Enhancing Total Conductivity of La2NiO4+δ Epitaxial Thin Films by

Jun 27, 2008 - High quality epitaxial c axis oriented La2NiO4+δ thin films have been prepared by the pulsed injection metal organic chemical vapor de...
0 downloads 0 Views 555KB Size
Download PDF
10982

J. Phys. Chem. C 2008, 112, 10982–10987

Enhancing Total Conductivity of La2NiO4+δ Epitaxial Thin Films by Reducing Thickness Mo´nica Burriel,† Jose´ Santiso,† Marta D. Rossell,‡ Gustaaf Van Tendeloo,‡ Albert Figueras,† and Gemma Garcia*,§ ICMAB and Centro de InVestigacı´on en Nanociencia y Nanotecnologı´a (CSIC), Campus UAB, 08193 Bellaterra, Spain, EMAT, UniVersity of Antwerp, Groenenborgerlaan, B-2020 Antwerp, Belgium, and Group de Nanomaterials i Microsistemes, Departament de Fı´sica, UAB, 08193 Bellaterra, Spain ReceiVed: October 19, 2007; ReVised Manuscript ReceiVed: April 16, 2008

High quality epitaxial c axis oriented La2NiO4+δ thin films have been prepared by the pulsed injection metal organic chemical vapor deposition technique on different substrates. High-resolution electron microscopy/ transmission electron microscopy has been used to confirm the high crystalline quality of the deposited films. The c-parameter evolution has been studied by XRD as a function of time and gas atmosphere. The high temperature transport properties along the basal a-b plane of epitaxial La2NiO4+δ films have been measured, and the total conductivity of the layers has been found to increase as the thickness is reduced. Layers of 50 nm and thinner have shown a maximum conductivity larger than that measured for single-crystals, in particular, the 33 nm thick films with a conductivity of 475 S/cm in oxygen correspond to the highest value measured to date for this material. 1. Introduction The use as cathodes for solid oxide fuel cells (SOFCs) of mixed ionic and electronic conducting (MIEC) oxides instead of the conventional electron conducting LaSrMnO3 (LSM) or composite LSM-yttria stabilized zirconia (YSZ) cathodes leads to enhanced performances, since the triple phase boundary between gas phase and oxide ionic and electronic conducting material, necessary for the electrochemical oxygen reduction, extends to the whole cathode surface. Therefore, they have attracted great attention as promising candidates for their use not only as cathodes, but also in several other electrochemical industrial applications, such as oxygen permeation membranes, catalytic oxidation systems, gas sensors, and so forth. Recently, the perovskite-related oxides of the K2NiF4 (A2BO4+δ)-type structure, such as La2NiO4, have attracted the interest of the scientific community.1,2 This structure is formed by stacks of perovskite ABO3 layers separated by AO rocksalt-type layers. These La2NiO4-based ceramics prepared as polycrystalline pellets or single crystal have shown high oxygen diffusivity at high temperatures,3–15 along with a high electronic conductivity, making them attractive candidates as mixed ionic and electronic conducting materials. The layered structure is responsible for the great flexibility in the oxygen stoichiometry as it permits the incorporation of oxygen excess (δ) involving interstitial sites along the rocksalt layers, therefore, rendering a high ionic conductivity to these materials. In the case of La2NiO4+δ, as occurs to all of the K2NiF4-type oxides containing both rare earth and transition metals, an electric field gradient is present between the La2O22+ and the NiO22- layers. Therefore, the Coulomb potential may prevent the (large) interstitial O2- ions from leaving the rocksalt layer, leading to large diffusion anisotropy.1 The ionic conductivity based on interstitial defects represents an attractive alternative to the vacancy-based conduc* Corresponding author. Phone: +34 935811481. Fax: +34 581 2155. E-mail: [email protected]. † ICMAB and Centro de Nanociencia y Nanotecnologı´a. ‡ University of Antwerp. § Departament de Fı´sica, UAB.

tion mechanism primarily present in perovskites and fluorite oxides, where defect interactions and associations limit the conductivity. The electrical conductivity of La2NiO4+δ is characterized by a semiconductor-type electronic conductivity, which occurs via hopping of p-type charge carriers between mixed-valence nickel cations, with a small activation energy (typically in the range 50-100 meV) below 600 K.8,11,14,15 They show a maximum between 600 and 700 K, above which a smooth apparent change from semiconducting to metallic-like behavior is observed.6,11,16 Bassat et al.13 presented evidence that the resistivity increase, previously attributed to a metal-insulator transition above 600 K, was actually due to the decrease of the hole density related to an oxygen loss in the structure. The reported values for the maximum conductivity (of 80 S/cm for polycrystalline bulk ceramics and films 8,10,12,13,17,18 and around 200 S/cm for single crystal samples along the a-b plane 3,11) turn this material into a very interesting cathode for SOFCs. The expected high anisotropy of the electrical properties was experimentally confirmed by the studies on single crystals.2 Because of this large anisotropy, final applications of this material with enhanced performance will strongly depend on the ability of preparing highly textured samples. In this work, we present the results on the preparation and characterization of epitaxial La2NiO4+δ thin films by the liquid pulsed injection metal organic chemical vapor deposition technique (PIMOCVD) onto SrTiO3 (100) and NdGaO3 (110) substrates. The substrate orientations and the deposition conditions have been selected in order to obtain fully oriented c axis films, as confirmed by X-ray diffraction (XRD). High-resolution electron microscopy/transmission electron microscopy (HREM/ TEM) characterizations have been used to confirm the high crystalline quality, as well as the epitaxial growth relationship with the substrate. Preliminary studies of the structure evolution as a function of temperature in oxygen and nitrogen have also been performed. Total planar conductivity measurements along the a-b plane in different oxygen partial pressures and as a function of film thickness have been carried out. The conductiv-

10.1021/jp7101622 CCC: $40.75  2008 American Chemical Society Published on Web 06/27/2008

Enhancing Total Conductivity of La2NiO4+δ

J. Phys. Chem. C, Vol. 112, No. 29, 2008 10983

ity has been found to increase with reducing the film thickness, and to our knowledge the measured conductivity values for the thinnest films correspond to the largest values reported so far. 2. Experimental Section Films were grown by PIMOCVD in a setup reactor described elsewhere.19 Commercial solid β-diketonates La(thd)3 and Ni(thd)2 (thd ) tetramethylheptanedionate) from STREM Chemicals mixed and solved in 1,2-dimethoxyethane were used as metal precursors. The total concentration of the solution was 0.02 M while the molar ratio La/Ni in the precursor solution was fixed to 3, as the result of a previous calibration.20 The evaporation and deposition temperatures were fixed to 280 and 750 °C, respectively. The total gas flow was maintained at 1 L/h with a 1:1 ratio of oxygen and nitrogen. The opening time of the injector was 2 ms with a pulse frequency of 1 Hz. Two different single crystal substrates have been used to grow the films: SrTiO3 (100) (noted STO) and NdGaO3 (110) (noted NGO). STO has a cubic structure with cell parameter of 3.905 Å, and NGO is orthorhombic with pseudocubic in-plane parameters of ap ) 3.863 Å and bp ) 3.854 Å for the 110 plane cut. These substrates and orientations were chosen in order to promote epitaxial growth of the layers due to their low mismatch: 1% on STO and about -0.19% on NGO substrates (values calculated in comparison with La2NiO4.18 orthorhombic with a ) 5.4652, b ) 5.4686, c ) 12.678 Å; basal plane parameter ap ) 3.866 Å).21 The phase composition, structure, and orientation of the films were analyzed by wavelength dispersive spectrometry (WDS) microprobe analysis and X-ray diffraction. The lattice parameters were accurately determined performing area scans of asymmetric reflections of the layers using a Bruker D8 DISCOVER with a four-angle goniometer and GADDS detector. Reciprocal space maps were systematically recorded for both film and substrate peaks and fitted to a Gaussian curve to extract the in- and out-of-plane cell parameters. XRD experiments at high temperature were performed in a Siemens D500 powder diffractometer under a controlled atmosphere of either pure air or N2 gas. Electron diffraction (ED) and HREM studies were performed using a JEOL 4000EX microscope. The planar resistance of the films was measured using a HP 4192A impedance analyzer at 1 kHz during the cooling process from 700 °C to room temperature in which different temperature levels were programmed each 50 °C to allow data acquisition at a constant and stabilized temperature. Measurements were performed either in pure oxygen, in pure nitrogen, or in a controlled O2/N2 mixture. Two painted parallel silver contacts were used as electrodes. 3. Results and Discussion 3.1. Structural Characterization. Epitaxial La2NiO4+δ (LNO) thin films of different thickness were deposited on STO and NGO by varying the number of droplets injected. The film thickness was determined by X-ray reflectivity for the thinnest films, corresponding to a growth rate of 1.0 ( 0.2 Å per droplet. By varying the number of droplets injected from 200 to 3000, we prepared films ranging from 20 to 335 nm, approximately. The high crystalline quality of the La2NiO4+δ deposited thin films were confirmed by cross section TEM images of the film/ substrate interfaces along with the corresponding ED patterns, as can be seen in Figure 1. The ED patterns are formed by a superposition of the patterns from substrate and film. The more intense reflections correspond to the substrate, whereas the weaker reflections are due to the film. For all of the observed ED patterns, the reflections from the film can be indexed in a

Figure 1. (a) Cross section TEM image of a c axis oriented LNO film deposited on NGO (110). (b) Cross section HREM image of a c axis oriented LNO film deposited on NGO (110) and corresponding ED pattern, along the [100] direction. Antiphase boundary generated by a surface step on the substrate is marked with an arrow.

tetragonal unit cell with cell parameters equal to those determined from the XRD data. In the real space, the visible periodic layers along the c axis correspond to the La2NiO4+δ layer sequence NiO2, LaO, LaO, NiO2. Figure 1b reveals a sharp and well-defined substrate/film interface, in which it is possible to observe one substrate surface step (indicated by a white arrowhead) produced by the miscut angle. These steps generate antiphase boundaries (APB) in the film that propagate during growth reaching the La2NiO4+δ film surface, as can be clearly observed in the low magnification top image (Figure 1a). The APBs are periodically separated by approximately 50 to 60 nm. As previously mentioned, by varying the number of pulses during the deposition, epitaxial LNO films of different thickness were deposited on STO and NGO. The in-plane (a and b) and out-of-plane (c) parameters were determined by measuring the angular position of the 2210 reflection of the LNO in the reciprocal space maps (equivalent to the 203 reflection for the cubic perovskite structure) like those shown in Figure 2 for films of about 350 nm deposited on STO and NGO. In order to avoid systematic errors in the measurements, the position of the 203 STO and 334 NGO reflections of the substrate were used as a reference. In the NGO case, as the cell is orthorhombic, we distinguished two different in-plane cell parameters, a and b, along the [001] and [-110] directions in the NGO substrate, respectively, for which we used two different reference reflections, namely, 334 and 150. The c parameters were also determined from the 0014 reflection positions in a standard θ-2θ configuration. The reciprocal space maps are expressed in Qx, Qy reciprocal coordinates (reciprocal lattice units: 1 rlu ) 1.298 Å-1), where Qx and Qy are calculated from the measured angular positions following the equations:

Qx ) sin θ sin(θ - ω) Qy ) sin θ cos(θ - ω)

(1)

where 2θ is the detector angular position and ω is the incidence angle at which the reflection from reciprocal lattice points hkl

10984 J. Phys. Chem. C, Vol. 112, No. 29, 2008

Burriel et al.

Figure 2. Reciprocal space maps of 2210 LNO reflection of 350 nm thick films deposited on STO and NGO substrates, along with the 203 STO and 334 NGO reference substrate reflections.

Figure 3. Variation of the in-plane cell parameter and strain with thickness for as-deposited LNO films on both substrates.

can be observed. The in- and out-of-plane cell parameters, a and c, respectively, were calculated from eq 1 applying the following equations:

a)

λh λh ) 2Qx 2 sin θ sin(θ - ω)

c)

λl λl ) 2Qy 2 sin θ cos(θ - ω)

(2)

where λ ) 1.5406 nm for the Cu KR wavelength and h and l ) Miller reflection indexes hkl, equal to 2 and 10 in this case (2210 reflection). The 203 STO and 334 NGO reflections from the substrate clearly show two components: KR1 and KR2, while the La2NiO4+δ 2210 reflection consist of a single peak although slightly asymmetric, also due to the two KR components of the X-rays. The a/b in-plane and c out-of-plane lattice parameters determined as a function of the layer thickness are shown in Figures 3 and 4, respectively. In LNO/NGO, and as a general trend, the a parameters increase from 3.849 Å at 22 nm to 3.856 Å at 373 nm, and b increases from 3.856 Å at 33 nm to 3.860 Å at 330 nm, as depicted in Figure 3, while the opposite behavior is observed for the LNO/STO cell parameters, where a decreases from 3.890 Å at 33 nm to 3.877 Å at 350 nm. A

similar, but inverted dependence is observed in Figure 4 for the c parameter that decreases from 12.719 Å at 33 nm to 12.697 Å at 375 nm on LNO/NGO, whereas it increases from 12.617 Å at 33 nm to 12.646 Å at 350 nm, on LNO/STO. The variation of the cell parameters with thickness was related to the in-plane strain effect induced by the film/substrate mismatch and depends on the selected substrate surface structure. The thinnest films on NGO are submitted to a biaxial compressive strain in the a-b plane, while the strain is tensile in the case of STO, in accordance with the expected mismatch. Therefore, for both substrates, thinner films are strained (in an opposite direction for each substrate), while thicker films are more relaxed and tend to common cell parameter values. From these experimental values, we also estimated the films strain as a function of the thickness (right axes in Figures 3 and 4) using the following equations:

Exx )

as - a0 bs - b0 × 100 Eyy ) × 100 a0 b0 Ezz )

cs - c0 × 100 (3) c0

where Exx, Eyy, and Ezz are the strain in the x, y (a-b plane) and z (c axis) directions; as, bs, and cs are the strained cell parameter values, and a0, b0, and c0 are the equilibrium cell parameters values (a0 ) 3.8645 Å, b0 ) 3.8670 Å, c0 ) 12.678 Å) selected from the structure reported by Metha and Heaney for La2NiO4.18 (δ ) 0.18),21 and consistent with the observed tendency of the cell parameters for the thicker films. We have measured a tensile strain in the a-b plane for the La2NiO4+δ films grown on STO, which decreases with thickness from approximately 0.7 to 0.3%, and a quite smaller compressive strain, which varies from approximately -0.4 to -0.2% for films grown on NGO (Figure 3). In the perpendicular direction, the opposite behavior occurs: for films deposited on STO, a strain from approximately -0.5 to -0.3% is observed, while on NGO the strain varies from approximately 0.4 to 0.2% (Figure 4). As expected, the natural evolution of thick films tends to fully relax cell parameters (a0, b0, c0). We have to point out that there has been observed a considerable spread in the parameter

Enhancing Total Conductivity of La2NiO4+δ

J. Phys. Chem. C, Vol. 112, No. 29, 2008 10985

Figure 4. Variation of the out-of-plane cell parameter and strain with thickness for as-deposited LNO films on both substrates.

determination, in particular, for LNO/STO, as it is shown in Figure 3 and 4. This might be related either to oxygen content deviations in the as-deposited samples or to the slightly lower epitaxial quality observed in LNO/STO films. This is evidenced by the elliptical shape of the 2210 reflection in the reciprocal space map of the LNO film on STO (in Figure 2), which is an indication of the larger orientation spread in the c axis domains for films on STO than those deposited on NGO. The misalignment of the crystalline domains might influence the way the oxygen is incorporated into the structure, thus, modifying the oxygen composition. In bulk or ceramic samples, it is known from literature that there is a correlation between lattice parameters and oxygen content, that is, the higher the oxygen content, the larger the c parameter.18,22,23 However, the fact that the epitaxial films are strained makes it very difficult to establish such a correlation. Unfortunately, because of the small mass of our samples, standard methods to determine the oxygen stoichiometry could not be applied. Thereby, we could not assign any δ values to our films or even relate the cell parameter changes with any oxygen content. 3.2. Total Conductivity at High Temperature. Planar transport measurements have been performed as a function of temperature on films deposited on STO and NGO substrates with thicknesses ranging from 33 to 335 nm. In addition to pure oxygen, the in-plane conductivity of our samples was also measured in air and nitrogen atmospheres. As can be seen in Figure 5, the total conductivity increases linearly with temperature following a thermally activated behavior, reaching maximum values around 400-450 °C. As expected, at higher temperatures (g450 °C), the material looses oxygen leading to a Ni3+ content decrease along with a reduction of carrier density.6,7,14,17 Furthermore, the conductivity in oxygen shows slightly higher values than in air and considerably higher values than in nitrogen. The dependence with temperature and oxygen partial pressure confirms that the conductivity is predominantly p-type electronic and occurs via a small-polaron mechanism.6,13 Figure 6, shows the Arrhenius representation of the conductivity dependence on the reciprocal temperature measured in pure oxygen atmosphere for the different film thicknesses. We have to notice that the conductivity of the substrate measured under the same conditions is several orders of magnitude smaller than that of the layers: ∼10-3 S/cm for STO and ∼10-6 S/cm (for NGO conductivity) at 700 °C. Thus, even for the thinner films, the conductance of the substrate (0.5 mm thick) is negligible, and the conductivity of the layer can be measured without any significant influence of the substrate. Conductivity values reported for single crystals and polycrystalline samples are

Figure 5. Logarithmic representation of σ*T of LNO epitaxial films as a function of the gas atmosphere for a 33 nm thick film deposited onto STO (left) and onto NGO (right).

Figure 6. Total conductivity of LNO epitaxial films as a function of the inverse of the temperature for films of different thickness deposited onto STO (left) and NGO (right) in oxygen atmosphere and comparison with literature data for single crystals12 and bulk ceramics in air.

plotted for comparison.11,18 We can clearly observe that, independent of the thickness, the film conductivity values are higher than the bulk polycrystalline samples measured by Amow and Skinner.18 Only the thickest films (335 nm) deposited on STO showed similar conductivity values to that of the bulk material at low temperatures. Furthermore, films below 50 nm thick and grown on both substrates showed even larger conductivities than that of La2NiO4+δ single crystals along the a-b plane. The maximum conductivity, obtained for the 33 nm thick films in oxygen atmosphere is close to 475 S/cm, far above the reported values to date for polycrystalline bulk ceramics and films (around 80 S/cm)8,10,12,13,18,21,24 and single crystal (200 S/cm).2,11 This fact seems to confirm the high quality of our samples. Since the transport properties are highly anisotropic, it is easy to understand that conductivities along the a-b plane of our c axis oriented films are larger than those measured on randomly oriented polycrystalline ceramics or films. The corresponding values of activation energy (Ea) have been calculated

10986 J. Phys. Chem. C, Vol. 112, No. 29, 2008

Burriel et al.

TABLE 1: Activation Energies of the Conductivity (σT) for the Films with Different Thickness Deposited on STO and NGO Substrates thickness (nm)

LNO/STOa(meV)

LNO/NGOa(meV)

33 50 200 335

62 81 84 87

75 76 79 106

a

Measured in O2 atmosphere.

from the measured conductivity σ using the adiabatic small polaron hopping model:

σ)

A0 Ea exp(- ) T RT

(4)

The apparent activation energies as a function of thickness for the temperature region below 400 °C are shown in Table 1. The small activation energies obtained (from 62 to 106 meV) fall in the typical range found by other authors below 600 K.1,2,10,13 On both substrates, the activation energy decreases slightly when reducing the film thickness. The reason for the high conductivity values obtained for very thin films is still unknown. One possibility is that the higher conductivity might be related to a higher oxygen content, which would increase the hole carrier density. However, as we noted in a previous paragraph, there is no estimate of the oxygen content in the films. Another possibility is that the strain, either tensile (on STO) or compressive (on NGO), increases the electronic carrier mobility. This is consistent with the apparent lower activation energies (slope of the conductivity curve) for the thinner films on both STO and NGO substrates in Table 1. The microstructure of the films with the presence of APBs, as observed by HREM in Figure 1, could also play a role in the transport properties of the films. However, the density of those defects seems to depend only on the substrate surface steps density and is thus independent of the film thickness. Any implication of such microstructure defects on the observed variations of the transport properties can thus be neglected. In summary, many parameters affecting the transport mechanisms of the La2NiO4+δ epitaxial thin films in a competing or cooperative way seem to give rise to an optimal strained state with unexpectedly high properties, compared with the bulk and single crystal relaxed material. Whether the observed increase is related to the carrier density or mobility variations still remains unresolved. 3.3. Chemical Stability Studies. To complete the study of the La2NiO4+δ structure and its stability in different gas atmospheres, we have measured the evolution of the c parameter at 973 K (700 °C) by switching the gas atmosphere from air to nitrogen, as shown in Figure 7. These values have been determined from the LNO 0010 reflection position and the 300 or 330 reflections of the STO and NGO substrates, respectively. The films were heated up to 973 K (700 °C) in air and, once this temperature stabilized, an XRD pattern was acquired every 8 min. After approximately 1 h, the gas was rapidly switched to pure nitrogen, and a similar process was repeated for another 60 min. This procedure was repeated two or three times depending on the sample. Different observations can be extracted from this experiment. The c parameter seems to be fairly stable in air for the time period studied, whereas in nitrogen atmosphere, it decreases continuously with annealing time. This structural change seems to be reversible, since after a new gas change to air, the c parameter recovers its previous value, except

Figure 7. Evolution of the c parameter as measured by XRD at 973 K by switching the gas atmosphere from air (open symbols) to nitrogen (filled symbols), for 33 and 330 nm thick LNO films deposited on NGO and STO substrates.

for the thinner film on NGO. This irreversibility may be due to partial strain release in the film, which showed the largest strain in the as-deposited state. The stabilization kinetics of the c parameter after a gas change depends on the final gas atmosphere and film thickness. The c parameter stabilizes in a few seconds when the gas atmosphere is switched from nitrogen to air (oxidation), while from air to nitrogen (reduction) the c parameter decreases continuously and takes 15 to 30 min to reach a quasi-stable value. Since the c parameter varies because of the oxygen exchange with the surrounding atmosphere, we can deduce that the oxygen incorporation in the lattice is kinetically more favorable than the oxygen loss. From the differences between the 33 and the 330 nm films, we could say that oxygen exchange in thick (relaxed) films seems to be kinetically more favorable than in thinner (strained) ones, at least for the reduction process. In order to simulate realistic working conditions, the films were submitted to high temperature treatment under an oxidizing atmosphere for longer periods but without applying an electrical field. A thick La2NiO4+δ/NGO film (330 nm) was heated at 600 °C during 400 h, followed by a 800 °C heat treatment for the same period of time. The aging effects on the structure were characterized by XRD patterns taken at room temperature before and after each treatment. No major structure variations were detected after the heat treatment at 600 °C for 400 h. Nevertheless, after the heat treatment at 800 °C, a broadening and loss of intensity of the 00l La2NiO4+δ reflections were observed, along with the appearance of some weak peaks assigned to La2O3 impurity phase. In order to study the stability of the electronic transport in the films, the conductivity at high temperature was recorded for long periods. For this purpose, we have measured the conductivity of a 200 nm thick La2NiO4+δ film deposited on NGO at 574 °C as a function of time switching periodically the gas atmosphere from oxygen to nitrogen, as shown in Figure 8. Initially, the sample was heated up in oxygen, and once the conductivity was stabilized, the atmosphere was switched to pure nitrogen. The atmosphere changes were accompanied by a decrease in the conductivity. After 80 min, the atmosphere was switched again to pure oxygen, and this procedure was repeated 3 times before stabilizing the sample again in pure oxygen. At this point, we observe a slight conductivity decrease of about 2% from its initial value. No further variation was observed in oxygen for at least 14 h. The procedure was repeated

Enhancing Total Conductivity of La2NiO4+δ

J. Phys. Chem. C, Vol. 112, No. 29, 2008 10987 enhancement seems to be the thickness and the absolute strain value. From the evolution of the c parameter as a function of the gas atmosphere, we have observed slower oxygen exchange kinetics in nitrogen, compared with air. The films were stable in oxidizing atmospheres, especially at intermediate temperatures (e600 °C). As a general conclusion, epitaxial thin La2NiO4+δ films present more attractive properties than bulk material, making them adequate for their use in electrochemical devices.

Figure 8. Variation of the conductivity with time at 847 K by changing the gas atmosphere from air to nitrogen, for a 200 nm LNO film deposited on NGO.

again, stabilizing the sample in nitrogen atmosphere. In this case, a continuous drop in the conductivity was observed during the experiment elapsed time. After a new gas switch cycle and the sample stabilization in oxygen, we observed a conductivity loss from its initial value estimated at 18%. This change with cycles seems to be an irreversible process. This irreversibility might be associated either to the c parameter changes observed at high temperature in nitrogen (in Figure 7) or to the film composition segregation. The XRD patterns taken after this experiment showed again the appearance of small peaks corresponding to the La2O3 phase, similar to those observed after anneal in air at 800 °C during several weeks. This lanthanum oxide could derive from the decomposition of La2NiO4 in LaNiO3 plus La2O3, but as the LaNiO3 has similar parameters than the substrate, we were not able to detect by XRD the possible presence of these perovskite impurities. 4. Conclusions By using the PIMOCVD technique, we have deposited epitaxial c axis La2NiO4+δ films onto NGO and STO substrates. The high quality of the layers was assessed by HREM. We have characterized by XRD the variation with thickness of the inand out-of-plane cell parameters showing an evolution from thin strained to thick relaxed films on both substrates. Planar transport measurements allowed us to determine the conductivity of La2NiO4+δ films along the a-b plane. We have observed a semiconducting type behavior for temperatures up to 425 °C, and the activation energy of the thermally activated process has been found to increase with film thickness. The conductivity has been found to increase with reducing thickness showing a maximum conductivity of 475 S/cm for 33 nm thick films, corresponding, to our knowledge, to the highest value reported to date for this compound. The physical mechanism explaining the higher conductivity values measured for the thinner films is still unclear as no real evidence of the effect of lattice parameters or oxygen stoichiometry has been found. Actually, the only parameters playing a role in the total conductivity

Acknowledgment. This work was partially funded by the MAT2005-02601 project from the Spanish MEC, and the Spanish Fuel Cell Network. M.B. acknowledges the I3P program for financial support. Authors thank Dr. J. Bassas from the SCTUB for high temperature XRD measurements. References and Notes (1) Bassat, J. M.; Gervais, F.; Odier, P.; Loup, J. P. Mater. Sci. Eng., B 1989, 3, 507–514. (2) Rao, C. N.R.; Buttrey, D. J.; Otsuka, N.; Ganguly, P.; Harrison, H. R.; Sandberg, C. J.; Honig, J M. J. Solid State Chem. 1984, 51, 266– 269. (3) Bassat, J. M.; Odier, P.; Villesuzanne, A.; Marin, C.; Pouchard, M. Solid State Ionics 2004, 167, 341–347. (4) Minervini, L.; Grimes, R. W.; Kilner, J. A.; Sickafus, K. E. J. Mater. Chem. 2000, 10, 2349–2354. (5) Kharton, V. V.; Viskup, A. P.; Naumovich, E. N.; Marques, F. M. B. J. Mater. Chem. 1999, 9, 2623–2629. (6) Kharton, V. V.; Yaremchenko, A. A.; Shaula, A. L.; Patrakeev, M. V.; Naumovich, E. N.; Logvinovich, D. I.; Frade, J. R.; Marques, F. M. B. J. Solid State Chem. 2004, 177, 26–37. (7) Skinner, S. J.; Kilner, J. A. Solid State Ionics 2000, 135, 709–712. (8) Ishikawa, K.; Shibata, W.; Watanabe, K.; Isanaga, T.; Hashimoto, M.; Suzuki, Y. J. Solid State Chem. 1997, 131, 275–281. (9) Boehm, E.; Bassat, J. M.; Steil, M. C.; Dordor, P.; Muavy, F.; Grenier, J. C. Solid State Sci. 2003, 5, 973–981. (10) Sayer, M.; Odier, P. J. Solid State Chem. 1987, 67, 26–36. (11) Dembinski, K.; Bassat, J. M. ; Coutures, J. P.; Odier, P. J. Mater. Sci. Lett. 1987, 6, 1365–1367. (12) Vashook, V. V.; Tolochko, S. P.; Yushkevich, I. I.; Makhnach, L. V.; Kanomyuk, I. F.; Altenburg, H.; Hanck, J.; Ullmann, H. Solid State Ionics 1998, 110, 245–253. (13) Bassat, J. M.; Loup, J. P.; Odier, P. J. Phys. Condens. Mater. 1994, 6, 8285–8293. (14) Odier, P.; Nigara, Y.; Coutures, J. J. Solid State Chem. 1985, 56, 32–40. (15) Bassat, J. M.; Odier, P.; Villesuzanne, A.; Marin, C.; Pouchard, M. Solid State Ionics 2004, 167 (3-4), 341–347. (16) Shinomori, S.; Kawasaki, M.; Tokura, Y. Appl. Phys. Lett. 2002, 80 (4), 574–576. (17) Skinner, S. Solid State Sci. 2003, 5, 419. (18) Amow, G.; Skinner, S. Solid-State Electron. 2006, 10, 538. (19) Burriel, M.; Garcia, G.; Santiso, J.; Hansson, A. N.; Linderoth, S.; Figueras, A. Thin Solid Films 2005, 473 (1), 98–103. (20) Burriel, M.; Garcia, G.; Rossell, M. D.; Figueras, A.; Van Tendeloo, G.; Santiso, J. Chem. Mater. 2007, 19 (16), 4056–4062. (21) Mehta, A.; Heaney, P. J. Phys. ReV. B 1994, 49 (1), 563–571. (22) Rice, D.; Buttrey, D. J. Solid State Chem. 1993, 105, 197. (23) Jorgensen, J.; Dabrowski, B.; Pei, S.; Richards, D.; Hinks, D. Phys. ReV. B 1989, 40 (4), 2187–2199. (24) Aguadero, A.; Alonso, J. A.; Martı´nez-Lope, M. J.; Ferna´ndezDı´az, M. T.; Escudero, M. J.; Daza, L. J. Mater. Chem. 2006, 16, 3402.

JP7101622