 0 < x < 1, delafossite phases with lattice parameters in accordance with the iron-and chromiumrich delafossites were obtained after only 10 h of treatment at 900 °C under a nitrogen atmosphere. Longer treatment times and higher temperatures with intermittent grindings were required to obtain a single delafossite solid solution. Variation of the unit cell volume as a result of the thermal treatment is determined from lattice parameter Rietveld refinement as shown in Figure 2. Even though two delafossite phases are present for x = 0.165 and 0.835, they are not represented in Figure 2 because of the weak Bragg peaks in the XRD pattern.

The lattice parameters and unit cell volumes are indicated in Table 1. The XRD pattern for x = 2 / 3 (CuFe 0.333 Cr 0.667 O 2 ) is shown in Figure 3. The unit cell parameters are a = 3.0344(2) A ˚and c = 17.158(3) A ˚for x = 0 (CuFeO 2 ) and a = 2.9742(2) A ˚and c = 17.090(2) A Table for x = 1 (CuCrO 2 ). These values are in excellent agreement with those listed in the literature. 1,2,11 The decrease in the unit cell parameters and cell volume with increasing x in CuFe 1-x Cr x O 2 well agrees with Vegard's law and with the decrease of the ionic radius as iron with r(Fe 3+ coord VI ) = 0.645 A ˚is replaced by chromium with r(Cr 3+ coord VI ) = 0.615 A ˚. This change in the unit cell volume (-4.3% from x = 0 to 1) is rather anisotropic because it is mainly due to the shrinking of the a parameter (-2% from x = 0 to 1), whereas the c parameter remains more or less constant (-0.3% from x = 0 to 1). The O-Cu-O bond lengths along the c axis are quite unaffected by the iron-to-chromium substitution, whereas the a parameter decreases proportionally to the edge of the MO 6 octahedral when x increases. This anisotropy induces a constant decrease of the distortion of the MO 6 octahedra from iron to chromium, which can be correlated to the covalence of the metal-to-oxygen bond. 2 The x = 2 / 3 (CuFe 0.333 Cr 0.667 O 2 ) sample has been investigated by TEM. The electron diffraction (ED) pattern has been indexed in the 3R delafossite structure with a ≈ 3.0 A ˚and c ≈ 17.1 A ˚. As illustrated in Figure 4, corresponding to the [-110] zone axis, a high-resolution transmission electron microscopy (HRTEM) micrograph and the corresponding Fourier function transform (FFT) exhibit well-defined regular stacking along the c axis. No other phase or polytype has been observed. In particular, no iron-or chromium-rich delafossite phase can be detected in the ED patterns and/or EDS measurements. Only occasionally has disorder in the stacking of successive cationic layers been observed in accordance with XRD pattern anisotropic peak broadening.

The nitrogen adsorption of the as-synthesized powders was measured by the BET method. The specific surface areas S w for all of the CuFe 1-x Cr x O 2 (0 e x e 1) samples are listed in Table 2. The specific surface areas are between 0.5 and 2 m 2 /g in accordance with powder synthesis at such high temperature. The larger the value of x, the more the specific surface area continuously increases according to Vegard's law, i.e., from 0.5 m 2 /g for x = 0 (CuFeO 2 ) to 2 m 2 /g for x = 1 (CuCrO 2 ). From the BET specific surface areas, the particle diameters d BET were calculated according to the equation d BET = 6/FS w (F is the crystallographic density determined from the lattice parameters). For this calculation, it is necessary to assume that the grains are spherical with a monodisperse distribution. d BET continuously decreases from 1.9 μm for x = 0 to 0.6 μm for x = 1, as indicated in Table 2. These particle size estimations are in agreement with the medium grain size d SEM observed by scanning electron microscopy (SEM) for all of the compositions (Table 2). SEM images for x = 0, 2 / 3 , and 1 (Figure 5) are in quite good agreement with the assumption about the spherical shape and the narrow distribution of the particle size distribution of the particles.

CuFe 1-x Cr x O 2 Thermostructural Behavior. Significant quantities of oxygen can be quickly and easily intercalated into the Cu 1+ M 3+ O 2 delafossite only when the M cations exceed a certain size, leading to off-stoichiometric Cu 1+ M 3+ O 2+δ . This is because O -anions have enough space to be intercalated between copper atoms in the copper plane. As the M cations become smaller, the Cu-Cu distances become too short for such oxygen passage. 11 Thus, oxygen intercalation into Cu 1+ M 3+ O 2 delafossite is forced or is even not possible if the M cations are too small.

The phase stability and thermal behavior of the CuFe 1-x Cr x O 2 solid solution was studied by TGA and high-temperature XRD under an air atmosphere up to 1000 °C in order to characterize the oxygen intercalation with a decrease in the M 3+ cation size from iron with r(Fe 3+ coord VI ) = 0.645 A ˚to chromium with r(Cr 3+ coord VI ) = 0.615 A ˚. TGA curves are plotted in Figure 6. For x = 0, i.e., for CuFeO 2 , the TGA experiment shows two weight gains. A slight oxidation of cuprous ions appears, leading oxygen intercalation from about 400 to 500 °C. An off-stoichiometric CuFeO 2+δ delafossite phase is then formed. 11 Above 500 °C, a phase transition occurs, leading to copper monoxide (CuO) and cuprospinel (Cu x Fe 3-x O 4 ) phases. Quantitative XRD analysis indicates that the copper content is equally distributed between the CuO and cuprospinel phases. It can then be deduced that the composition of the cuprospinel phase is x = 1 (CuFe 2 O 4 ) in accordance with the thermodynamic phase diagram. 13,14 XRD experiments at high temperature confirm this phase transition (Figure 7a). For M 3+ = Fe 3+ , oxygen intercalation is possible but is limited to low δ values. Although the radius of the ferric iron ions [r(Fe 3+ ) = 0.645 A ˚] is below the theoretical critical radius, allowing oxygen intercalation (r c = 0.70 A ˚)11 at room temperature, off-stoichiometric CuFeO 2+δ can be obtained at intermediate temperature because of the unit cell dilation. Figure 7b represents the linear variation of 0.003% per degree Celsius of the a parameter with temperature, as soon as it can be measured properly. For example, the a parameter is equal to 3.05 A ˚at T = 550 °C, which is 0.5% higher than the room temperature value. With a quite constant variation along the c axis (less than 0.1%), this leads to a volumic variation of 1.2% in the same temperature range. Even if additional space for oxygen intercalation becomes available at elevated temperature, off-stoichiometric delafossite intercalation is forced. This induces defects in the Cu-Cu plane first before leading to the delafossite-to-spinel phase transition. This is confirmed by the continuous decrease in the intensity and the heterogeneous enlargement of the Bragg peaks belonging to the delafossite phases during oxidation, followed by the apparition of the characteristic peaks belonging to the spinel and CuO phases. Quantification of the relative change of the main crystallographic phases is done at all temperatures by measuring the intensity of two isolated characteristic Bragg peaks [(006) at 31.5°in 2θ for delafossite and (220) at 30°in 2θ for spinel]. The values for the delafossite and spinel phases for x = 0 are reported in Figure 7c.

By analogy to a previous work on the structural analysis of the transformation mechanisms in iron oxides, 36 a detailed structural mechanism is proposed to explain this delafossite-to-spinel phase transition. This mechanism is schematically illustrated in Figure 8, and the main steps are labeled from 1 to 6. The delafossite structure (R3m) is a compact stacking of O -anions of the ABC type along the c axis. The first layer of the FeO 6 octahedra is the reference of the structural mechanism; it will then be assumed to be static in our model. The intercalation of extra oxygen atoms within the copper layers (step 1) induces a shearing of the [FeO 2 ] layers. The , respectively, with respect to the first layer (step 2). Half of the total copper atoms and their corresponding extra oxygen atoms get out of the structure and are simultaneously oxidized into CuO (step 3). A quarter of the total copper atoms stay in their environment, which is becoming an octahedral sublattice (step 4). The last quarter of the total copper atoms diffuse between the [FeO 2 ] layers to be stabilized in the tetrahedral sites (step 5). Because of the relative stabilities of the Fe 3+ and Cu 2+ species in the octahedral and tetrahedral sites of the spinel structure, 37 Cu 2+ ions diffuse into the FeO 6 octahedral layer, replacing the Fe 3+ cations, which themselves diffuse to the tetrahedral sites located between the two lower oxygen layers. Finally, Fe 3+ cations diffuse from the first FeO 6 octahedral layer to the tetrahedral sites into the two first oxygen layers (step 6). This is the same as the cationic diffusion into the second and third octahedral layers. For x = 1, i.e., CuCrO 2 , no weight gain and no phase transition have been observed up to 1000 °C in air. CuCrO 2 is thermally stable in air up to 1000 °C. 38 The Rietveld analysis of high-temperature XRD measurements shows that CuCrO 2 and CuFeO 2 exhibit the same linear variation of the a parameter of 0.003% per degree Celsius (see Figure 9). For CuCrO 2 , the cell parameter a is then never high enough to intercalate O -anions (a max = 3.00 A ˚), i.e., to allow cuprous ion oxidation.

For all of the intermediate compositions of CuFe 1-x Cr x O 2 where 0 < x < 1, complex redox reactions and phase transitions in both delafossite and spinel have been observed with TGA measurements (Figure 6). First, from about 500 to 700 °C, the TGA measurements show a weight gain for CuFe 1-x Cr x O 2 compounds corresponding to the oxidation of cuprous ions. Simultaneously, we observed by hightemperature XRD that the delafossite phase is transformed into spinel and CuO phases. Nevertheless, as the chromium content shifts to higher values, the oxidation and phasetransition temperatures increase. This can be explained by variation of the M 3+ cation size. The substitution of iron by chromium leads to a decrease in the average M 3+ radius, r (M 3+ ). The oxygen intercalation then becomes more difficult and requires higher temperatures to sufficiently dilate the delafossite structure to allow oxygen insertion.

The as-formed spinel and CuO phases are then reduced to delafossite above approximately 800 °C. TGA measurements show a weight loss of about 5% corresponding to the reduction of all Cu II ions to Cu I ions. This reduction is systematically correlated to a phase transition from spinel and monoxide to delafossite as a reverse option of the first phase transition. Figure 10 shows the thermodiffractograms of a sample with x = 2 / 3 (CuFe 0.333 Cr 0.667 O 2 ), which is subjected to such reversible redox reaction phenomena and related phase transformations.

In the case of M 3+ = Fe 3+ , the CuFe 2 O 4 spinel phase is only transformed into the CuFeO 2 phase after annealing at 1050 °C. From the thermodynamic data, one sees that the reduction of Cu 2+ into Cu + occurs above 1000 °C in air. 39 Indeed, the TGA curve of CuFeO 2 (Figure 6) shows a slight weight loss at 1000 °C. For 0 < x < 1 as well as for increased chromium quantities, the reduction temperature is lowered. This means that chromium facilitates the reduction of the spinel and CuO phases. The thermodynamic data of Navrotsky and Kleppa 37 and the structural description of the delafossite and spinel phases can explain this behavior: when x > 0, the delafossite-tospinel phase transition occurs mainly with the same structural mechanism as that described for x = 0. By contrast, however, Cr 3+ ions are more stable than Cu 2+ ions in an octahedral environment. In this case, the Cr 3+ and Cu 2+ cations do not diffuse into the pure octahedral and mixed tetrahedral/octahedral layers in the spinel structure, keeping the stacking sequence almost unchanged. The delafossite structure becomes more stable than the spinel structure as soon as chromium is substituted for iron. Thus, chromium ions tend to stabilize the stoichiometric delafossite phase, whereas ferric ions make cuprous ion oxidation possible and favor the delafossite-to-spinel phase transition.

Three successive thermal treatments of the sample with x = 2 / 3 (CuFe 0.333 Cr 0.667 O 2 ) under an air atmosphere at 1000 °C lead to three successive oxidations and reductions always associated with delafossite-to-spinel and spinelto-delafossite phase transitions, clearly showing the reversibility of the process. CuFe 0.2 Ga 0.8 O 2 delafossite was prepared in order to compare the effect of preferential coordination of M 3+ to the stability of the delafossite structure. In fact, this mixed gallium/iron delafossite is very interesting because the calculated mean radius of M 3+ ions in this phase is close to 0.625 A ˚, the value previously determined for CuFe 0.333- Cr 0.667 O 2 according to Shannon's table of ionic radii. 8 The stability difference between CuFe 0.2 Ga 0.8 O 2 and CuFe 0.333 Cr 0.667 O 2 is then only due to the coordination preference of gallium and chromium cations for octahedral or tetrahedral sites because no significant interference from the dimensional effect is present in this case.

CuFe 0.2 Ga 0.8 O 2 exhibits a single delafossite phase with lattice parameters similar to those of CuFe 0.333 Cr 0.667 O 2 , according to what was expected from the calculated mean M 3+ radii. Figure 11 shows that both delafossites are oxidized in the same temperature range. CuO and spinel oxide are precipitated for both gallium-and chromiumcontaining samples when the weight gain becomes higher than about 4.5%. At higher temperature (from about 900 to 1000 °C), however, the delafossite structure is formed again for the chromium-containing sample but not for the gallium-containing one. For the chromium-containing spinel, because of very strong stabilization of the chromium ions in the octahedral sites, 37 the copper ions are mainly located in the tetrahedral sites of the spinel structure. In other words, the mechanism previously described for CuFeO 2 in Figure 8 stops at step 5, i.e., before diffusion of the cationic species between the pure octahedral layers. For the gallium-containing spinel, because of stabilization of the gallium ions in the tetrahedral sites, the copper ions are mostly located in the pure octahedral layers. This corresponds to step 6 already described for the iron-containing samples. The reduction of cupric ions to cuprous ions is probably easier when they are located in tetrahedral sites because tetrahedral cuprous ions are more difficult to oxidize than octahedral cuprous ions. 40 Consequently, the cupric ion reduction would be easier for the chromium-containing sample. The migration of cuprous ions in 2-fold-coordinated sites would then be enough to induce the formation of the delafossite structure. For the gallium sample, more energetic activation would be required.

Conclusion

The complete CuFe 1-x Cr x O 2 (0 e x e 1) delafossite solid solution has been prepared by solid-state reaction at high temperature in an inert atmosphere. According to Vegard's law and a decrease of the ionic radius from r(Fe 3+ coord VI ) to r(Cr 3+ coord VI ), the unit cell parameters and cell volume decrease as x in CuFe 1-x Cr x O 2 increases. A thermostructural study has revealed that, with heating in air up to 1000 °C, the CuFe 1-

x Cr x O 2 (0 e x e 1) compounds of the solid solution are oxidized into spinel and CuO phases, except for CuCrO 2 , which is stable in this range of temperature. Chromium ions make it difficult for the oxidation and phase transition to proceed. Temperature-dependent XRD experiments have also shown that delafossite cell parameters vary linearly up to the phase transition into spinel and CuO phases. At high temperatures, the spinel and CuO phases can be reduced back to delafossite when the chromium quantity is enough. A structural mechanism has been proposed to explain this delafossiteto-spinel transition. For x = 0, the spinel structure is stable and the phase transition is not reversible. For x > 0, the substitution of Fe 3+ by Cr 3+ makes the formation of the delafossite phase easier. Preferential coordination of M 3+ to the tetrahedral site aids in the stabilization of the delafossite structure, as is confirmed by cationic substitution with gallium. It is noted that the Cu 2+ cations migrate in the delafossite structure very readily. 41 Half of the total copper ions in the delafossite gets out of the structure in order to oxidize into CuO and, subsequently, the copper ions in CuO reenter the delafossite structure during the spinel-to-delafossite transition.

Figure 2 .

 2 Figure 2. Unit cell volumes of the CuFe 1-x Cr x O 2 series with 0 e x e 1 after 10 h of treatment at 900 °C (in gray) and after 30 h at 900 and 1000 °C (in black).

Figure 1 .

 1 Figure 1. Schematic drawing of the 3R CuFeO 2 delafossite structure.

Figure 3 .

 3 Figure 3. Room temperature XRD pattern of CuFe 0.333 Cr 0.667 O 2 .

Figure 4 .

 4 Figure 4. [-110] zone axis HRTEM and corresponding FFT images of the CuFe 0.333 Cr 0.667 O 2 compound.

Figure 5 .

 5 Figure 5. SEM images of the CuFe 1-x Cr x O 2 solid solution when x = (a) 0, (b) 2 / 3 , and (c) 1.

Figure 6 .

 6 Figure 6. TGA measurements in air for the CuFe 1-x Cr x O 2 solid solution series with 0 e x e 1.

Figure 7 .

 7 Figure 7. (a) Change in the XRD patterns of CuFeO 2 with respect to temperature (b). Corresponding change in the a lattice parameter with respect to the temperature. (c) Relative change in the phase intensities of delafossite (circle) and spinel (square).

Figure 8 .

 8 Figure 8. Schematic representation of the structural mechanism for the delafossite-to-spinel phase transition.

Figure 9 .

 9 Figure 9. (a) Change in the XRD patterns of CuCrO 2 with respect to temperature. (b) Corresponding change in the a lattice parameter with respect to the temperature. (c) Relative change in the phase intensities of delafossite (circle) and spinel (square).

Figure 10 .

 10 Figure 10. (a) Change in the XRD patterns of CuFe 0.333 Cr 0.667 O 2 with respect to temperature. (b) Corresponding change in the a lattice parameter with respect to the temperature. (c) Relative change in the phase intensities of delafossite (circle) and spinel (square).

Figure 11 .

 11 Figure 11. Comparative TGA measurements in air for CuFe 0.2 Ga 0.8 O 2 and CuFe 0.333 Cr 0.667 O 2 .

  Lattice Parameters for a CuFe 1-x Cr x O 2 Solid Solution Series with 0 e x e 1

		a (A ˚)	c (A ˚)	V (A ˚3)
	CuFeO 2	3.0344(2)	17.158(3)	136.82(2)
	CuFe 0.835 Cr 0.165 O 2	3.0268(2)	17.154(3)	136.10(2)
	CuFe 0.667 Cr 0.333 O 2	3.0159(2)	17.141(2)	135.02(2)
	CuFe 0.5 Cr 0.5 O 2	3.0055(2)	17.128(2)	133.99(2)
	CuFe 0.333 Cr 0.667 O 2	2.9942(2)	17.109(2)	132.84(3)
	CuFe 0.165 Cr 0.835 O 2	2.9826(2)	17.097(2)	131.71(3)
	CuCrO 2	2.9742(2)	17.090(2)	130.93(2)

1.

Table 2 .

 2 Specific Surface Areas (S w ) Calculated from BET Measurements and Estimated SEM Particle Sizes for a CuFe 1-x Cr x O 2 Solid Solution Series with 0 e x e 1

		S w (m 2 /g)	d BET (μm)	d SEM (μm)
	CuFeO 2	0.57(2)	1.92(7)	1.6(4)
	CuFe 0.835 Cr 0.165 O 2	1.01(3)	1.08(3)	1.0(3)
	CuFe 0.667 Cr 0.333 O 2	1.36(4)	0.80(2)	0.8(2)
	CuFe 0.5 Cr 0.5 O 2	1.41(4)	0.77(2)	0.7(3)
	CuFe 0.333 Cr 0.667 O 2	1.56(5)	0.69(2)	0.6(2)
	CuFe 0.165 Cr 0.835 O 2	1.75(5)	0.61(2)	0.6(3)
	CuCrO 2	1.94(6)	0.55(2)	0.5(2)


Antoine Lonjon

Lydia Laffont

Philippe Demont
email: demont@cict.fr.

Eric Dantras

Colette Lacabanne

New Highly Conductive Nickel Nanowire-Filled P(VDF-TrFE) Copolymer Nanocomposites: Elaboration and Structural Study

New highly electrical conductive nanocomposites were prepared by dispersing nickel nanowires into a poly(vinylidene difluoride)-trifluoroethylene P(VDF-TrFE) matrix. A suspension of individual nickel nanowires with a regular high aspect ratio (ξ ≈ 250) was elaborated. The nickel nanowires were fabricated by electrodeposition using templates in anodic aluminum oxide with a nominal pore diameter of 200 nm, allowing a close control of nanowire crystallinity. Polycrystalline or single crystal nickel nanowires were obtained. An oxide layer was observed on nanowire surfaces after their extraction from the template. Physical and chemical treatments were used to completely remove the oxide layer. Scanning and high resolution transmission electron microscopy studies were performed. The elemental composition and the nature of the nanowires surface were investigated by electron diffraction and energy dispersive spectroscopy. Nickel nanowires without oxide layers were elaborated. The electrical conductivity of nanocomposite films was performed as a function of treated nickel nanowire volume fraction. A very low percolation threshold of 0.75 vol % was determined. Percolated nanocomposites filled by treated nanowires displayed a highly electrical conductivity value. The conductivity value obtained above the percolation threshold is the highest value known up to now in the case of a conductive nanoparticle dispersion.

Introduction

There is great interest for one-dimensional nanoparticles with high aspect ratio ξ (ξ ≈ 100-1000) in nanocomposites elaboration. Nanomaterials such as carbon nanotubes (CNT), 1 nanorods, 2 and nanowires 3 enable the formation of highperformance nanocomposites [4][5][6] at a much lower concentration previously shown with standard microsized fillers. Metallic nanowires have attracted much attention because of their specific properties as mechanical resistance, 7,8 stiffness, 9 electrical conductivity, 10 and magnetic susceptibility. 11 These specific properties show that nanowires are excellent candidates to substitute the standard nanofillers in the formulation of multifunctional polymer nanocomposites. 12 Several methods exist to synthesize metallic nanowires. One of them consists of growing a metallic germ in a solution by reducing metal ions. A growth direction is induced by using a capping agent. Silver nanowires have been obtained by this chemical method. 13 The other one uses the template technique as a support. Two ways are described in the literature: covering a high aspect ratio object by a metallic layer (positive template) 14 or filling a high aspect ratio support (negative template) 15 such as an anodic aluminum oxide (AAO) porous membrane. 16 This technique has been used for several years to produce metallic nanowires. The nanostructure shape depends on the template used. Template electrodeposition studies 17 have been focused on the maximal ratio obtained with the smallest nanometer pores.

One of the aims of this work is to synthesize metallic nanowires by the negative template technique without defects and so elaborate polymer nanocomposites having metallic behavior for electrical conductivity. The nanofiller volume fraction must be as low as possible to maintain the mechanical properties of the pristine polymer.

A great interest has been devoted to the chemical nature of the nanowire surface. In spite of their high specific chemical reactivity, the metallic behavior of the surface has to be promoted.

Experimental Methods

Nanowire Synthesis. Nickel nanowires (Ni NWs) were synthesized by electrochemical deposition in an anodic aluminum oxide (AAO) porous template. The electrolyte used for nickel deposition was a standard Watts bath with an accurate pH value of 4. Direct current (DC) electrodeposition was carried out at room temperature using a nickel wire as anode with 1.0 mm of diameter. A porous (AAO) membrane of 200 nm diameter and 50 μm thickness was supplied by Whatman. One side of the AAO membrane (silver back side) was coated with a 35 nm thickness silver layer by using a sputtering technique as cathode for electrodeposition. Nickel NW growth was controlled by the deposition time and the direct current intensity using a Keithley 2420 source meter. The AAO membrane was dissolved in NaOH 6 M for 30 min, releasing Ni NWs from the template. Then, Ni NWs were washed with distilled water and filtrated. A following sonication led to a very good dispersion of NWs in water or toluene.

An oxide layer was observed on Ni NWs which is responsible for the decrease of NW conductivity. Two methods were used to remove oxide layer. The first one was annealing under reducing atmosphere at 200 °C under H 2 atmosphere for 30 min. The second was dissolving of the oxide layer in a 0.25 M solution of H 2 SO 4 at 60 °C and stirring at 400 rpm for 5 min. The treated nanowires were filtered, rinsed with water, and stored in toluene or acetone. The same process was performed with H 3 PO 4 acid as an alternative treatment.

Nanocomposite Preparation. Poly(vinylidene difluoride)trifluoroethylene P(VDF-TrFE) 70/30 mol % copolymer was supplied by Piezotech S.A. (Saint Louis, France). The melting temperature was 150 °C determined by using a differential scanning calorimeter (Perkin-Elmer DSC7), and the density was about 1.8 g cm -3 . The nanocomposites were prepared by using a solvent casting method. P(VDF-TrFE) copolymer was dissolved in acetone. Treated Ni NWs were then poured into the polymer solution, and the mixture was submitted to a short pulse of sonication for 5 s, corresponding to a dissipated power of 25 W, to obtain a suitable dispersion of the nanowires in the polymer solution. A slight curvature of nanowires can be observed when a strong sonication is applied because of the metal nanowires' ductility and aspect ratio. The mild sonication used in this study allows nanowires to release from each other and their dispersion in the matrix, and does not affect the aspect ratio of nickel nanowires as seen in the different SEM images.

The solvent was evaporated using a magnetic stirrer at 80 °C for 1 h. Pellets of unaligned nickel nanowires randomly dispersed in P(VDF-TrFE) matrix were obtained. Films with a thickness of 200 μm were obtained using compression molding in a hot press at 200 °C under a pressure of about 0.3 MPa. No traces of solvent were observed during a thermogravimetric analysis (TGA) scan, indicating a complete removal of acetone.

Electron Microscopy. The morphology of the nickel nanowires was examined by scanning electron microscopy (SEM) using a JEOL JSM 6700F instrument equipped with a field emission gun (SEM-FEG). Nanocomposite samples were fractured at the liquid nitrogen temperature for observation by SEM. The structure and texture of the Ni NWs were determined by high-resolution transmission electron microscopy (HRTEM). A drop of the different NW suspensions was deposited on a copper grid coated with lacey carbon films to obtain transparent samples for electron irradiation. TEM and HRTEM images were obtained by using a FEI Tecnai F20 S-TWIN microscope. The elemental composition was also determined by energy dispersive spectroscopy (EDS) at a nanometre scale. The diffraction patterns were obtained using the selected area electron diffraction (SAED) mode.

Electrical Conductivity. Electrical conductivity measurements were carried out by recording the complex conductivity σ*(ω) using a Novocontrol broadband spectrometer. The measurements were done in the frequency range from 10 -2 to 10 6 Hz at room temperature. The real part, σ′(ω), of the complex conductivity σ*(ω) was investigated. For all the nanocomposite samples considered in this study, the phase lag between the measured impedance and the applied ac voltage was negligible at low frequencies, so that the reported impedance at 0.01 Hz is equivalent to the dc resistance. The dc conductivity σ dc of samples was determined from the independent frequency part of σ′(ω) (low frequency plateau). Films 200 μm thick were placed between two circular gold platted electrodes (20 mm in diameter). To reduce contact resistance with the cell electrodes, a thin layer of gold (100 nm) was sputtered onto both sides of the films using a BOC Edwards scancoat six SEM sputter coater.

Electrical Conductivity Percolation. The electrical performance of nanowire-filled polymer nanocomposites is strongly dependent on the volume fraction p and the aspect ratio (length to radius) L/r of the filler. In disordered systems, the insulator to conductor transition is driven by the percolation law where a critical volume fraction or percolation threshold of the filler is necessary for the onset of electrical conduction. According to the percolation theory, the dc conductivity σ dc increases by a power law above the percolation threshold p c : where σ 0 is a constant and t is the critical exponent. [START_REF] Stauffer | Introduction to Percolation Theory[END_REF] For percolation in a lattice, t depends on the lattice dimension d: t ) 1.1-1.3 for d ) 2 and t ) 1.6-2 for d ) 3. The percolation threshold model in a three-dimensional sticks system, proposed by Balberg and Binenbaum, [START_REF] Balberg | [END_REF] shows that the percolation threshold is dependent on the aspect ratio and the anisotropy of sticks through their excluded volume V ex . For that, the nanowires are capped cylinders with L/r .1 and the associated excluded volume is defined as where θ is the angle between two NWs and 〈sin θ〉 describes the degree of alignment of the NWs. For randomly oriented NWs, 〈sin θ〉 is equal to π/4 and eq 2 is written as where V NW is the NW volume. At the percolation threshold, V ex is found to be constant. When NWs are randomly oriented, the critical excluded volume V ex cr is between 1.4 and 1.8 and is given by where N c and f c are the NWs critical concentration and critical volume fraction for percolation, respectively.

Results and Discussion

Much interest has been devoted to the chemical nature of the nanowire surfaces. In spite of their high specific chemical reactivity, the metallic behavior of the surface has to be promoted.

The surface morphology of the AAO membrane was observed by SEM. Figure 1 shows an average pore diameter of about 200 nm with a very narrow distribution, and a pore density of about 5 × 10 8 pores cm -2 was estimated. The filling of the AAO membrane pores with nickel is performed by using direct current intensity ranging from 0 to 105 mA under a voltage of 3 V. The current density controls the deposition time. Ni NW growth (1)

V ex ) 32 3 πr 2 + 8πLr 2 + 4L 2 r〈sin θ〉 (2) V ex ) πL 2 r ) L r V NW (3) V ex cr ) L r V NW N c ) L r f c ( 4
)
is dependent on the electron transfer which is regulated by the speed of the electron through the interface metal/solution and the kinetic transport of metallic ions to the electrode. Figure 2 shows the dependence of the time of deposition on intensity current. An extrapolation of the nanowire length L is also possible. A linear relation between the deposition time and the nanowire length is characteristic of behavior of charge transfer. The nickel nanowire morphology was studied by SEM to determine the aspect ratio ξ. Figure 3 shows that Ni NWs are still attached to the silver coated layer backside of the alumina membrane, forming a "carpet fakir". NWs have a uniform length of 50 μm and a regular diameter of 200 nm. The silver coated layer backside and the AAO membrane were dissolved in HNO 3 and NaOH solution, respectively. Then the nanowire suspension was filtrated on 200 nm diameter polyamide membrane filter, and 12 mg of nickel nanowires per cm -2 was obtained. Nickel nanowires are well released from the membrane. Figure 4 shows the nickel nanowires releasing from the membrane. The density of Ni NWs is very high due to their suitable straightness. SEM confirms that Ni nanowire surfaces are smooth and continuous.

The conductivity of NWs is limited by two factors: the bulk resistivity and their contact resistance. The bulk resistivity is negligible compared to the contact one. An oxide layer formation occurs when the nanowires are released from the membrane and silver coated layer back side. The presence of the oxide layer is confirmed by conductivity measurements on pressed Ni NW samples for which a value of only 20 S m -1 is reached, a value drastically lower than that of nickel bulk: 1.56 × 10 7 S m -1 . The conductivity of Ni NWs treated with H 2 SO 4 and H 3 PO 4 acids is performed after rinsing with water then ethanol and dried at 50 °C for 30 min. NWs reduced with H 2 SO 4 display a higher conductivity (115 S m -1 ) than those reduced with H 3 PO 4 (24 S m -1 ).

Filtered nickel nanowires have been stored in acetone and dispersed using an ultrasonic bath. A droplet of this solution has been dried and deposited on an SEM pin. A good dispersion of the suspended nanowires without bundles is observed. The SEM image clearly shows that nickel nanowires are not damaged by sonication.

TEM images of a single nickel wire are shown in Figure 5a andc with their corresponding SAED patterns in Figure 5b andd. The average diameter is about 300 nm ((a) 330 nm; (c) 294 nm). The SAED pattern in Figure 5b shows single bright spots characteristic of nickel single crystals, while Figure 5d reveals polycrystallinity especially for the [111] direction, that is, the growth direction that is in good agreement with the literature. 20 Nevertheless, TEM characterizations of nickel nanowires have been performed in the literature, 17,20,21 but no results about the nature of the layer on the nanowire surfaces have been reported from the bright field TEM images. So the chemical nature and structure of the nanowires surface need to be well characterized.

The nickel nanowires, released after the silver backing and the alumina membrane dissolution, were examined by HRTEM. Figure 6 shows the bright field image of a nickel nanowire of 350 nm in diameter. One can see on the nanowire surface a discontinuous layer of varying thickness. Inset 1 in Figure 6 shows that the SAED pattern, taken from the end of the single nickel nanowire, confirms the nanowire to be Ni. The diffraction spots are well indexed to the face centered cubic (fcc) structure (JCPDS 65-2865). Inset 2 reveals both Ni and NiO phases. The NiO indexation has been done thanks to the JCPDS File (NiO: 75-0269; Fm3 j m). The thin layer, visible only in the portion of the nanowire corresponding to SAED 2, confirms the surface of the NW to be oxidized. It is well-known that nickel oxidation can be limited by physical and chemical treatment.

The nickel nanowires, released after the silver backing and the membrane dissolution, were dispersed into water. Figure 7a shows a HRTEM image of a single nanowire of 270 nm in diameter with some filaments on the surface. One part of the bright field image was acquired via high resolution, showing that the NW is covered by a Ni(II) hydroxylated oxide Ni(OH) 2 (JCPDS 03-0177, hexagonal) layer. To remove this oxide layer, the Ni nanowires are treated with sulfuric acid. The HRTEM image at low magnification (Figure 7b) shows that the surface of the nanowire is homogeneous and no filaments or layer are visible at this magnification. A high resolution image of one part of the NW is characteristic of a well-crystallized wire. The two-dimensional Fourier transform calculated from high resolution images of the inner and the surface of the nanowire shows diffraction spots representative of Ni and NiO phases. The passivation of nickel NW is topotactic because both phases have the [111] direction in common. A H 2 reduction has no effect on the nanowire surface as seen in Figure 7c (diameter of 241 nm). Nevertheless, the spacings in the SAED pattern are in good agreement with spacings of the fcc NiH (JCPDS 65-0099). H 2 reduction has modified the chemical nature of the nanowires. Nickel nanowires have shown an ability to absorb hydrogen that confirms the Ni NWs as potential candidates for hydrogen storage.

The effect of phosphoric acid on NWs has also been investigated by TEM (Figure 7d). The aspect ratio of the NWs is not modified, and no oxide layer has been observed at high magnification. The corresponding SAED pattern of the bright field image is indexed by the fcc structure of the Ni phase, and the growth direction of the pure nickel nanowire is [200]. Sulfuric acid prevents massive oxidation and confers to nickel nanowires higher intrinsic electrical conductivity. Table 1 summarizes the chemical nature of the surface nanowire, observed by TEM, as a function of the chemical or physical treatment applied to nanowires.

Ni NW/P(VDF-TrFE) composites were prepared with a volume fraction of Ni NWs varying from 0 to 5 vol %. Figure 8 shows the SEM image of the polymer matrix filled by 1, 2, 3, and 4 vol % of Ni NWs. Nickel nanowires are clearly individual in acetone, and this homogeneous dispersion is conserved in P(VDF-TrFE) copolymer. The temperature and pressure conditions applied for the compression molding nanocomposite film processing involve a slight orientation of the nanowires.

The electrical conductivity of the insulating P(VDF-TrFE) copolymer is 10 -12 S m -1 at room temperature. The dc electrical conductivity of the Ni NW/P(VDF-TrFE) samples at room temperature is plotted as a function of the NW volume fraction (Figure 9). The electrical conductivity increases drastically by 13 orders of magnitude at a very low percolation threshold of 1 vol % in comparison with the high value of 20 vol % observed for Ni spherical particles. The data are best-fitted by a scaling law according to eq 1 with p c ) 0.75 vol %, σ 0 ) 3.55 × 10 -3 S m -1 , and t ) 1.57 ( 0.16. The critical exponent value t is smaller than the universal value for three-dimensional percolation systems which is equal to 1.94. From the scaling law, σ 0 is an extrapolation at 100 vol % of NWs, which corresponds to the conductivity of Ni NW bundles. σ 0 is 5 orders of magnitude lower than the conductivity previously measured in Ni NWs. The Ni NWs display a narrow distribution of their aspect ratio around ξ ≈ 250. Assuming the stick is Ni nanowire, we find a critical volume fraction of 0.64 ( 0.08 vol % according to eq 4, which is slightly lower than the experimental value of 0.75 vol %. This weak difference can be explained by the formation of a few poorly dispersed NW bundles. It is wellknown that the percolation is dependent on the ability to disperse individually and uniformly the NWs in the polymer matrix. Such a low filler volume fraction is sufficient to establish a threedimensional conducting network because of the high aspect ratio of Ni NWs. The measured conductivity arises from the electrical properties of nickel and the surface treatment of Ni NWs. A conductivity of 10 2 S m -1 is reached above p c even if a value of only 10 -1 S m -1 is measured on powder of spherical micrometer-sized nickel particles. 22 As a comparison, a conductivity value of about 1 S m -1 is obtained in Ni NWs without surface treatment.

Conclusion

We have successfully developed an easy method to produce a great quantity of nickel nanowires with controlled geometry, crystallinity, dispersion, and surface layer without oxide. Nickel nanowires with a high aspect ratio (∼250) were electrodeposited in an AAO template membrane. Elaboration of individual and uniformly dispersed nickel nanowires was reported. This electrodeposition method allows keeping the geometry of nanowires after the process and the dispersion in solvent. A quantity of 12 mg cm -2 of nickel nanowires was obtained in one time, suggesting an easy way to produce several grams. The electrodeposition parameters need to be adjusted in order to control the crystallinity of the Ni nanowires. The nanocomposite prepared by dispersing Ni NWs in P(VDF-TrFE) copolymer melt shows a very low percolation threshold: 0.75 vol %, preventing a damage of flexibility, weak density, and light weight of the copolymer matrix. An electrical conductivity of 10 2 S m -1 at such low filler content is the highest value to our knowledge. These results underline the key points and the essential parameters to elaborate highly conductive nanocomposites. The aspect ratio and the dispersion of nanofillers govern the percolation threshold, while the electrical conductivity of final conductive Ni NW/P(VDF-TrFE nanocomposites is supplied by the intrinsic conductivity of nickel nanowires. The composite conductivity value obtained slightly above the percolation threshold is the highest measured with high aspect ratio conductive nanomaterials. 23

Figure 1 .

 1 Figure 1. SEM images of the AAO membrane at two different magnifications before nickel electrodeposition.

Figure 2 .

 2 Figure 2. Nickel nanowire length as a function of the intensity of applied current for three electrodeposition times.

Figure 3 .

 3 Figure 3. SEM image, at two magnifications, of nickel nanowires after complete removal of the AAO membrane but still anchored to the silver back side.

Figure 4 .

 4 Figure 4. SEM image of nickel nanowires after complete removal of AAO membrane and silver back side and dispersion in acetone.

Figure 5 .

 5 Figure 5. (a,c) Bright field TEM images of an individual nickel nanowire. SAED patterns (b) and (d) reveal, respectively, the single crystalline and polycrystalline nature of nickel.

Figure 6 .

 6 Figure 6. HRTEM image of the surface of an individual nickel nanowire and the SAED pattern taken from different portions of the individual wire.

Figure 7 .

 7 Figure 7. TEM and HRTEM images of individual nickel nanowire obtained after chemical treatments: (a) water, (b) sulfuric acid, (c) H 2 reduction, and (d) phosphoric acid. Insets in (a) an (b) show the twodimensional Fourier transforms of the selected area at the surface of the nanowire. Insets in (c) and (d) show the corresponding SAED patterns.

Figure 8 .

 8 Figure 8. SEM images of the fracture surface of nickel nanowire/ P(VDF-TrFE) copolymer composites at different Ni NW volume fractions: (a) 1 vol %, (b) 2 vol %, (c) 3 vol %, and (d) 4 vol %.

Figure 9 .

 9 Figure 9. Dependence of the dc conductivity (σ dc ) on the nickel nanowire volume fraction in P(VDF-TrFE) matrix at 25 °C. The solid line connecting symbols is a guide for the eyes. The inset shows the log-log plot of σ dc versus (pp c ) with p c ) 0.75% and t ) 1.57. The solid line corresponds to the best-fitted line.

TABLE 1 : Chemical Nature of the Surface Nickel Nanowire as a Function of the Treatment

 1

	treatment	dispersion in H 2 O	dispersion in C 7 H 8	H 2 reduction	H 2 SO 4 reduction	H 3 PO 4 reduction
	chemical nature of surface wire	Ni(OH) 2	NiO	NiH	NiO	Ni

Acknowledgment. The work was supported by Direction Ge ´ne ´rale des Entreprises (DGE) and Conseil Re ´gion Midi-Pyre ´ne ´es (CRMP) through NACOMAT contract.


Seng Kian Cheah

Emilie Perre

Mårten Rooth

Mattis Fondell

Anders Hårsta

Leif Nyholm

Mats Boman

Torbjo ¨rn Gustafsson

Jun Lu

Patrice Simon

Kristina Edstro ¨m

Self-Supported Three-Dimensional Nanoelectrodes for Microbattery Applications

A nanostructured three-dimensional (3D) microbattery has been produced and cycled in a Li-ion battery. It consists of a current collector of aluminum nanorods, a uniform layer of 17 nm TiO 2 covering the nanorods made using ALD, an electrolyte and metallic lithium counter electrode. The battery is electrochemically cycled more than 50 times. The increase in total capacity is 10 times when using a 3D architechture compared to a 2D system for the same footprint area.

Li-ion batteries are today used for a wide range of applications such as portable electronics, medical devices, and in prototypes for hybrid electric vehicles (HEV). The chemistry in the battery can be tuned depending on the performance requirement for the different applications and hence the chemical content of an HEV battery is not in detail the same as for a mobile phone battery. This tuning allows for a miniaturization to make small scale power supplies and energy storage devices. However, to make a high energy and power unit on a small footprint area (for example <1 cm 2 ) requires the use of nanostructured architectures, threedimensional (3D) microbatteries. Three-dimensional microbatteries can be used in powering, for instance, MEMS-based sensors and actuators, 1 biosensors, and other similar applications. There are still only a few examples in literature of 3D microbatteries. [2][3][4] In this work, we show for the first time how a 3D Li-ion battery can be made based on a nanostructured current collector of aluminum in the form of aluminum nanorods, and that it is possible to uniformly deposit a cathode material (in this case TiO 2 ) on the rods despite the narrow distance between them (∼20 nm for homemade membranes and a minimum 50 nm for commercial membranes) and then to cycle it in a Li-ion battery. We selected TiO 2 based on its simplicity to synthesize thin films with the anatase structure with a broad selection of different techniques. Anatase can host up to 0.5 Li per formula unit and with TiO 2 particles <40 nm up to 0.8 Li can be stored. 5 TiO 2 is also abundant, nontoxic, and cheap.

We have earlier shown that we can grow aluminum nanorods using a direct electrodeposition procedure into a nanoporous alumina membrane used as a template. 6 The rods can be synthesized in both commercial nanoporous membranes (Anodisc 47, Whatman) (Figure 1a,b) and in membranes synthesized in-house through either mild-or hard anodization as described in literature. 7,8 Our results presented here are based on commercial membranes. All the steps for the preparation of aluminum nanorods were carried out in an argon-filled glovebox (O 2 and H 2 O <2 ppm) since the ionic liquid, 1-ethyl-3-methylimidazolium chloride ([EMIm]-Cl)/aluminum chloride (AlCl 3 ) (1:2 ratio), used as the deposition electrolyte can easily decompose in contact with water.

The acidic electrolyte was prepared by slow addition of aluminum chloride (anhydrous powder, 99.99%, Sigma-Aldrich) into 1-ethyl-3-methylimidazolium chloride (98.5%, Fluka) under continuous stirring. Since the reaction of AlCl 3 with [EMIm]Cl is highly exothermic, a light white smoke could be observed if the addition of the AlCl 3 was too fast. Therefore, care must be taken to prevent the decomposition of the electrolyte. A brownish liquid was formed at the end of the mixing.

To prepare the aluminum nanorods, a three-electrode setup connected to a potentiostat/galvanostat Autolab PGSTAT30 was used for the voltage pulse electrodeposition. 6 Two 1 cm × 1 cm aluminum plates were used as the working and counter electrodes. A thick glass fiber separator was filled with the electrolyte. The reference electrode was an aluminum wire with 1 mm diameter (99.999%, Goodfellow) immersed in the electrolyte ([[EMIm]Cl/AlCl 3 , molar ratio 1:2). The free-standing aluminum nanorods were then obtained by dissolving the alumina template outside the glovebox in an aqueous solution of CrO 3 (1.8 wt %) and H 3 PO 4 (6 wt %) 9 followed by rinsing in water.

The aluminum nanorods, however, do not grow as evenly as the Cu nanorods reported by Taberna et al. 10 There is instead an "island-like" growth resulting in an uneven heightdistribution. This can depend on nucleation problems on the flat substrate or of the nonhomogenous distribution of pores in the commercial nanoporous aluminum membrane. The best nanorods, however, were obtained when applying a pulse potential deposition sequence, -0.4 V versus Al for 50 ms then -0.1 V versus Al for 200 ms.

The electrode material TiO 2 was then deposited on to the aluminum nanorods by the use of atomic layer deposition (ALD). ALD is a method that can meet the high demands of controlled thickness and high uniformity on nonplanar substrates. 11 ALD is therefore an ideal method for depositing TiO 2 on aluminum nanorods and in the present study TiO 2 was deposited from TiI 4 and H 2 O (see, e.g., ref 12). An Al substrate covered with aluminum nanorods was placed on a substrate holder in the ALD reactor. The evaporation zone and the deposition zone were preheated to 100 and 200 °C, respectively, for half an hour. The deposition temperature was 200 °C and the TiI 4 was evaporated at 108 °C. These temperatures were maintained through out the rest of the experiment.

The ALD process contained four different pulses in which each lasted for 15 s. In the first pulse, TiI 4 was evaporated and carried to the deposition zone by inert nitrogen gas where it reacted with the surface of the substrate. As TiI 4 adsorbed on the surface, one or more of the ligands were broken off and entered the gas phase in form of I or I 2 depending on the temperature. The second pulse was a purging pulse where the reaction chamber was flushed with N 2 to remove reaction biproducts (I, I 2 ) and the excess of TiI 4 . The third pulse was the water pulse where water vapor reacted with the adsorbed titanium iodide forming TiO 2 and HI. The fourth and last pulse was another purging pulse in which the HI and excess water vapor was removed by a N 2 pulse. These four pulses formed a cycle. One hundred and fifty cycles were used for the deposition of TiO 2 .

X-ray diffraction characterization was then carried out on the deposited film using a Siemens D5000 grazing incidence diffractometer and Cu KR radiation at 1°grazing angle. The TiO 2 film showed to have an anatase structure.

The Al nanorods with the deposited anatase were further examined with HRTEM (see Figure 2). The HRTEM (Tecnai F30 ST, FEI) was used to determine the thickness and crystalline phases of the film. It was shown that a 17 nm thick layer of TiO 2 was deposited onto the rods. In Figure 2a, the HRTEM micrograph shows that the whole surface of the Al nanorods is covered by a conformal layer of TiO 2 deposited by ALD; even the edges are covered with TiO 2 . In Figure 2b, the narrow spacing between nanorods, around 40 nm, is also covered by the conformal layer of TiO 2 . In Figure 2c, a thin layer of Al 2 O 3 , approximately 4 nm, can be observed in between the TiO 2 layer and the Al nanorods. The oxide layer forms due to the exposure to air before the deposition of TiO 2 . This oxide layer not only provides extra protection to the Al nanorods from reactions with the electrolyte in a battery but also facilitates the growth and adhesion of the TiO 2 layer on the rods. The HRTEM study also shows the structure to be more complicated than seen from the X-ray diffraction experiment only. The combination of the techniques shows the deposited TiO 2 film to consist of amorphous matrix with nanocrystalline anatase.

The electrode was then tested in a Li-ion battery context. Galvanostatic cycling tests were performed using the TiO 2 on the aluminum nanorods as the working electrode, lithium metal as a counter electrode, and a glass fiber separator soaked with an electrolyte of 1 M LiClO 4 -PC (propylene carbonate) (the salt was dried at 80 °C overnight and the PC from Merck was used as received). The water content is <10 ppm (Karl Fischer titration). The electrochemical cycling tests were carried out with a Digatron battery testing system between 3.0 and 1.4 V. Cells were assembled in a glovebox under argon atmosphere (<2 ppm H 2 O and O 2 ) and packed into a polymer-coated aluminum pouch.

The galvanostatic charge-discharge curve of TiO 2 deposited on the nanostructured current collector is presented in Figure 3. The curve shows the typical behavior of a semicrystalline sample. The plateaus observed at 1.77 V during discharge and at 1.9 V when charging, correspond to the insertion/extraction of lithium within the crystalline part of TiO 2 .

To be able to compare the cycling results for the 2D and 3D architectures, the amount of deposited TiO 2 needs to be determined. Considering for both 2D and 3D substrates, a TiO 2 layer of 17 nm and a density of 3.895 g•cm -3 , the mass of deposited TiO 2 can be calculated. For a 1 cm 2 × 1 cm 2 flat Al substrate, the mass deposited is equal to 0.0066 mg, while for an Al substrate containing nanorods onto a footprint area of 1 cm 2 × 1 cm 2 the mass deposited is equal to 0.0667 mg. The amount of TiO 2 deposited onto the 3D substrate has been evaluated based on the calculation of area gain reported by Perre et al. 6 Assuming a constant nanorod diameter of 200 nm, an average height of 2 μm, and an interpore distance of 400 nm, the area gain is equal to 10.07.

The theoretical capacity of TiO 2 is 167.7 mAh•g -1 for 0.5 inserted Li-ion, thus corresponding to theoretical capacity values of 0.0011 mAh•cm -2 for the 2D microbattery and 0.0112 mAh•cm -2 for the 3D microbattery.

As expected the results of the battery cycling show an increase in total capacity when using a 3D architectured electrode compared to a 2D electrode (see Figure 5) due to the higher amount of TiO 2 deposited onto the 3D nanostructured current collector. This increase in capacity by a factor of 10 is consistent with the calculated area gained on a 3D substrate compared to a 2D substrate.

For both the 3D and the 2D electrodes, the measured capacity values are close to the theoretical values and are maintenained upon cycling. Thus showing that due to the small thickness deposited, TiO 2 accommodates easily the structural changes taking place during the insertion of 0.5 Li ion. This result is consistent with the reported result by Sudant et al. 13 The good correlation between the calculated and measured capacity values also confirms the efficiency of the ALD process when a conformal layer needs to be deposited onto a complex 3D substrate.

The rate capability of the nanostructured TiO 2 electrode was also tested and shown to be excellent. This is shown in Figure 5, where TiO 2 was cycled at different rates and still retained its initial capacity values after 50 cycles. The high irreversible capacity observed in the first cycle can be attributed to the presence of defects within the TiO 2 amorphous structure that trap the Li + ions. 4 All these results show the 3D nanostructured electrode to be a promising system. The electrodes are able to provide approximately 40 and 35% of the initial capacity (cycling at C/5 rate) even though the cycling current has been increased by 50 (10 C) and 100 (20 C) times, respectively. The outstanding performance is the result of the nanostructured active material and the conformal deposition. Despite of the small grain size and short diffusion length, the electrode still shows the typical intercalation behavior for a crystalline TiO 2 . In a previous report, 5 the biphasic property of the TiO 2 during charge and discharge started to disappear when the size of the TiO 2 is smaller than 120 nm.

The results revealed in this study are promising for the future development of 3D nanostructured architectures for lithium-based microbatteries. The arrays of aluminum current collector coated by TiO 2 showed higher volumetric capacity as compared to the results reported by Min et al. 2 and higher rate capability compared to the design proposed by Golodnistky et al. 3 Ortiz et al. 4 recently reported the cycling data of dense arrays of TiO 2 anatase nanotubes grown by anodization of a Ti substrate. They presented high capacity values per footprint area. However, in such architectures power performances are driven by the height of the active material pillars; 1 accordingly power capability was limited to 2C.

Our study opens new paths for the study of deposition processes of other cathode materials than TiO 2 on complicated nanostructured current collectors.

Figure 1 .

 1 Figure 1. SEM image of aluminum nanorods synthesized by -0.4 V vs Al pulses for 50 ms and -0.1 V vs Al pulses for 200 ms 14 400 times in a commercial membrane top view (left) and side/tilted view (right).

Figure 2 .

 2 Figure 2. HRTEM images of Al nanorods with deposited TiO 2 at 200 °C for 150 cycles. A homogeneous TiO 2 layer (dark color) is covering the whole surface of the nanorod tip (a). The TiO 2 layer is also covering the narrow spacing in between the nanorods (b). There is a 4 nm thick Al 2 O 3 layer between the TiO 2 layer and the nanorod (c).

Figure 3 .

 3 Figure 3. Cycles 3 and 4 in the galvanostatic charge-discharge curve of TiO 2 deposited on the Al nanostructured current collector. Cycled at 1 μA.

Figure 4 .

 4 Figure 4. Normalized capacity (mAh/cm 2 of geometrical surface area) for the 3D microbattery with TiO 2 deposited on an Al nanorod current collector and the 2D electrode with TiO 2 deposited on a flat Al plate. The galvanostatic cycling current is 0.001 mA.

To whom correspondence should be addressed. E-mail: Kristina.edstrom@ mkem.uu.se. Tel: +46-18-4713713. Fax: +46-18-513548.† Department of Materials Chemistry, Uppsala University. ‡ Universite ´de Toulouse. § Department of Engineering Sciences, Uppsala University.

Acknowledgment. This work is supported by the Swedish Research Council (VR), ALISTORE-ERI within EU-FP6 framework, and SUPERLION within EU-FP7. We also acknowledge the EUROPEAN ERASMUS MUNDUS PRO-GRAM from the European commission and the French Ministry of Research for, respectively, the Erasmus Mundus Grant of S.K.C. and the Ph.D. Grant of E.P.