Article pubs.acs.org/cm
Large Magnetoresistance and Electrostatic Control of Magnetism in Ordered Mesoporous La1−xCaxMnO3 Thin Films Christian Reitz,† Philipp M. Leufke,† Reinhard Schneider,‡ Horst Hahn,†,§ and Torsten Brezesinski*,† †
Institute of Nanotechnology, Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany ‡ Laboratory for Electron Microscopy, Karlsruhe Institute of Technology, Engesserstr. 7, 76131 Karlsruhe, Germany § Helmholtz Institute Ulm, Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany S Supporting Information *
ABSTRACT: Ferroic nanomaterials have received much attention in recent times because of their potential for novel applications. Here, we report a facile bottom−up synthetic approach to the first ordered mesoporous mixed-valence manganese oxide. Continuous thin films of perovskite-type La0.68Ca0.30Mn1.02O3−δ have been prepared from common inorganic salt precursors by taking advantage of the superior templating properties of a polyisobutylene-block-poly(ethylene oxide) diblock copolymer. This novel solution-processed mesostructured material is well-defined at the nanometer and micrometer levels and behaves superparamagnetically above 230 K. Furthermore, it exhibits large magnetoresistance over a wide range of temperatures due to complex percolation pathways for electron transport imparted by the unique pore−solid architecture, and it is ferromagnetic metallic below 180 K. We also demonstrate reversible magnetization modulation of up to 8.5% by electrostatic charge carrier doping in the polymer-templated thin films. This value is the highest thus far reported for electrolyte-gated mixed-valence manganese oxides.
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INTRODUCTION Perovskite-type mixed-valence manganese oxides of the form A1−xBxMnO3 (with A = rare earth, B = alkaline earth), such as La1−xCaxMnO3 (LCMO) and La1−xSrxMnO3 (LSMO), are highly interesting from the viewpoint of material science and have been studied for more than 6 decades.1−3 This is due, in part, to the fact that they exhibit interesting magnetotransport properties and have a rich phase diagram that includes, among others, ferromagnetic metallic and insulating phases.4−7 In general, these multimetallic oxides can be considered as solid solutions between A3+Mn3+O3 and B2+Mn4+O3, that is, A-site doping leads to mixed-valence states of the manganese ions in the lattice. Over the years, several studies have demonstrated that the magnetic behavior is governed by the ratio of Mn4+ to Mn3+ (the Mn3+−O2−−Mn4+ coupling via double exchange is ferromagnetic in nature), and there is a strong correlation between the magnetic properties and the charge carrier transport through the material.8 After the discovery of colossal magnetoresistance (CMR) in perovskite-type mixed-valence manganese oxides in the early 1990s, much effort has been spent to understand the magnetotransport properties.9 Both in single crystals and epitaxially grown layers, the magnetoresistive effect occurs only at high fields near the ferromagnetic metallic to paramagnetic insulating phase transition, thus limiting their application potential.10 In contrast, polycrystalline forms of © 2014 American Chemical Society
these materials with small grain sizes may exhibit significant magnetoresistance at low fields over a wide temperature range.11 This low-field effect is technologically important and has been shown to be strongly determined by electron scattering at grain boundaries and magnetic domain walls.12 As a result, it can be tailored to a large extent through grain boundary engineering, which is one of the reasons magnetoresistive nanomaterials and composites thereof have been of interest in recent times, with much attention paid to interface effects.13−15 Moreover, it has been reported recently by Mishra and co-workers that the magnetization of electrolyte-gated nano-LSMO can be controlled by electrostatic modulation of the surface charge carrier density.16,17 This is an interesting result, as it might pave the way for novel applications of LSMO and other related materials. They showed that the relative change in magnetization upon capacitive charging (i.e., formation of a Helmholtz double-layer at the solid/liquid electrolyte interface) can be as large as 2.5%, and this kind of electrostatic carrier doping is reversible, unlike chemical doping. Since the magnetization modulation is dependent upon the change in carrier density, decreasing the grain size of the Received: July 31, 2014 Revised: September 4, 2014 Published: September 18, 2014 5745
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sample-to-detector distance of 26 cm). For Rietveld refinement of GIXRD data, FullProf software was used. X-ray photoelectron spectroscopy (XPS) data were acquired on a VersaProbe PHI 5000 Scanning ESCA Microprobe from Physical Electronics equipped with a monochromatic Al Kα X-ray source and a hemispherical electron energy analyzer at an electron takeoff angle of 45°. The C 1s signal from adventitious hydrocarbon at 284.8 eV was used as an energy reference to correct for charging. Rutherford backscattering spectrometry (RBS) measurements were carried out using 2 MeV 4He+ ions accelerated by a Van de Graaff accelerator, with the parallel beam (at normal incidence) irradiating a cross-sectional area of 1 mm2. The backscattered ions were analyzed at an angle of 152°. The divergence of the ion beam (ion current of approximately 15 nA) was smaller than 0.02°. The solid angle of the detector was in the range of 10−3 sr. For data analysis, SIMNRA software was used. Magnetic susceptibility measurements were carried out on a Quantum Design MPMS XL-5 superconducting quantum interference device (SQUID) magnetometer in fields of up to ±45 kOe. For magnetotransport measurements, the standard vacuum plug sealing the sample tube of the SQUID magnetometer was replaced by a custom-made PMMA plug with LEMO 0S series straight coupler with 4 contacts as electrical feedthrough. For Van der Pauw four-probe resistance measurements, flexible Kapton-coated copper wires were used, and the LCMO thin films on quartz glass substrate were cleaved into smaller pieces (typically 0.5 × 0.5 cm2 in area) to fit the lateral dimension of the SQUID sample tube. Good electrical contact was established by fixing the Kapton-coated copper wires at the corners of the sample with temperature stable carbon epoxy resin (EC 261 C from Polytec PT GmbH). After curing the epoxy resin at 90 °C for 12 h, the copper wires were connected internally to the electrical feedthrough and externally to a Keithley SourceMeter 2601a. The SourceMeter was connected to a personal computer via a GPIB bus and controlled by the MPMS MultiVu application. Electrostatic carrier doping experiments using diethylmethyl(2methoxyethyl)ammonium bis(trifluoromethylsulfonyl)imide as nonaqueous electrolyte were performed in situ inside the SQUID magnetometer at T = 230 K. For these measurements, a custombuilt sample tube was utilized (see Figure S6 of the Supporting Information) and assembled in an argon-filled glovebox with [O2] < 1 ppm and [H2O] < 1 ppm to keep the impact of water and associated degradation effects at a minimum. Cyclic voltammetry was carried out in a two-electrode cell using an Autolab PGSTAT302 potentiostat. High surface area carbon fiber cloth (SBET ≈ 2400 m2/g) and a glass fiber filter disc (GF/A) from Whatman served as the counter electrode and separator, respectively.
material and introducing nanoscale porosity should increase the specific surface area and thus the overall effect. High-quality perovskite-type mixed-valence manganese oxide layers have been produced by several physical deposition techniques, such as sputtering, pulsed laser deposition (PLD), and others.18−20 Although these techniques allow control over both grain size and structure, they typically produce rather dense films with a low surface-to-volume ratio. In contrast, chemical deposition techniques, particularly sol−gel, have been shown to be well-suited for the preparation of functional nanoporous solids.21−29 However, reports on multimetallic oxide thin films having a regular pore network are scarce, mainly due to the lack of control over crystallization. In the present work, we describe for the first time the evaporation-induced self-assembly (EISA)30,31 synthesis of cubic mesoporous LCMO thin films using a polyisobutyleneblock-poly(ethylene oxide)32−34 diblock copolymer (referred to as PIB107-b-PEO150 hereafter) as the structure-directing agent. This novel mesostructured material can be described by the chemical formula La0.68Ca0.30Mn1.02O3−δ and exhibits large lowand high-field magnetoresistance because of the unique morphology. Furthermore, we demonstrate significant tuning of the magnetization upon double-layer charging, which is unprecedented for polymer-templated thin films.
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EXPERIMENTAL SECTION
Materials. Mn(OAc) 2 ·4H 2 O (99.999%), La(NO 3 ) 3 ·6H 2 O (99.99%), Ca(NO 3 ) 2 ·4H 2 O (99.9995%), diethylmethyl(2methoxyethyl)ammonium bis(trifluoromethylsulfonyl)imide (≥98.5%), glacial acetic acid (99.99%), ethanol, tetrahydrofuran, and 2-methoxyethanol were purchased from Sigma-Aldrich and ABCR. Polished 700 μm thick Si(001) wafers were purchased from SiMat. H[C(CH3)2CH2]107C6H4(OCH2CH2)150OH (PIB107-b-PEO150) was obtained from BASF SE and used as the polymer structure-directing agent. Thin Film Synthesis. A solution containing 50 mg of PIB107-bPEO150, 0.6 mL of ethanol, 0.15 mL of glacial acetic acid, 0.2 mL of tetrahydrofuran, and 1.3 mL of 2-methoxyethanol was combined with 127 mg of Mn(OAc)2·4H2O, 157 mg of La(NO3)3·6H2O, and 37 mg of Ca(NO3)2·4H2O. After stirring the colorless solution for 30 min, sub-micrometer-thick films were produced via dip coating on both Si(001) and quartz glass substrates at a relative humidity of 30%. After the films had been dried for 4−5 min at room temperature, they were transferred to an oven at 140 °C for 15 min and then heated to 300 °C using a 90 min ramp, followed by aging for 12 h. Lastly, the films were subjected to calcination at 750 °C for 2 min both to fully remove the polymer template and achieve crystallization of the amorphous wall structure. Methods. High-resolution transmission electron microscopy (HRTEM) was performed with an FEI Titan3 80−300 microscope operated at 300 keV. The microscope is housed in an enclosure to dampen acoustic and temperature variations and is equipped with a correction-lens system for spherical and other aberrations (Cs image corrector, CEOS GmbH). This allows obtaining HRTEM images with a point resolution of 0.08 nm. The residual objective lens aberrations were determined on the basis of the procedure described by Uhlemann and Haider.35 Images were taken by means of a 4 megapixel CCD camera (Gatan UltraScan 1000 P), which can be positioned on the optical axis. DigitalMicrograph (Gatan) software was used to control the camera. For each image, the exposure time was set at 0.5 s. In addition, the LCMO thin films were imaged in scanning TEM (STEM) mode using an electron probe of nominal 0.3 nm diameter. Field-emission scanning electron microscopy (SEM) images were recorded on a MERLIN from Carl Zeiss operated at 5 keV. Grazing incidence X-ray diffraction (GIXRD) measurements were carried out on a Bruker D8 Discover diffractometer equipped with a Soller slit and a LynxEye strip detector (wavelength of 0.15418 nm,
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RESULTS AND DISCUSSION As mentioned above, the films studied here were prepared by solution processing using a soft-templating approach. Several precursors were tested during the course of this work, and it turned out that the combination of hydrated manganese(II) acetate with hydrated lanthanum and calcium nitrate salts gives the best results in terms of nanoscale structure. In addition, the sol made from these reagents is stable over several days, thereby facilitating the synthesis. After sol−gel film deposition by dip coating on Si(001) or quartz glass substrates, the samples were directly subjected to a calcination treatment to get pure perovskite-type LCMO with a mesoporous morphology. The pore structure of the PIB107-b-PEO150-templated LCMO thin films with amorphous and crystalline walls was studied ex situ by means of electron microscopy. Panels a−d of Figure 1 show cross-sectional scanning electron microscopy (SEM) images at different magnifications collected after heating films at various temperatures for 2 min. These images show both 80° and 90° views of the polymer-templated structure and depict the evolution of the pore−solid architecture with increasing 5746
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have a thin solid sealing layer, as indicated by the SEM image in Figure 1a. The reason for the formation of such a surface layer is unclear at present. Figure 1 presents further bright-field (BFSTEM) and high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) and highresolution transmission electron microscopy (HRTEM) images of a PIB107-b-PEO150-templated LCMO thin film heated at 750 °C. These data corroborate the SEM results. Specifically, they confirm that the films are mesostructured with cubic pore symmetry, and the walls are highly crystalline with a randomly oriented grain structure [see also electron microscopy data, including selected-area electron diffraction (SAED), in Figure S1 of the Supporting Information]. The HRTEM image along the [001] zone axis in panel g shows the atom columns very clearly, with no amorphous regions. The fast Fourier transformation (FFT) pattern of this image provides a d-spacing of 2.76 Å, which matches that of the (002) lattice plane of orthorhombic LCMO.36 Taken together, the results from electron microscopy collectively demonstrate that LCMO can be templated to produce continuous perovskite-type thin films with a well-defined mesoporous morphology. This sol−gelderived material represents an ideal model system to study both the magnetoresistive properties and the impact of electrostatic carrier doping on magnetization as it contains a large amount of air−solid and solid−solid interfaces. In the following sections, we show data only for films that had been heated at 750 °C. The reason is that a high degree of crystallinity is crucial for achieving the desired physical properties and the structural stability of the samples in general. The crystallinity and phase composition as well as the chemical composition and state of the PIB107-b-PEO150-templated LCMO thin films were investigated via grazing incidence Xray diffraction (GIXRD), X-ray photoelectron spectroscopy (XPS), and Rutherford backscattering spectrometry (RBS). A representative GIXRD pattern is given in Figure 2a, confirming that the mesostructured material crystallizes in the orthorhombic crystal system, as expected for the nominal composition with x = 0.3. All of the diffraction peaks can be indexed to the orthorhombic phase in space group Pnma. A schematic of the unit cell, with the manganese ions occupying the octahedral 4b sites (Wyckoff notation), oxygens on the 4c and 8d sites, and both lanthanum and calcium ions on the 4c lattice sites, is shown in the inset of Figure 2a. The GIXRD data were also analyzed by Rietveld refinement to retrieve the lattice parameters and average crystallite size. The profile was fitted using the Thompson−Cox−Hastings pseudo-Voigt profile function, and the background level was fitted with a linear interpolation between 27 given points with refinable heights. The quality of the refinement was assessed both graphically by comparing the observed and calculated patterns and analytically by the magnitude of the weighted profile R-factor (Rwp = 9.44%) and the goodness-of-fit parameter (χ2 = 0.142). The fact that the patterns are in good agreement and the discrepancy values are low indicates that the fit is of high quality overall and the solution-processed LCMO is indeed single phase. The lattice parameters derived from the Rietveld analysis are a = 5.449(3) Å, b = 7.690(4) Å, and c = 5.4749(4) Å. The average crystallite size is 17(1) nm. This dimension is within the range of wall thicknesses determined via electron microscopy. The elemental composition, bonding configuration, and oxidation state of manganese were examined by XPS. There appears to be no detailed XPS studies on LCMO in the
Figure 1. Electron microscopy of the PIB107-b-PEO150-templated LCMO thin films heated at 550 °C (a), 650 °C (b), and 750 °C (c− g). Cross-sectional SEM images collected with a sample tilt of 90° (a, b) and 80° (c, d) show the structural evolution of the 3D cubic pore network with increasing calcination temperature. (e) BF-STEM and (f) HAADF-STEM images of the same film region. (g) [001] zone axis HRTEM image.
calcination temperature, particularly during the amorphous-tocrystalline transition. Films heated at 550 °C are fully amorphous, whereas those at 750 °C are highly crystalline; the temperature of 650 °C seems to mark the onset of crystallization. As can be seen from the SEM data, amorphous material exhibits a well-developed cubic mesostructure. The average pore size is 20 nm in-plane and 7 nm vertical. The wall thickness is in the range of 13−19 nm. The fact that the pores are highly anisotropic in shape is due to unidirectional lattice contraction caused by thermally triggered hydrolysis and condensation reactions. Also, it is evident that the films undergo restructuring throughout the crystallization process. However, the nanoscale porosity is preserved to a large extent, that is, the architecture shows no typical signs of structural collapse, as is often observed for polymer-templated nonsilicate oxides. Nevertheless, the initial periodicity is lost. Furthermore, we find that the films are crack-free on the micrometer level and the majority of pores at the top surface are open when the material is crystalline (see top view SEM image in Figure S1 of the Supporting Information); fully amorphous films tend to 5747
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ions resulting from the substitutional incorporation of divalent Ca in LCMO. The peaks at lower binding energies of (653.5 ± 0.05) eV and (641.9 ± 0.05) eV for the 2p1/2 and 2p3/2 levels, respectively, can be assigned to Mn3+, and those centered at (655.3 ± 0.05) eV and (643.7 ± 0.05) eV represent the Mn4+ oxidation state. The Mn4+ content from quantitative analysis matches the expected Mn4+ to Mn3+ ratio for La1−xCaxMnO3 with x = 0.3, thus indicating that charge compensation is predominantly achieved by formation of Mn4+ species rather than oxygen vacancies. We are aware that there is a certain error margin involved in these data. However, this margin should be comparatively small in absolute terms. The Ca 2p spectrum in Figure 2d also shows a doublet due to spin−orbit splitting. In contrast to the Mn 2p level, the 2p1/2 and 2p3/2 peaks at (349.9 ± 0.05) eV and (346.4 ± 0.05) eV, respectively, are symmetric, that is, they can be fitted as single-component peaks. The atomic La/Ca/Mn ratio was calculated by comparing the total integrated area of the peaks. In this way, a ratio of 0.70:0.28:1.02 is obtained, which is close to the targeted composition of La0.7Ca0.3Mn1.0O3. Since XPS is a surface analysis technique, RBS measurements were also carried out to get insight into the bulk chemical composition of the PIB107-bPEO150-templated LCMO thin films. Figure 2e shows an RBS spectrum taken with a 2 MeV 4He+ ion beam at normal incidence. Unfortunately, because of the silicon edge, a precise determination of the oxygen content in the films is virtually impossible. Quantitative analysis of the data reveals a La/Ca/ Mn ratio of 0.68:0.30:1.02. Accordingly, the films can be described by the chemical formula La0.68Ca0.30Mn1.02O3−δ, in agreement with the results from XPS. Overall, the data in Figure 2 lead us to conclude that the sol−gel-derived LCMO thin films are chemically pure and the presence of second phases can be ruled out within the detection limits of the techniques. Furthermore, Rietveld analysis indicates that there is no mismatch between crystallite size and wall thickness, which helps to explain why the mesoporous morphology is retained after the films had been heated at 750 °C. The in-plane magnetic properties of the PIB107-b-PEO150templated LCMO thin films were characterized by SQUID magnetometry. Zero-field-cooled (ZFC) and field-cooled (FC) magnetization curves at an applied field of 100 Oe are presented in Figure 3a. The ZFC data show a peak centered at Tmax = 188 K, whereas the FC magnetization continues to increase as the temperature is lowered further and begins to level off at about 30 K. In addition, it can be seen that the FC and ZFC curves diverge at Tirr = 196 K, that is, the films exhibit magnetic irreversibility below Tirr. This behavior is typical of superparamagnetic materials.40 Figure 3b shows magnetic hysteresis loops in the field range of ±2 kOe measured at 5, 100, and 200 K (see Figure S3a of the Supporting Information for all isothermal hysteresis curves collected in this work). The coercive fields (HC) are 670, 268, and 58 Oe, respectively. While a comparison with literature reports is difficult, the lowtemperature HC value appears to be slightly larger than that of other LCMO nanomaterials. A closer look at the data shows that HC varies as T1/2 (see Figure S3b of the Supporting Information). This kind of dependence is generally observed for noninteracting single-domain particles below the blocking temperature (TB). TB was determined by fitting the data according to HC = HC0·[1 − (T/TB)1/2], where HC0 is the coercive field at zero temperature.41 In so doing, we obtain TB ≈ 230 K, in agreement with the nonhysteretic behavior (unblocked state) found for T ≥ 230 K. As expected, the M(H)
Figure 2. Composition and chemical state of the PIB107-b-PEO150templated LCMO thin films heated at 750 °C. (a) Observed (black circles) and calculated (red line) XRD patterns showing that LCMO adopts an orthorhombic crystal structure. The difference curve is plotted in blue below the patterns. The presence of a spurious background peak is denoted with an asterisk. A schematic of the unit cell with oxygen in red, manganese in turquois, and lanthanum/ calcium in yellow is shown in the inset. (b−d) XPS detail spectra of the La 3d, Mn 2p, and Ca 2p core excitations. Solid black curves are fits to the data, and solid red curves are the sums of the fits. (e) RBS spectrum showing the composition corresponds to the formula La0.68Ca0.30Mn1.02O3−δ. The solid red curve is a fit to the data using SIMNRA.
literature.37,38 A wide-scan survey spectrum, with peaks evident only for lanthanum, manganese, calcium, oxygen, and carbon, is presented in Figure S2 of the Supporting Information. Detail spectra of the La 3d, Mn 2p, and Ca 2p levels are shown in panels b−d of Figure 2. These spectra were fitted using splitting parameters with relative area ratios of 1:2 and 2:3 (based on the degeneracy ratios) for the 2p and 3d component peaks, respectively, and the full width at half-maximum (fwhm) was constrained to be equal for all peaks of a given core level. As can be seen from Figure 2b, the La 3d spectrum is rather complicated. However, similar to the 3d core excitation in La(OH)3 and La2O3, the data can be fitted well assuming two final states.39 The main peaks at binding energies of (850.4 ± 0.05) eV and (833.7 ± 0.05) eV are due to emissions from the 3d3/2 and 3d5/2 levels (final state without charge transfer), respectively. The satellite structure at the higher binding energy side of the core level peaks can be deconvoluted into two components (final state with charge transfer from the oxygen 2p valence band to an empty lanthanum 4f orbital). The peaks at (852.4 ± 0.05) eV and (835.7 ± 0.05) eV represent the antibonding component and those at (854.7 ± 0.05) eV and (838.0 ± 0.05) eV the bonding component of the final state. We note that the energy separation of both the different peak pairs and the satellites from the main peaks was kept constant during the fitting procedure. The peaks centered at (847.0 ± 0.05) eV and (851.7 ± 0.05) eV can be attributed to the La 3d5/2 plasmon loss and Mn LMM Auger line, respectively. The Mn 2p photoelectron spectrum in Figure 2c contains an asymmetric doublet. This is due to the mixed valence of the Mn 5748
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electron scattering at the magnetic domain walls (grain boundaries) results in high overall electrical resistance. By applying a larger magnetic field, the electrical resistance drops sharply because of facilitated electron transport (spin-polarized tunneling) across grain boundaries. As a result, the MR effect increases. A schematic representation of the MR effect in the PIB107-b-PEO150-templated LCMO thin films is shown in Figure 4. After the saturation magnetization is reached, the
Figure 4. Schematic representation of the MR effect in the PIB107-bPEO150-templated LCMO thin films indicating that more current can be drawn through the film at constant voltage with applied magnetic field. It should be noted that the magnetic field direction was in-plane during the transport measurements; the out-of-plane configuration was chosen for graphical reasons.
magnetoresistance shows a linear dependence with applied field. Similar observations have been made for LCMO films in bulk form with grain sizes in the nanometer range and are attributed to the presence of a noncollinear magnetic surface layer.42 Furthermore, it is apparent that the absolute low-field magnetoresistance decreases with increasing temperature: 15% at 5 K, 13% at 50 K, and 9% at 150 K (∼5% at 300 K) with a magnetic field of 4 kOe applied parallel to the current. In contrast, the high-field magnetoresistance reveals the opposite trend. This can be clearly seen in Figure 3f, showing the temperature-dependent negative HFMR effect.42 The electrical resistance measured with and without applied field is shown as well. In both cases, the electrical resistance increases with decreasing temperature until a maximum is reached at the insulator-to-metal transition at TIM ≈ 180 K. As the temperature is lowered further, the electrical resistance decreases until T = 44 and 52 K for H = 0 and 45 kOe, respectively. Thereafter, it increases again (the upturn in resistance is less pronounced when a magnetic field is applied), in agreement with literature reports.42−44 In some of these works, this behavior is related to electron−electron interaction or spin-disorder scattering at low temperatures, but the underlying mechanisms remain elusive. As is evident from the data in Figure 3f, the negative HFMR effect (up to −53% at T = 215 K) does not change much over a wide temperature range. This behavior is different from epitaxially grown layers, but it has been reported for polycrystalline LCMO nanomaterials.42 Overall, the obtained values for the LFMR effect appear to be similar or slightly larger than those observed for other perovskite-type mixed-valence manganese oxides fabricated by chemical solution deposition. This is likely due to the large number of interfaces associated with both the small grains and the mesoporous morphology and the resulting complex percolation pathways for electron transport. However, we
Figure 3. Magnetic and magnetotransport properties of the PIB107-bPEO150-templated LCMO thin films heated at 750 °C. (a) ZFC/FC curves measured in an applied field of 100 Oe. (b) Hysteresis loops at 5, 100, and 200 K. Field-dependent MR effect at 5 K (c), 50 K (d), and 150 K (e). Red curves show the negative LFMR effect in units of percent, and black curves represent the corresponding hysteresis loops. (f) Temperature-dependent negative HFMR effect.
curve measured at 250 K follows the Langevin function (see Figure S4 of the Supporting Information); the magnetic cluster diameter of (13 ± 1) nm was extracted from the fit to the data. This dimension is in fair agreement with the average crystallite size. As mentioned previously, the magnetotransport properties of the PIB107-b-PEO150-templated LCMO thin films were also investigated in this work. Both the negative low-field (LFMR) and high-field magnetoresistive (HFMR) effects were determined by measuring the electrical resistance (R) as a function of temperature and applied field. The negative magnetoresistance is defined as follows: MR(%) = [(R(H) − R(0))/ R(H)]·100. Panels c−e of Figure 3 show the LFMR effect in the field range of ±4 kOe at different temperatures (see also Figure S5 of the Supporting Information for field-dependent MR up to 45 kOe) along with the magnetic hysteresis loops. At HC, the films have the highest electrical resistance and thus reveal the lowest absolute value for the MR effect, in agreement with theory. This is due to the low relative orientation of the spin-polarized charge carriers at the Fermi energy; strong 5749
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note that a direct comparison with magnetoresistive materials that have been described over the years is somewhat difficult because of different key parameters, such as measurement temperature and applied field. Since the pore structure provides the films with an active surface area that is much larger than the geometric area, the sol−gel-derived LCMO studied in this work represents an ideal model system to find out if, and to what extent, the magnetization can be tuned by electrostatic modulation of the surface charge carrier density.16,17 All measurements were carried out in situ inside the SQUID magnetometer at T = 230 K using a custom-built sample tube (see Figure S6 of the Supporting Information) with conventional two-electrode setup, with the PIB107-b-PEO150-templated LCMO thin film and high surface area carbon fiber cloth serving as the working and counter electrodes, respectively. Diethylmethyl(2-methoxyethyl)ammonium bis(trifluoromethylsulfonyl)imide was chosen as nonaqueous electrolyte. Part of the reason for this is the low freezing point of this particular ionic liquid (the ions are mobile down to 200 K according to Lee et al.) and the large operating potential range.45,46 As a first set of experiments, the mesostructured material was charged potentiostatically at either +1.2 or −1.2 V in the presence of a magnetic field of 4 kOe. When the magnetization had leveled off, hysteresis curves were measured in the field range of ±4 kOe, while keeping the potential constant (see Figure S7 of the Supporting Information). The data analysis reveals an increase in magnetization by 3.0% upon positive polarization and a decrease by 3.7% upon negative polarization. This behavior can be explained by electrostatic hole- and electron-doping, respectively, which causes an increase/decrease in the Mn4+ to Mn3+ ratio. Assuming that electrostatic carrier doping is, in first approximation, equivalent to chemical doping and considering the magnetic phase diagram of LCMO, a change in the aforementioned ratio results in a shift in the magnetic transition temperature TC (a change in magnetic anisotropy can be ruled out). Therefore, for isothermal measurements, the magnetic moment must either increase or decrease, depending on the polarization state. To gain more insight into the surface charge-induced effect, the magnetization change was monitored during voltammetric cycling experiments. The cutoff voltages were set at either ±1.2 or ±1.4 V. Representative voltammetry sweeps at a rate of 2 mV/s are shown in Figure 5a. As can be seen, the voltammograms are featureless and do not change during cycling; the rectangular-shaped profiles are characteristic of electrochemical double-layer capacitors. Panels b and d of Figure 5 show the relative change in magnetization at an applied field of 500 Oe along with the charge as a function of time. Both curves follow the very same general shape, that is, they reach their maximum values at the same time and are perfectly in-phase, which demonstrates that the overall process is fully reversible. This can also be seen from the data in Figure 5c: the surface charge/relative magnetization change relationship is rather linear (yet, the positive and negative branches have slightly different slopes). The small opening of the curve is almost certainly due to minor deleterious side reactions at the working electrode. However, kinetic effects might play a role as well. The fact that the magnetization change for a given amount of charge is slightly different for the positive and negative polarization states can again be explained on the basis of the magnetic phase diagram. TC does not change linearly with Ca2+
Figure 5. Reversible electrostatic carrier doping in the PIB107-bPEO150-templated LCMO thin films at T = 230 K. (a) Voltammetry sweeps at a rate of 2 mV/s. (b, d) Magnetization change (normalized to the magnetization at 0 V) upon capacitive charging measured in an applied field of 500 Oe. (c) Relative change in magnetization as a function of surface charge.
concentration around the range of interest; the relative change in magnetization should be higher upon negative polarization (electron-doping), which is exactly what we find for the mesoporous LCMO thin films.47 Lastly, the same carrier doping experiments were carried out in the reversed applied field of −500 Oe (see Figure S8 of the Supporting Information). As expected, the magnetic response and charge curves are not coincident any more, but antiphase. However, the relative change in magnetization follows perfectly the charge curve and is very similar to that in the positive applied field direction, thereby indicating that current-induced magnetic moments (e.g., due to leakage current or other undesired effects) are negligible, if even existent. Overall, the peak-to-peak magnetization modulation equals 5.9 and 8.5% in the potential windows of +1.2 to −1.2 V and +1.4 to −1.4 V, respectively. These values are, to our knowledge, the highest thus far reported for electrolyte-gated perovskite-type mixedvalence manganese oxide thin films and emphasize the benefits of nanoscale porosity.
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CONCLUSIONS In summary, we have shown that continuous large-pore mesoporous LCMO thin films can be prepared through a facile soft-templating method. This novel solution-processed mesostructured material exhibits interesting magnetotransport properties and further allows for large tuning of the magnetization upon capacitive charging due to the unique morphology. The synthetic approach employed in this work is of a more general nature in that it is extendable to other oxides with the same general formula, including lanthanum strontium manganese oxide. Overall, the findings are very interesting from a technological point of view, as they may lead to new design paradigms for low-field magnetoresistive materials and magnetically tunable nanostructures. 5750
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ASSOCIATED CONTENT
S Supporting Information *
Additional data from electron microscopy, XPS, magnetic/ magnetoresistive measurements, and electrostatic carrier doping experiments. This material is available free of charge via the Internet at http://pubs.acs.org.
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. Funding
Research at the Institute of Nanotechnology was supported by the German Research Foundation (T.B., grant no. BR 3499/31; H.H., grant no. HH 1344/28-1), the Fonds der Chemischen Industrie im Verband der Chemischen Industrie (C.R.), and the Helmholtz Association. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS We thank Ralf Witte, Robert Kruk, and Thomas Leichtweiss for their assistance with GIXRD, SQUID magnetometry, and XPS, respectively.
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