Room-Temperature Ferromagnetism in Antiferromagnetic Cobalt

Letter pubs.acs.org/NanoLett

Room-Temperature Ferromagnetism in Antiferromagnetic Cobalt Oxide Nanooctahedra Nerio Fontaíña-Troitiño,† Sara Liébana-Viñas,†,‡ Benito Rodríguez-González,† Zi-An Li,‡ Marina Spasova,‡ Michael Farle,‡ and Verónica Salgueiriño*,† †

Departamento de Física Aplicada, Universidade de Vigo, 36310, Vigo, Spain Fakultät fur Physik and Center for Nanointegration (CENIDE), Universität Duisburg-Essen, 47057, Duisburg, Germany

‡

S Supporting Information *

ABSTRACT: Cobalt oxide octahedra were synthesized by thermal decomposition. Each octahedron-shaped nanoparticle consists of an antiferromagnetic CoO core enclosed by eight {111} facets interfaced to a thin (∼4 nm) surface layer of strained Co3O4. The nearly perfectly octahedral shaped particles with 20, 40, and 85 nm edge length show a weak room-temperature ferromagnetism that can be attributed to ferromagnetic correlations appearing due to strained lattice configurations at the CoO/Co3O4 interface. KEYWORDS: Ferromagnetism, antiferromagnetism, strain, anisotropy

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Nanocrystals with cubic, octahedral, and rhombic dodecahedron geometries are particularly important because they are selectively terminated by {100}, {111}, and {110} facets on the surface. The atoms at these surfaces have lower coordination resulting in sometimes dramatically different surface magnetic properties than for the corresponding bulk material. Similarly, oxidation and the presence of surfactants have a strong bearing on the magnetic response of the nanoparticles (see, for example, refs 14 and 15). The magnetic surface anisotropy resulting from the lower symmetry at the surface can favor in plane and out plane orientation of the magnetic moments.16−18 As a consequence of the interplay between surface and bulk magnetic anisotropies and exchange interactions, however, noncollinearities between surface and volume magnetic moments can appear and be strongly modified by the presence of edges and different crystallographic surface orientations.18 Accordingly, the magnetism can vary with the local chemical composition and crystalline structure in the particle, for example, due to constrained surface facets19 causing local structural and electronic perturbations.20−22 Herein, we report on the synthesis and structural as well as magnetic characterization of nominally AFM CoO-Co3O4 octahedra, which display only the eight {111} surface facets of the cubic structure. We find a ferromagnetic response at room temperature not arising from spurious metallic Co clusters but from the particular morphology attained by the AFM nanocrystals.

odern spintronic devices such as magnetic tunnel junctions and high-density memories take advantage of antiferromagnetic (AFM) materials at the nanoscale, for example, to modify the switching behavior of adjacent ferromagnets (FM) via the exchange bias effect,1 to adjust their coupling for magneto-logic devices,2 or to tailor their interactions in dense nonvolatile storage of information.3 Consequently, the understanding and tuning of spin structures at AFM surfaces4 and interfaces is highly desirable not only for these advancing technologies but also for enabling new insights into the fundamental understanding of the related magnetic effects.5−8 Well-defined polyhedron-shaped nanocrystals have attracted intense attention as prototype systems to investigate facetdependent properties.9−12 The methods for the synthesis of these nanocrystals in solution should provide the constituent atoms in such a way that they can rearrange and grow along energetically strongly favored crystalline directions to provide well-defined surface facets. This rearrangement depends on the cohesive energy per atom and correlates with the melting temperature of the solid that becomes therefore a decisive factor in determining optimal conditions for nanocrystal growth, given predominantly by the boiling temperature of the solvent in which the colloidal synthesis proceeds.13 In this regard, the lower melting temperature of nanomaterials becomes a favoring factor in the synthesis of these nanocrystals, which can be grown at temperatures where organic surfactant molecules are stable (200−400 °C) and can adhere to the growing crystals. The organic−inorganic interface determines the crystalline growth directions of the nanoparticles by means of a dynamic solvation of these surfactants in solution, which generally contain electron-donating groups to allow coordination to electron-poor metal atoms at the nanocrystal surface. © XXXX American Chemical Society

Received: October 15, 2013 Revised: January 13, 2014

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dx.doi.org/10.1021/nl4038533 | Nano Lett. XXXX, XXX, XXX−XXX

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Figure 1. TEM images of the octahedron-shaped nanocrystals with 20 (a), 40 (b), and 85 (c) nm average edge length. X-ray photoelectron (XPS) spectra (d−f) and XRD pattern (g). EEL spectra (h) at the O-K and Co-L edges recorded in the core (spot 1) and shell (spot 2) regions of the octahedron as indicated in the TEM image (i).

(S2) with respect to the Co 2p3/2 main peak. They show relative intensities (S2 is rather weak but still discernible) that indicate that Co2+ is the predominant species and some Co3+ is also present. Figure 1f shows the O 1s signals, centered at 531.7 eV. As the Co 2p and O 1s core levels binding energies of cobalt oxides and hydroxides are in the same range, the spectra confirm again the absence of metallic cobalt.23 Figure 1g shows a representative XRD pattern of the octahedron-shaped nanocrystals with peaks matching the Bragg reflections of the standard face-centered cubic (fcc) structure, which corresponds to cubic CoO. No other phases such as the hexagonal CoO wurzite or the Co3O4 spinel types are discernible. It is worth noting that the ratio of intensities of the (200) and (111) diffraction peaks is higher than the tabulated value (1.36 instead of 1.19), which indicates that the nanostructures are abundant in {100} planes, likely due to an accelerated growth along the ⟨100⟩ directions.24 Spatially resolved EELS of the octahedra were collected in scanning transmission electron microscopy (STEM) mode. Figure 1h shows the EELS spectra at the O K-edges and Co L2,3-edges after background subtraction for two positions marked as spot 1 (1 nm probe size) and spot 2 in the TEM image of the octahedron included in Figure 1i. At the O K-edge, the peaks labeled as a, b, and c arise from the electron transition from O

Figure 1 includes transmission electron microscopy (TEM) images of the as-synthesized cobalt oxide nanocrystals. Because TEM gives the two-dimensional (2D) projection of the objects, the octahedron-shape particles look like hexagons, rectangles, or rhombi. For an overview of the size and polydispersity, we perform statistical measurements (log−normal) considering the edge length of the nanocrystals from three samples (see Figure S1 in the Supporting Information) with average edge length of 19.40 nm*/1.33, (contains 95.5% of the nanocrystals), 40.95 nm*/1.56, (95.5%) and 84.43 nm*/1.51, (95.5%) (referred to in the following as 20, 40, and 85 nm samples). The crystalline structure and chemical composition was characterized by X-ray photoelectron spectroscopy (XPS) (Figure 1d−f), X-ray diffraction (XRD) (g), and electron energy loss spectroscopy (EELS) (h,i). The XPS spectra indicate the presence of cobalt and oxygen as expected. The Co 2p region (Figure 1e) includes two broad sets of signals corresponding to 2p3/2 (782 eV) and 2p1/2 (798 eV) with many similarities to CoO and Co(OH)2 confirming the presence of Co2+ while excluding the presence of metallic Co (778.0 eV) (within the resolution of XPS). The cobalt oxidation states are highly related to the energy gaps between the Co 2p main peak and satellite peaks. The spectra included in Figure 1e include two distinguishable satellite peaks shifted by ∼6 (S1) and ∼9 eV B


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Figure 2. HRTEM image (scale bar: 2 nm) of an octahedron showing core and shell areas with different contrast (a). Fourier filtering of this TEM image using frequencies that correspond only to Co3O4 planes (b) and using frequencies that correspond both to Co3O4 and to CoO (c). Insets demonstrate the Fourier transform of each of the images.

Figure 3. HRTEM images of the cobalt oxide octahedra showing the planes ordered up to the surface in the [211] and [110] zone axis respectively, including the views of a {111} facet in section (a,b). Three different projections of the octahedra in the [011], [112], and [111] zone axis (sketches in yellow correspond to projection schemes of an octahedron model in the same corresponding axis (c−e). Corresponding SAED patterns obtained from the same octahedron shown in the upper images. In plane rotation of the diffraction patterns were compensated (f−h).

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1s → 2p states hybridized with Co unoccupied 3d, 4s, and 4p orbitals, respectively. Relative intensities and energies of the peaks a, b, and c allow to discriminate Co-oxide phases (CoO or Co3O4). In O K-edge data, the high intensity of the prepeak a recorded in the shell region (spot 2) is indicative of Co3O4, while in the one from spot 1 in which the intensity of the prepeak a is low is typical for the CoO phase. The white line ratio of the Co L3 to L2 edge intensities in the two regions is 4.32 and 4.35, respectively. In comparison to the bulk reference ratios, 4.51 for CoO and 2.42 for Co3O4, this would suggest that the atomic Co electronic structure in the core and shell region is the same and similar to CoO (Co2+ state). Because the CoO presence in the shell can be unambiguously excluded from high-resolution transmission electron microscopy (HRTEM) and XPS results, we attribute this unexpected Co L3/L2 ratio in the shell region to a structurally distorted, nonstoichiometric, and/or defective phase Co3O4-like phase. This could explain the high intensity of prepeak a in O K-edge and the satellite peak S2 in the XPS Co 2p spectra, but approximately the same L3/L2 ratio for the core and shell regions. Figure 2 shows a higher magnification image of an octahedron that exhibits bright-dark patched contrast. Because the particles present a nearly perfect octahedron shape and are oriented along their zone axes, the contrast variation in the shell arises from changes in scattering condition due to the core/ shell structural overlay. The existence of the shell with a different crystalline lattice is demonstrated by the Fourier filtered images shown in Figure 2b,c. In Figure 2b, we show an image that is the result of the Fourier filtering of the TEM image in Figure 2a, using the spots or frequencies that are exclusive for the Co3O4; while in Figure 2c the filtering was done using the spots that are due both to Co3O4 and CoO. This provides convincing evidence for the presence of two different lattices; CoO in the core region and Co3O4 spinel-like phase in the shell. No planar defects such as twinning or stacking faults are detectable. The hypothesis of local strain is also supported in these atomically resolved images on which the black patches display better the crystalline lattice than the bright ones. Figure 3a,b includes HRTEM images with enlarged views of the octahedron edges (the regions marked by the red squares in the insets show the imaged area on the octahedron). These HRTEM images show well-defined lattice fringes with the interplanar spacing of the (111) planes resolved at 2.46 Å. These images also show really smooth {111} facets of the cubic structure, demonstrating planar boundaries ordered up to the surface. The flatness of the limiting surfaces of the octahedra is outstanding and is an indication of the excellent crystallinity of the nanoparticles. Figure 3c−e includes different perspectives of the octahedra and their respective Fourier transform analysis (f−h) to obtain the corresponding crystalline zone axis. These projections look like a rhombus, a rectangle, and a hexagon and correspond to the [011], [112], and [111] zone axis, respectively. Figure 3c,d depicts the position of the {111} facet now parallel to the electron beam, which allows us to appreciate the facets of the octahedra. The TEM images permit also the verification of the presence of bevels in some corners of the nanooctahedra (red arrows in Figure 3c−e). The selected area electron diffraction (SAED) patterns obtained from single nanocrystals (Figure 3f−h) show a set of spots that can be indexed on the basis of a single crystal CoO zone axis and no other phases are detectable. This result confirms again that the octahedra are single crystalline.

Figure 4 shows the results from three-dimensional (3D) electron tomography.25 Figure 4a shows a pair of projected 2D

Figure 4. HRTEM (a, left) and 3D reconstruction images (a, right, and b−d) of the octahedra. Black arrows indicate the presence of voids inside the octahedra.

and reconstructed 3D images of a chain of octahedra. The reconstructed 3D structure that is volume-rendered enables us to measure the edges and angles (70.1 and 117° dihedral angles between (101) and (110) and between (111) and (110) facets) of the individual octahedron-shaped particles. The analysis reveals the almost ideal octahedron shape with equal edge lengths. An extracted octahedron is color coded and oriented along major axes as shown in Figure 4b,c together with octahedron schematics. It is noteworthy that according to the perfect octahedron shape, exact {111} crystallographic facets terminated on the eight facets of the octahedron emerge. In comparison to previous work by Park and co-workers who synthesized hexagonal and cubic cobalt oxide nanoparticles and nanoparallelepipeds from a similar reaction (cobalt acetylacetonate in benzylamine or oleylamine,26−28 instead of cobalt acetate in trioctylamine), our method is different in terms of (i) the nucleophilic attack between the acetate and the amine groups, (ii) the growth of the nanoparticles (using oleic acid), and (iii) the solvent that reaches a much higher refluxing temperature. All three aspects favor the production of perfectly shaped octahedra. Oleic acid as capping agent controls the formation of densely packed {111} facets,29 because in its absence no octahedra were produced. For an understanding, one has to consider the surface energies and polarizabilities of different crystallographic facets of ionic crystals.30,31 The capping molecules with negatively charged head groups can selectively stabilize the (111) planes (containing Co2+ (or Co3+) cations only) because electrostatically the interaction with the charged {111} facets is favored in comparison to the uncharged {100} facets (containing Co and O atoms). The polar (111) surface is energetically not stable. Goniakowski et al. pointed out that nature tends to avoid such a polarity catastrophe by changing the distribution of surface charges or D


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Figure 5. (a) Magnetization as a function of the applied field for the 20, 40, and 85 nm samples recorded at 300 K. (b) The same M(H) curves in a low-field range showing the presence of ferromagnetic hysteresis for all three samples.

magnetic field recorded at 300 K, well above the magnetic ordering temperatures for both cobalt oxides (Néel temperature of CoO TN1 = 291 K and of Co3O4 TN2 = 40 K). The shape of the M(H) curves clearly indicates the presence of different magnetic phases. The linear increase of the magnetization with increasing field above 1 T is a response of paramagnetic and also antiferromagnetic phases. Surprisingly, all three samples show a well-defined hysteresis at 300 K (see Figure 5b). Coercive fields μ0HC = 370, 290, and 320 mT are measured for 20, 40, and 85 nm samples, respectively. The intrinsic magnetization was determined by extrapolation of the linear part of the M(H) curve to H = 0. Table 1 summarizes the

the modification of the composition of the surface region.32 In our case, surface stabilization seems to be attained by the formation of the spinel Co3O4 configuration, which can be seen as four layers of a hypothetical spinel (with three out of four top oxygen atoms missing) on top of CoO. Similarly to an Oterminated octopolar configuration, the spinel-like termination is characterized by three-fourths of the oxygens missing in the surface layer and one-fourth of the cations missing in the subsurface layer and thus fulfills the electrostatic compensation condition. A chemical stabilization of the surface facets by OH groups can also be considered. Figure 4d shows the oblique slicing along an octahedron [001] zone-axis, which reveals some voids inside the octahedron (indicated by black arrows), also appreciated in some TEM images. The presence of holes inside the octahedra and located around a main and centered cavity (5−10 nm in diameter) can be explained considering the synthetic process as kinetically controlled (nucleation at T ∼ 170 °C). In that case, an organometallic intermediate phase would match the premises for creating some kind of porosity using a chemical gradient, because according to Smigelkas and Kirkendall solid diffusion in a concentration gradient occurs through a vacancy exchange mechanism.33 If the diffusion coefficient of the two species (the intermediate organometallic compound and the CoO) is different, the net directional flow of vacancies results in the formation of pores with the outward diffusion of the core material (the intermediate organometallic compound) through the CoO being faster than the inward diffusion of the outer material (the CoO). Gösele and co-workers claimed a general fabrication route for hollow nanostructures provided that the Kirkendall effect should be generic,34,35 and Dilger et al. reported a similar case for the synthesis of aerogel-like ZnO with organometallic methylzinc methoxyethoxide ([MeZnOEtOMe]4) and ZnO.36 The magnetic characterization of dried powders fixed in a capsule was performed using a superconducting quantum interference device (SQUID) magnetometer. Initial tests had shown that these nominally AFM nanoparticles moved in magnetic field gradients yielding a first indication of a ferromagnetic response. The measured magnetic response of the powder is determined by a complex interplay between the distribution of magnetic properties of individual octahedra and their interactions. Figure 5 shows the magnetization of the 20, 40, and 85 nm octahedral particles as a function of the applied

Table 1. High-Field Magnetic Mass Susceptibility χ0, Ferromagnetic Saturation Magnetisation Msat, coercivity μ0Hc, and Ferromagnetic Moment Per One Particle μp Determined from Magnetization Measurements at 300 Ka

a

sample

χ0 [*10−7 m3/kg]

FM Msat [A* m2/kg]

μ0Hc [mT]

FM moment per particle [μB]

20 nm 40 nm 85 nm

8.87 8.6 8.3

0.165 2.4 1.58

370 290 320

430 49970 384276

The error bar for the data is 10%.

result of magnetization measurements at 300 K. With increasing size, the mass susceptibility is the same within the error bar indicating the size independent contribution of an AFM and/or paramagnetic phase. The saturation magnetization Msat as well as the coercive field shows a nonuniform behavior though the total magnetic moment per particle increases with size. The nonmonotonous values of coercivity as increasing size can stem from the different contributions of magnetic anisotropy present in these nanocrystals, such as surface anisotropy (changing monotonously) and magnetostriction effects (changing nonmonotonously), which vary with the octahedron average edge. With increasing size, the configuration of magnetic domains as well as the domain reversal mechanism can change in the nanocrystals. Given the fact that we have an antiferromagnet in contact with a ferromagnetic shell constrained by the octahedron shape, the domain configuration and the reversal mechanism, which determine the effective HC, cannot be predicted by simple arguments. The nonmonotonic behavior indicates, however, that the configurations and mechanisms change as a function of size. An initial E


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Figure 6. Temperature dependence of ZFC-FC magnetization of the 20 nm (left) and 40 nm (right) samples. The measuring and cooling field is 10 mT.

model based on the formation of a (surface/interface)-induced ferromagnetism would predict a linear increase of the volume magnetization with the surface or interface area. However, although the total magnetic moment per particle increases, the mass normalized magnetization does not. One possible reason might be that the strain effects, which stabilize the ferromagnetic order, are only present up to a certain lateral size of the (111) facets. Strain may not be maintained for larger interface areas and misfit dislocations may destroy a ferromagnetic long-range order. Such a model would also explain why ferromagnetic order has not been discussed for CoO/Co3O4 (111) surfaces, for example, in thin films. Additionally, we like to emphasize again that based on the chemical synthesis route and on preliminary data from elementspecific X-ray spectroscopy we can exclude the presence of Co0 (metallic Cobalt) as the origin for the observed ferromagnetism. Another clear evidence for the room temperature ferromagnetic response of the samples is provided by the zero-fieldcooled (ZFC) and the field-cooled (FC) magnetization as a function of temperature from 5 up to 370 K for the 20 and 40 nm samples, as shown in Figure 6. For the ZFC measurements, the sample was cooled without an external magnetic field from 370 to 5 K, then a magnetic field of 10 mT was applied and the magnetization of the sample was measured with increasing temperature. For FC measurements, the sample was cooled in a magnetic field of 10 mT from 370 to 5 K, and then the magnetization was measured at 10 mT with increasing temperature. A distinct difference between ZFC and FC magnetization is observed in the temperature range from 5 up to 370 K, characteristic of a ferromagnetic phase present in the system. No evidence for superparamagnetism such as a maximum in the temperature dependence of the magnetization is observed when measured in a low field of 10 mT. Nevertheless, the temperature dependence of the high-field susceptibility (Figure 7), determined by the slope of M(H) measurements in fields between 2 and 3T shows a characteristic maximum at T = 285 K, close to the Néel temperature of bulk CoO for all three samples. No obvious maxima near the Néel temperature of Co3O4 are observed. These temperaturedependent magnetic data confirm the presence of the AFM CoO. In summary, octahedron-shaped CoO nanoparticles with different tunable sizes were synthesized by thermal decom-

Figure 7. Temperature dependence of the high-field magnetic mass susceptibility of the 20, 40, and 85 nm samples of octahedron-shape cobalt oxide nanoparticles.

position. High-resolution chemical and structural analyses confirm that the eight polar {111} facets of the octahedral are stabilized by the formation of a strained Co3O4 shell. The particles show a ferromagnetic response well above the antiferromagnetic ordering temperature of CoO, which can be attributed to the formation of an ∼4 nm thick strained surface layer of Co3O4. The lateral size of the facets, that is, the surface area, appears to be a critical parameter for the stabilization of ferromagnetic order in otherwise antiferromagnetic CoO@Co3O4 nanoparticles. Experimental Section. Chemicals. Cobalt(II) acetate tetrahydrate (Sigma-Aldrich, 98%), trioctylamine (TOA) (Aldrich, 90%), oleic acid (OA) (Aldrich, 90%), oleylamine (Aldrich, 70%), trioctyl phosphine oxide (TOPO) (Aldrich), octyl ether (Aldrich, 97%) and ethanol anhydrous. Synthesis. 40 and 75 nm (average edge length) CoO nanooctahedra were obtained using the hot-injection thermodecomposition of cobalt acetate (2.66 and 4 mmol) (dissolved in 5 mL of ethanol) in trioctylamine (25 mL, 57.18 mmol) and oleic acid (5.12 and 8 mmol) at 170 °C and then left to reflux (at T = 300 °C) for two hours. 20 nm (average edge length) CoO nanooctahedra were similarly synthesized dissolving cobalt acetate (2 mmol), oleic acid (8 mmol), oleylamine (20 F


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support through EU Network “REFreePerMag”. The Xunta de Galicia (Regional Government, Spain) has supported this work under projects 10PXIB312260PR and 2010/78 (Modalidade Emerxentes).

mmol) and TOPO (0.2 mmol) in a mixture of trioctylamine (5 mL) and octyl ether (15 mL), and then heating to reflux (at T = 300 °C) for six hours. Once cooled to room temperature, the nanostructures were separated by centrifugation, washed several times, and finally stored in ethanol. Characterization. TEM, HRTEM, and SAED were used to investigate particle size, morphology, shape, and crystalline microstructure of the octahedron-shaped CoO nanocrystals. HRTEM analyses were carried out on a JEOL JEM2010F TEM with a field emission gun working at 200 kV. Samples for TEM and HRTEM observations were prepared by dropping a diluted suspension of the particles onto a standard carbon coated copper grid. A second TEM (FEI, Tecnai F20) with a Schottkytype field emission gun operating at 200 kV and equipped with Gatan Imaging Filter (GIF) system by which EELS can be performed was also used. EELS were acquired in a STEM mode with a minute electron probe of around 1 nm. The core level EEL spectra of O K-edge (electron transition from O 1s → 2p orbitals) and Co L2,3- edge (electron transition from Co 2p → 3d) assist to determine the chemical states of the cobalt atoms in oxides. The energy resolution in EELS was measured to be around 1 eV or better, as evaluated from the full-width at halfmaximum of the zero-loss peak. Spectra were recorded at dispersion of 0.3 eV/channel and the convergence and collection semiangles are estimated to be 10 and 11.2 mrad, respectively. All the EEL spectra were corrected for dark current and gain variation of CCD camera. X-ray diffraction patterns were collected using a Panalytical X’Pert Pro diffractometer (Cu Kα radiation (λ = 1.54056 Å) in the 5− 100° 2θ range) and compared with crystallographic information files (CIF) from the crystallographic open database (COD). XPS surface measurements (Co2p and O1s spectra) were recorded using a Thermo Scientific K-Alpha instrument. Monochromatic X-ray source Al Kα (1486.6 eV) was used for all samples and experiments. The calibration of the binding energies (BEs) was done using the carbon 1s peak at 285.0 eV for spurious carbon on unsputtered surfaces as reference. The atomic concentrations were determined from the XPS peak areas using the Shirley background subtraction technique and the Scofield sensitivity factors. Magnetic measurements were performed using SQUID magnetometry using powder samples. The diamagnetic signal from the capsules holding the powder samples was subtracted.

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ASSOCIATED CONTENT

S Supporting Information *

Figure S1 includes the graphs showing the size distribution analysis and the log−normal fit parameters, considering 208, 358, and 418 nanoparticles in the 20, 40, and 85 nm samples, respectively. This material is available free of charge via the Internet at http://pubs.acs.org.

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REFERENCES

(1) Parkin, S. S. P.; Roche, K. P.; Samant, M. G.; Rice, P. M.; Beyers, R. B.; Scheuerlein, R. E.; O’Sullivan, E. J.; Brown, S. L.; Bucchigano, J.; Abraham, D. W.; Lu, Y.; Rooks, M.; Trouilloud, P. L.; Wanner, R. A.; Gallagher, W. J. J. Appl. Phys. 1999, 85, 5828−5833. (2) Wolf, S. A.; Awschalom, D. D.; Buhrman, R. A.; Daughton, J. M.; von Molnar, S.; Roukes, M. L.; Chtchelkanova, A. Y.; Treger, D. M. Science 2001, 294, 1488−1495. (3) Loth, S.; Baumann, S.; Lutz, C. P.; Eigler, D. M.; Heinrich, A. J. Science 2012, 335, 196−199. (4) Mariño-Fernández, R.; Masunaga, S. H.; Fontaíña-Troitiño, N.; Morales, M. P.; Rivas, J.; Salgueiriño, V. J. Phys. Chem. C 2011, 115, 13991−13999. (5) Sugiyama, I.; Shibata, N.; Wang, Z. C.; Kobayashi, S.; Yamamoto, T.; Ikuhara, Y. Nat. Nanotechnol. 2013, 8, 266−270. (6) Wang, K.; Smith, A. R. Nano Lett. 2012, 12, 5443−5447. (7) Bode, M.; Vedmedenko, E. Y.; von Bergmann, K.; Kubetzka, A.; Ferriani, P.; Heinze, S.; Wiesendanger, R. Nat. Mater. 2006, 5, 477− 481. (8) Kuch, W.; Chelaru, L. I.; Offi, F.; Wang, J.; Kotsugi, M.; Kirschner, J. Nat. Mater. 2006, 5, 128−133. (9) Tao, A. R.; Habas, S.; Yang, P. Small 2008, 4, 310−325. (10) Lim, B.; Jiang, M.; Tao, J.; Camargo, P. H. C.; Zhu, Y.; Xia, Y. Adv. Funct. Mater. 2009, 19, 189−200. (11) Xia, Y.; Xiong, Y.; Lim, B.; Skrabalak, S. E. Angew. Chem., Int. Ed. 2009, 48, 60−103. (12) Salazar-Á lvarez, G.; Qin, J.; Sepelak, V.; Bergmann, I.; Vasilakaki, M.; Trohidou, K. N.; Ardisson, J. D.; Macedo, W. A. A.; Mikhaylova, M.; Muhammed, M.; Baró, M. D.; Nogués, J. J. Am. Chem. Soc. 2008, 130, 13234−13239. (13) Yin, Y.; Alivisatos, A. P. Nature 2005, 437, 664−670. (14) Jamet, M.; Wernsdorfer, W.; Thirion, C.; Dupuis, V.; Melinon, P.; Perez, A.; Mailly, D. Phys. Rev. B 2004, 69, 024401. (15) Salafranca, J.; Gazquez, J.; Pérez, N.; Labarta, A.; Pantelides, S. T.; Pennycook, S. J.; Batlle, X.; Varela, M. Nano Lett. 2012, 12, 2499− 2503. (16) Néel, L. J. Phys. Radium 1954, 15, 376−378. (17) Garanin, D. A.; Kachkachi, H. Phys. Rev. Lett. 2003, 90, 065504. (18) Yanes, R.; Chubykalo-Fesenko, O.; Kachkachi, H.; Garanin, D. A.; Evans, R.; Chantrell, R. W. Phys. Rev. B 2007, 76, 064416. (19) Mamatkulov, M.; Filhol, J.-S. J. Phys. Chem. C 2013, 117, 2334− 2343. (20) Sander, D.; Kirschner, J. Appl. Phys. A 2007, 87, 419−425. (21) Choi, W. S.; Kwon, J.-H.; Jeen, H.; Hamann-Borrero, J. E.; Radi, A.; Macke, S.; Sutarto, R.; He, F.; Sawatzky, G. A.; Hinkov, V.; Kim, M.; Lee, H. N. Nano Lett. 2012, 12, 4966−4970. (22) Rondinelli, J. M.; Spaldin, N. A. Adv. Mater. 2011, 23, 3363− 3381. (23) Yang, J.; Liu, H.; Martens, W. N.; Frost, R. L. J. Phys. Chem. C 2010, 114, 111−119. (24) Wang, N.; Cao, X.; Guo, L.; Yang, S.; Wu, Z. ACS Nano 2008, 2, 184−190. (25) Midgley, P. A.; Weyland, M. Ultramicroscopy 2003, 96, 413− 431. (26) Seo, W. S.; Shim, J. H.; Oh, S. J.; Lee, E. K.; Hur, N. H.; Park, J. T. J. Am. Chem. Soc. 2005, 127, 6188−6189. (27) Nam, K. M.; Shim, J. H.; Ki, H.; Choi, S.-I.; Lee, G.; Jang, J. K.; Jo, Y.; Jung, M.-H.; Song, H.; Park, J. T. Angew. Chem., Int. Ed. 2008, 47, 9504−9508. (28) Shim, J. H.; Nam, K. M.; Seo, W. S.; Song, H.; Park, J. T. Chem. Asian J. 2011, 6, 1575−1581. (29) Pimpinelli, A.; Villain, J. Physics of Crystal Growth; Cambridge University Press: Cambridge, U.K., 1998.

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS V.S. acknowledges the financial support from the Ramón y Cajal (Ministerio de Ciencia e Innovación, Spain) Program. M.F., S.L.-V., Z.-A.L. and M.S. acknowledge the financial G


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(30) Wang, Z. L. J. Phys. Chem. B 2000, 104, 1153−1175. (31) Zhang, J. M.; Ma, F.; Xu, K. W. Appl. Surf. Sci. 2004, 229, 34− 42. (32) Goniakowski, J.; Finocchi, F.; Noguera, C. Rep. Prog. Phys. 2008, 71, 016501. (33) Smigelskas, A. D.; Kirkendall, E. O. Trans. Am. Inst. Min., Metall. Eng. 1947, 171, 130−142. (34) Fan, H. J.; Knez, M.; Scholz, R.; Nielsch, K.; Pippel, E.; Hesse, D.; Zacharias, M.; Gosele, U. Nat. Mater. 2006, 5, 627−631. (35) Fan, H. J.; Knez, M.; Scholz, R.; Hesse, D.; Nielsch, K.; Zacharias, M.; Gosele, U. Nano Lett. 2007, 7, 993−997. (36) Dilger, S.; Lizandara-Pueyo, C.; Krumm, M.; Polarz, S. Adv. Mater. 2012, 24, 543−548.

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Room-Temperature Ferromagnetism in Antiferromagnetic Cobalt

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