Structure-Correlated Exchange Anisotropy in Oxidized Co80Ni20

Article pubs.acs.org/cm

Structure-Correlated Exchange Anisotropy in Oxidized Co80Ni20 Nanorods Sara Liébana-Viñas,†,‡ Ulf Wiedwald,*,† Anna Elsukova,† Juliane Perl,† Benjamin Zingsem,† Anna S. Semisalova,†,§,∥ Verónica Salgueiriño,‡ Marina Spasova,† and Michael Farle† †

Faculty of Physics and Center for Nanointegration (CENIDE), University of Duisburg-Essen, 47057 Duisburg, Germany Departamento de Física Aplicada, Universidade de Vigo, 36310 Vigo, Spain § Faculty of Physics, Lomonosov Moscow State University, 199911 Moscow, Russia ‡

S Supporting Information *

ABSTRACT: Rare earth-free permanent magnets for applications in electro-magnetic devices promise better sustainability and availability and lower prices. Exploiting the combination of shape, magnetocrystalline and exchange anisotropy in 3D-metals can pave the way to practical application of nanomagnets. In this context, we study the structural and magnetic properties of Co80Ni20 nanorods with a mean diameter of 6.5 nm and a mean length of 52.5 nm, prepared by polyol reduction of mixed cobalt and nickel acetates. Structural analysis shows crystalline rods with the crystallographic caxis of the hexagonal close-packed (hcp) phase parallel to the long axis of the Co80Ni20 alloy rods, which appear covered by a thin oxidized face-centered cubic (fcc) shell. The temperature dependence of the surprisingly high coercive field and the exchange bias effect caused by the antiferromagnetic surface oxide indicate a strong magnetic hardening due to alignment of anisotropy axes. We identify a temperature dependent local maximum of the coercive field at T = 250 K, which originates from noncollinear spin orientations in the ferromagnetic core and the antiferromagnetic shell. This might be useful for building four way magnetic switches as a function of temperature.

H

Due to their low magnetocrystalline anisotropy energy density (MAE), purely metallic Fe or Ni nanorods do not have the potential to fulfill the needs of permanent magnets. Co crystallized in the hcp structure has about 1 order of magnitude larger MAE, rendering it a candidate for permanent magnets based on nanorods.13 Further, Co-based alloys may be tailored in a way comprising high magnetization, large magnetocrystalline anisotropy, and sufficient oxidation resistance by forming a thin surface oxide layer passivating the metallic core.14,15 Previous experiments have shown that Co nanorods with the magnetic easy axis (c-axis) aligned parallel to the long axis of the nanorods can be prepared.16,17 The thin CoO surface layer that chemically protects the metal core also offers an exchange bias system.18,19 At low temperatures, the interfacial exchange anisotropy leads to the characteristic shift of the hysteresis loop accompanied by an increase of the coercive field. With increasing temperature this magnetic hardening of the nanorods weakens and the gain in the coercive field vanishes above the Néel (blocking) temperature of the antiferromagnet, respectively. CoO has Néel temperature TN = 291 K in the bulk.19 In thin oxide layers of 3 nm and below or in nanoparticles and nanostructures, TN might be significantly

ighly elongated nanomagnets like nanorods and nanowires have recently attracted huge interest due to possible applications in magnetic memories,1 sensors,2 and microwave devices3 or the use as permanent magnets.4 The magnetic properties of such nanorods with a length-to-diameter aspect ratio larger than 5 are strongly governed by the shape anisotropy, and thus, the easy axis of magnetization aligns parallel to the long axis of a nanorod. If the preparation technique or the synthetic method exerts control over the crystallographic orientation of the material along the nanorod geometry, one expects large coercive fields and high remanent magnetization stemming from the magnetocrystalline easy axis aligned parallel to the axis of the rod. Additionally, besides the shape and magnetocrystalline, any further increase of magnetic anisotropy, e.g., by exchange coupling to an antiferromagnetic material (exchange bias), will increase the energy product of the magnetic hysteresis of these nanostructures.5,6 Systems matching the above criteria are very interesting for permanent magnet applications as already suggested in the late 1950s.7−9 Nowadays, strong permanent magnets for various applications usually contain a significant amount of rare earth elements like Nd or Sm.10 For better sustainability, availability, and lower prices, the above idea of exploiting the combination of shape, magnetocrystalline, and exchange anisotropy in rare earth-free nanomagnets has received renewed attention over the last years.11,12 © 2015 American Chemical Society

Received: March 15, 2015 Revised: May 13, 2015 Published: May 13, 2015 4015

DOI: 10.1021/acs.chemmater.5b00976 Chem. Mater. 2015, 27, 4015−4022

Article

Chemistry of Materials reduced due to finite size effects or thermal excitations, leading to a vanishing exchange bias.19,20 Thus, in order to exploit such effects for permanent magnets operating at ambient temperature, materials with significantly higher TN are needed while maintaining the large MAE. In this regard, a good compromise between large MAE and high TN in the antiferromagnetic shell is the admixture of Ni, since NiO has a TN = 525 K in the bulk.13,21 For example, (Co59Ni41)O thin films have shown TN = 383 K.22 Different chemical methods based on electrodeposition in anodized alumina nanopores,23 hydrothermal processes,24,25 or organometallic chemistry26−28 have been used to synthesize 3d elemental and alloy nanorods and nanowires. In this work we have chosen the polyol process for the production of nanostructures of this alloy,16,29 optimizing both morphology and stoichiometry for the improvement of the magnetic properties. Accordingly, we report on the structure and magnetic properties of Co80Ni20 nanorods with a mean length of 52.5 nm and an aspect ratio of about 8, which according to standard micromagnetic simulations is close to the one for optimum coercive field. The stoichiometry has been chosen to exploit the highest magnetocrystalline anisotropy among the 3d-elements (hcp Co), the still large magnetization at low Ni content, and the good long-term stability of CoNi alloys.14 Indeed, the oxides form a passivating layer around the ferromagnetic core in ambient conditions and can promote a larger hysteretic energy product by exchange anisotropy between an antiferromagnetic oxide shell and the ferromagnetic metallic core. We quantitatively compare the present results on Co80Ni20 core/oxide shell nanorods with Co/CoO core−shell rods.

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NiCo2O4 powder were taken from literature33 and shifted by 1.4 eV toward higher binding energies accounting for different calibration. The information depth as defined by the inelastic mean free path of photoelectrons in matter at kinetic energies of about 700 and 630 eV for the Co and Ni 2p doublets is about 1.1 nm.34 Thus, we predominantly probe the chemical state of the nanorod’s surface, i.e., the oxide shell. The magnetic properties were characterized by superconducting quantum interference device (SQUID) magnetometry using a Quantum Design MPMS XL system and by vibrating sample magnetometry using a Quantum Design Dynacool PPMS system. Dried powder of Co80Ni20 nanorods was weighted and compressed in a capsule. The mass magnetization was calculated neglecting the weight of residual ligands with an estimated maximum error bar smaller than 10%. For all hysteresis measurements, the background signal from the capsule, which had been measured separately, was subtracted. The maximum field, temperature, and sample history are indicated in the figure captions. Before cooling, the sample was carefully demagnetized at 350 K by applying a decreasing alternating magnetic field.

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RESULTS AND DISCUSSION Morphology. The CoNi specimen has been investigated by TEM. Figure 1a presents a typical TEM micrograph showing nanostructures mainly with an elongated shape, though two different morphologiesnanorods and spherical particlescan be clearly identified. Single nanorods show stronger contrast variations compared to the spherical particles, suggesting a

EXPERIMENTAL METHODS

Materials. Cobalt acetate tetrahydrate (reagent grade), nickel acetate tetrahydrate (99%), oleic acid (90%), trioctylphosphine (90%), and ruthenium trichloride hydrate (45−55% Ru) were purchased from Sigma-Aldrich. Sodium hydroxide was obtained from Riedel-de Haën. 1,2-Butanediol (98%) was purchased from Fluka. Absolute ethanol (99.97%) and tetrahydrofuran (99.6%) were purchased from Analar Normapur. All the chemicals were used as received without further purification. Synthesis. Cobalt−nickel alloy nanorods were synthesized by the polyol reduction of the metal precursors according to a recipe developed previously.16,30−32 In a typical synthesis, cobalt acetate tetrahydrate (4.8 mmol), nickel acetate tetrahydrate (1.2 mmol), sodium hydroxide (1.2 mmol), and ruthenium trichloride hydrate (0.15 mmol) were dissolved in 75 mL of 1,2-butanediol. The solution was heated up to 170 °C using magnetic stirring. Once the mixture reached this temperature, oleic acid (6 mmol) and trioctylphosphine (6 mmol) were added, keeping the solution at this temperature for 20 min. After cooling to room temperature the precipitate was collected and washed with absolute ethanol several times. The particles were finally redispersed in tetrahydrofuran. Characterization. Transmission electron microscopy (TEM) including spatially resolved chemical analysis was conducted in a Tecnai F20 Supertwin microscope operated at 200 kV with a field emission gun and equipped with an energy dispersive X-ray (EDX) detector. Samples for structural and compositional analysis were prepared by slow evaporation of a highly diluted suspension on a standard carbon-coated copper grid in ambient conditions. For X-ray photoelectron spectroscopy (XPS), as-prepared CoNi nanorods have been deposited on a Si/SiO2 substrate at submonolayer coverage and imaged by scanning electron microscopy (not shown). We investigated the specimen using nonmonochromatic Al Kα radiation (1486.6 eV). The spectrometer has been calibrated by Au and Cu XPS lines. Reference spectra of a CoO single crystal and

Figure 1. (a) TEM micrograph of CoNi nanorods and particles. Panel (b) presents the corresponding size (length) distribution. The aspect ratio distribution of the nanorods is displayed in panel (c). Fitting results of the log-normal distributions (red) are given in Table 1. 4016


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Chemistry of Materials polycrystalline structure. The size (length) distribution of the nanocomposites displayed in Figure 1b is bimodal. The two contributions arise from spherical nanoparticles (5−15 nm) and nanorods with the length varying from 20 to 80 nm, respectively. In addition to the overall size distribution, the aspect ratio of the nanorods defined as the ratio of length L to diameter D is shown in Figure 1c. We observe a rather broad distribution of the aspect ratio varying from 3 to 15. Size and aspect ratio distributions were fitted by log-normal functions.35,36 The resulting mean dimensions of the nanoparticles and -nanorods are listed in Table 1. In the following, we are Table 1. Results of the Morphological Characterization by TEMa

Figure 2. Co 2p XPS of the CoNi nanorods (black). Additionally, reference spectra of CoO (red) and NiCo2O4 (blue) powder samples are taken from the literature.33 Vertical dashed lines indicate the peak positions of CoNi nanorods. The reference spectra were vertically shifted for clarity.

nanoparticles 7 ± 2 vol %

nanorods 93 ± 2 vol % length L

diameter D

aspect ratio L/D

diameter d

52.5 ± 19.4 nm

6.5 ± 1.0 nm

8.2 ± 2.4

8.2 ± 2.3 nm

a

Size and aspect ratio distributions are determined for more than 300 CoNi nanorods and particles assuming lognormal distributions.

oxidized CoNi rods with NiCo2O4 powder as a reference33 for ferrites, it is obvious that the satellite peaks are absent for the ferrite reference and the doublet separation (15.0 eV) is significantly smaller as compared to the CoNi rods (16.0 eV). Thus, we can exclude the formation of ferrites. Further, we determined from the total XPS line intensities the Co-to-Ni elemental ratio to 5.6 ± 1.5. The error bar arises from the poor signal-to-noise ratio of the Ni 2p spectrum. XPS predominantly probes the surface stoichiometry of the specimen, thus confirming the larger Co content in the oxide surface as suggested by the EDX linescans (cf. Supporting Information Figure S2). Structural Analysis. Structure determination has been carried out by means of high-resolution (HR)TEM and selected area electron diffraction (SAED). Indexing the ring patterns obtained in SAED (Figure 3a, inset) indicate a hcp structure with lattice parameters of a = 0.247 ± 0.006 nm and c = 0.400 ± 0.010 nm corresponding to those of bulk Co80Ni20 alloy (a = 0.250 nm, c = 0.407 nm).15,38 A diffraction ring labeled “oxide” arises most probably from the surface oxide. The exact oxide could not be determined, since the diffraction ring might be associated with NiO, CoO, Co3O4, or a ternary oxide CoxNiyOz. More detailed information on the structure and composition of the CoNi nanorods can be obtained by analysis of the HRTEM images. Figure 3a shows a typical HRTEM image of the center part of a CoNi nanorod. The contrast distribution indicates a core−shell structure with a single crystalline metallic core covered with a polycrystalline oxide shell. The fast Fourier transform (FFT) pattern of the image (Figure 3b) consists of two sets of reflections. The redencircled diffraction spots correspond to the hcp structure viewed along the [21̅ 10̅ ] zone axis. The yellow-encircled reflection set is typical for the fcc structure imaged along the [011] direction. Thus, the diffraction pattern can be interpreted as the sum of overlapping hcp and fcc diffractions. The FFT pattern analysis yields the lattice constants a = 0.252 ± 0.006 nm and c = 0.408 ± 0.010 nm related to the hcp structure of Co80Ni20.15,38 The hcp ⟨0001⟩ direction is parallel to the long axis of the nanorod (see also Figure S3 in the Supporting Information) and coincides with the fcc ⟨001⟩ direction. This suggests an epitaxial relation between the surface oxide layer (formed in ambient conditions) and metal core with ⟨0001⟩hcp ∥ ⟨001⟩fcc. The lattice parameter of the cubic phase is found to

using the obtained mean values of 52.5 and 6.5 nm for the length L and diameter D of the nanorods, and d = 8.2 nm for the diameter of nanoparticles, respectively. Further, we determined the volume portions of nanorods (93 vol %) and particles (7 vol %) from the log-normal fits assuming cylinders and spheres, respectively. Thus, we assume the magnetic properties to be strongly governed by the nanorod morphology. Chemical Analysis. The elemental concentration in CoNi nanorods has been measured by means of EDX spectroscopy in the TEM (Figure S1 in the Supporting Information). The average composition is determined from a larger sample area containing hundreds of nanorods and -nanoparticles. We find a Co to Ni elemental ratio of 4:1 matching the initial precursor concentration in the synthesis. EDX spectra were also obtained from various points of individual nanostructures showing a relatively uniform distribution of Co and Ni within the particle core. See example in Figure S2 in the Supporting Information. Thus, we conclude that Co80Ni20 alloy nanorods were successfully prepared. Additionally, we qualitatively observe a reduced amount of Ni toward the nanorod surface when comparing the individual Co and Ni elemental intensities along the scan (cf. Supporting Information Figure S2). The as-prepared Co80Ni20 nanorods were further characterized by XPS after deposition on Si/SiO2 substrates. Note that the Ni content is small and, thus, spectra have a poor signal-to-noise ratio, and accordingly, herein we restrict the detailed discussion to the Co 2p XPS doublet. Figure 2 presents the Co 2p doublet of the specimen and two reference spectra. We observe binding energies of 781.8 (797.8) eV for the Co 2p3/2 (Co 2p1/2) lines and a doublet splitting of Δ = 16.0 eV. Moreover, satellite peaks are observed toward higher binding energies. According to the NIST reference database37 these observations suggest the formation of Co(OH)2 on the particle surface, while CoO has slightly lower binding energies at a doublet splitting of Δ = 15.5 eV. The reported values in the XPS database scatter by about ±0.5 eV, and thus, an unambiguous proof of a certain chemical state is not possible here. On the other hand, when CoNi rods are exposed to air one may consider the formation of NixCo3−xO4 ferrites with 0 < x < 1. However, from comparing the spectra of the surface4017


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Chemistry of Materials

Figure 4. Magnetic hysteresis loops after cooling from 350 to 5 K in zero field (ZFC) and in an applied field of 4.5 T (FC). The FC hysteresis has larger coercive field and a significant shift toward the negative field direction (exchange bias). The inset presents the high field susceptibility. The FC hysteresis is also shifted toward larger magnetization indicating pinned magnetic moments. Figure 3. HRTEM image of the center part of a CoNi nanorod with a clear core−shell contrast (a). The inset shows the electron powder diffractogram from a larger sample area proving the hcp structure of nanorods. Panel (b) is the FFT pattern from the HRTEM image in panel (a). Red (yellow) circles mark spots arising from hcp (fcc) structure, respectively. Note that the hcp (0002) spot overlaps with the fcc (2̅00) spot. The hcp c-axis is parallel to the long axis of the nanorods as indicated in (a). The lower panels present structurally filtered images by choosing hcp spots (c) and fcc spots (d), respectively.

T = 5 K. The FC loop is shifted opposite to the cooling field direction indicating an FC-induced unidirectional magnetic anisotropy. A large exchange bias field μ0HEB = μ0|HC1FC + HC2FC|/2 = 290 mT is measured at 5 K. The FC coercive field μ0HCFC = μ0|HC1FC − HC2FC|/2 = 305 mT is increased by 80% as compared to the ZFC value. The broadening of the FC hysteresis is due to an additional FC-induced uniaxial magnetic anisotropy. Besides, a vertical magnetization shift of +1.7% is found (Figure 4 inset). All three phenomena are well-known for systems with ferromagnet (FM)/antiferromagnet (AFM) interfaces due to interfacial exchange coupling between the magnetic moments of the FM and AFM regions.41 Thus, the Co-rich oxide layer as determined by TEM and XPS is antiferromagnetic and increases the magnetic anisotropy of the nanorods at least at low temperatures. The temperature dependence of the coercive field (μ0HC) after ZFC or FC and the exchange bias field (HEB) are displayed in Figure 5. The comparison of the coercive field values of ZFC and FC measurements (blue triangles and red dots) shows two different parts. Above T1 = 100 K both curves are congruent. Below T1, μ0HC of FC measurements increases sharply with decreasing temperature as compared to the ZFC

be 0.417 ± 0.008 nm and can correspond to the lattice parameter of either CoO or NiO oxides.39 Studies on the oxidation of Co−Ni alloys showed that Co, due to the higher diffusivity of Co2+ ions in the oxide, segregates to the surface.40 Taking into account the expected oxidation behavior of Co−Ni and the corresponding experimental data from EDX and XPS, we can exclude the possibility of NiO oxide formation. Most probably, a Co-rich fcc oxide shell forms. The HRTEM images in Figure 3c,d are reconstructed from the FFT spot pattern (Figure 3b) by filtering selected diffraction spots. Figure 3c shows the inverse FFT image using only the red-encircled reflections, which correspond to the hcp structure. The reconstructed image shows a uniform distribution of the hcp structure within the rod. The reconstruction of the image using yellow-encircled fcc reflections in Figure 3d shows a preferential allocation of this phase at the nanorod surface. Accordingly, we conclude that the core−shell structure arises from a metallic hcp CoNi core covered with a 1−2 nm polycrystalline, strongly textured fcc Co-rich oxide shell. It is worth noting that the shell thickness varies along the nanorods having the largest thickness at the tips. Magnetic Properties. Magnetic hysteresis loops of dried powder of Co80Ni20 nanorods were measured in capsules in the temperature range 5−350 K in external fields up to 4.5 T. In order to investigate the influence of the oxide surface layer on the magnetic properties, zero field cooling (ZFC) and field cooling (FC) in 4.5 T from T = 350 to 5 K has been applied. See representative hysteresis loops at different temperatures in Figures S4 and S5 in the Supporting Information. Note that the sample has been carefully demagnetized at T = 350 K before cooling. Figure 4 presents the ZFC and FC hysteresis loops at

Figure 5. Coercive fields and exchange bias field as a function of temperature after cooling the sample from 350 to 5 K in zero field (ZFC) and 4.5 T (FC). The experiments were performed with increasing temperature. 4018


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describing the average decay of the magnetization with time. This viscosity model is suitable for the description of long-time relaxation at times t > 100 s, as it was observed in the present study (Figure 6a). The linear slope S in logarithmic scale is a measure of the magnetic viscosity. Note that small values of S correspond to a high stability of M. It is worth mentioning that the remanent magnetization (as deduced from the viscosity measurements at t = 104 s) follows the identical temperature dependence as compared to the coercive field (cf. Figures 5 and 6a). The temperature dependence of S is displayed in Figure 6b. Error bars have been estimated by applying different fits to experimental data. S has a small value at low temperatures where the majority of nanorods and particles are thermally blocked, and, therefore, magnetic relaxation is very weak. With increasing temperature S grows to a local maximum near 150 K, which is slightly below temperature T2 in Figure 5, where a minimum of the coercive field was observed. Similar behavior has been reported for Co/ CoO nanowires.17 Above, S decreases to a local minimum at about T = 220 K, again slightly below the local maximum of the coercive field at T3, before S rises to the highest observed values around ambient temperature. It is straightforward to understand these opposing trends of HC and S dependencies with temperature, since the larger the stability of the magnetization (indicated here by a high HC), the slower its decay with time (low S). The significant increase of S observed at higher temperatures (T > 220 K) can be related to the temperature variation of the different anisotropy contributions in the system.5 However, the underlying mechanism for the formation of a local maximum of HC and the local minimum of S cannot be explained with these observations alone. In most exchange bias systems (films, nanorods, or nanoparticles) such an oscillatory temperature dependence of HC is not observed. Some systems, i.e., highly textured MnF2/ Fe films,43 highly ordered Ni/NiO nanorods,44 or Co/CoO nanorods,17 exhibit an increase of HC at intermediate temperature while for spherical particles this behavior has not been observed, yet. For MnF2/Fe films Leighton et al.43 found an enhancement of HC at the Néel temperature of the AFM MnF2 only when the FM Fe layer thickness is significantly larger as compared to the AFM layer, thus making the FM layer dominant upon reversal. They discussed this effect in the frame of the Stiles and McMichael model, in which competing anisotropy energies of the FM and AFM layers are linked to specific crystallographic orientations.45 Hsu et al. argued that in the case of Ni/NiO nanorods spin frustration at the FM/AFM interface is important and HC increases with increasing temperature as soon as the superparamagnetic blocking temperature of the AFM shell is crossed.44 Nevertheless, Ni/ NiO-coated carbon nanotubes working as nanowires and being spin-frustated at the FM/AFM interface do not show an increase of HC at an intermediate temperature, likely because of the randomly oriented Ni/NiO nanostructures forming the magnetic coating of the carbon nanotubes.46 For Co/CoO nanorods with an epitaxial relation of the Co core and the CoO shell, Maurer et al.17 obtained similar trends as presented here for oxidized CoNi nanorods. They observe a local minimum of μ0HC = 450 mT and a maximum of S at the temperature at which the exchange bias vanishes T2 = 100 K, while HC rises up to 600 mT at T3 = 200 K.17 In the present contribution we find very similar temperature dependencies of HC and S, however, shifted by about 75 K toward higher temperatures. This can be assigned to the small Ni content in the shell leading to higher

values. This effect is due to the existence of an interface between ferromagnetic material (CoNi alloy core) and antiferromagnetic material (Co-rich oxide shell), which causes the unidirectional and uniaxial anisotropy. μ0HEB (black squares) gradually decreases from 290 mT at 5 K to vanish at T2 = 175 K. Most important is the increase of the coercive field from 70 mT to 130 mT above T2 up to T3 = 250 K. This behavior has only been observed in few systems (films and nanorods) all having a crystallographic relation between FM and AFM. For undirected oxide growth modes frequently observed for nanoparticles, nanorods, or films such intermediate increase of the coercive field is absent. Before starting the detailed discussion of the temperature dependence of HC, additional results on the thermomagnetic relaxation of the remanent magnetization M(t) as a function of time are presented. Such measurements give access to the socalled magnetic viscosity.5,42 After field cooling the sample from 350 K to the measuring temperature in an external field of 4.5 T, the time decay of the magnetization has been recorded at zero external magnetic field for temperatures in the 10−325 K range (Figure 6a). According to the model proposed by Street and Woolley42 the distributions of size, crystallographic orientation and anisotropy in an ensemble of magnetostatically interacting particles lead to the logarithmic decay of the magnetization M in time t: M(t) = M0 − S(T) ln(t/t0), where M0 is the magnetization at t0 = 0 s and S(T) is the slope

Figure 6. (a) Thermomagnetic relaxation of the remanent magnetization M(t) in logarithmic scale. All curves are measured at zero external magnetic field after cooling the specimen from 350 K to the indicated temperature in an external field of 4.5 T. Panel (b) shows the temperature dependence of the slope S taken from a fit M(t) = M0 − S(T) ln(t/t0) of the thermomagnetic relaxation curves. The spline function connecting the experimental points is a guide to the eye. Details are discussed in the text. 4019


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the shell without reaching TC and TN in the present experiments. Overall, one reasonable explanation of the magnetic hardening at intermediate temperatures is (i) there is a certain degree of crystallographic order between core and shell and (ii) the magnetic easy axis of the FM core (here c-axis and long axis of the rods) and the AFM shell (here presumably monoclinic Co-rich oxide below TN) are not collinear. These two conditions are fulfilled for the present study and all the examples discussed above. Thus, FM−AFM core−shell nanorods represent a prototype system for investigating the influence of SPM fluctuations in exchange bias. Vice versa, for spherical particles with core−shell symmetry, these conditions can hardly be matched, which is probably the reason why such behavior has not been observed so far.

TN and to a reduced effective magnetic anisotropy in the FM core and the shell of CoNi nanorods. Maurer et al. discuss these observations in a three regime model: CoO grains in the shell are AFM at low, become superparamagnetic (SPM) at intermediate, and get paramagnetic at high temperatures. Consequently, exchange bias is present until the SPM regime is reached (T2 = 175 K in Figure 5). Any further temperature increase leads to more and more SPM grains. Thus, its influence on the FM core decreases and the core’s magnetic moments at the interface can align easier along the long axis of the rods due to shape anisotropy. This leads to the higher squareness as indicated by the simultaneously increasing coercive field and remanent magnetization. Maurer et al. identified T3 as TN (cf. Figure 5), above which the influence of the AFM shell vanishes, and HC (S) decreases (increases) with temperature. We do not follow this point, but suggest that TN is higher and these observations should be discussed within a temperature-dependent vectorial effective anisotropy model comprising at least the major anisotropy components, i.e., magnetic anisotropy of the FM core and the AFM shell and the interfacial exchange. Since the above anisotropy values are unknown or scatter significantly for nanoscale systems in the literature and the present study cannot deliver them precisely, we restrict ourselves to a qualitative discussion in conjunction with the structural results. The magnetic easy axis of the FM CoNi core is the nanorod long axis with collinear orientation of magnetocrystalline and shape anisotropies. The oxide shell grows quasi-epitaxial on the rods with one ⟨001⟩ direction parallel to the c-axis of the core. In the bulk, CoO and presumably also the present Co-rich oxide with some Ni incorporated have rock salt structure above and a monoclinic structure below TN with a large tetragonal distortion along the cube edges with c/a < 1 and a small deformation along ⟨111⟩. These result in a pseudocubic structure with two CoO units per unit cell.47 In the AFM state the sublattice spins align collinear approximately in ⟨110⟩ directions of the pseudocubic structure.48 Although significant additional stress at the interface can be expected for the CoNi nanorods, we speculate that the spin orientation in the AFM shell is similar and thus not collinear to the ⟨100⟩ direction and the c-axis of the FM core. This misalignment of FM and AFM spins at the interface certainly has strong impact on the magnetic properties. Considering the stronger magnetic anisotropy of Co oxides as compared to CoNi at low temperatures it is likely that interface spins are governed by the AFM shell. From the fact that the Curie temperature of the CoNi core is much larger than TN of the shell, one may assume that the magnetic anisotropy contributions have strongly different temperature dependencies. Then it gets obvious that the major source of magnetic anisotropy should switch from the AFM shell to the FM core upon rising temperature. Below T1 = 100 K, fieldinduced orientation of rotatable AFM interfacial spins leads to the splitting of the coercive field after ZFC and FC. The temperature where the exchange bias vanishes (T2 = 175 K) is the temperature limit of pinned interface moments, and the above AFM moments become rotatable, contributing to the uniaxial anisotropy. For higher temperatures the influence of the AFM moments gets weaker and, consequently, more and more interface spins align with the FM core as suggested by the rising remanent magnetization up to T3 = 250 K. The decrease at higher temperatures is presumably due the significant decrease of all anisotropy contributions in both the core and

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CONCLUSIONS Co80Ni20 nanorods with average length (diameter) of 52.5 nm (6.5 nm) have been chemically prepared for testing them as permanent magnets, by exploiting the alignment of magnetocrystalline and shape anisotropy axes in conjunction with exchange bias. Structural, morphological, and chemical investigations revealed surface oxidized nanorods consisting of a metallic hcp CoNi core and a Co-rich oxide shell with a thickness of 1−2 nm. The crystallographic c-axis in the core is collinear with the long axis of the nanorods. Co-rich oxides have grown around the core with a defined crystallographic orientation, i.e., ⟨0001⟩hcp ∥ ⟨001 ⟩fcc. The magnetic characterization of this system shows strong unidirectional anisotropy at low temperatures after field cooling while the exchange bias vanishes at T = 175 K. Interestingly, in an intermediate temperature range 175 K < T < 250 K the coercive field rises by almost a factor of 2, before it starts decreasing again for higher temperatures. This effect is discussed within a three regime model suggested before, considering superparamagnetic fluctuation of antiferromagnetic grains in the Co-rich oxide shell. Such behavior is expected only when the magnetic easy axes of the ferromagnetic core and the antiferromagnetic shell are not collinear.

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ASSOCIATED CONTENT

S Supporting Information *

Additional information on the structural (S1−S3) and magnetic (S4−S5) properties is presented in the Supporting Information. Figure S1. EDX spectrum of CoNi nanorods on a carbon film supported copper TEM grid. Figure S2. (a) High-angle annular dark-field scanning TEM image of the CoNi nanorod. (b) Line scan profiles of the characteristic Ni−K and Co−K X-ray lines as a function of electron beam position on the nanorod in panel (a). The red line in the image (a) indicates the path of electron beam during the scan. Figure S3. (a) TEM image of the Co80Ni20 nanorods deposited from the solution on a carbon film supported copper grid in the presence of an external magnetic field B of 1.5 T (indicated by the blue arrow). (b) Corresponding electron diffraction pattern. The inhomogeneity of the diffraction rings indicates strong ⟨0001⟩ texture along the long axis of the nanorods. Representative magnetic hysteresis loops of the Co80Ni20 nanorod powder recorded after ZFC (Figure S4) and FC (Figure S5) at different temperatures between 5K and 300 K. This material is available free of charge via the Internet at http://pubs.acs.org/.The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.5b00976. 4020


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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Present Address ∥

(A.S.S.) Helmholtz-Zentrum Dresden-Rossendorf, Bautzner Landstr. 400, 01328 Dresden, Germany. Author Contributions

All authors have given approval to the final version of the manuscript. Funding

We acknowledge funding from the European Community’s Seventh Framework Programme (FP7-NMP) under Grant Agreement No. 280670 (REFREEPERMAG). V.S. also acknowledges funding from the Xunta de Galicia Regional Government (Spain) under Project EM2014/035 (Emerxentes). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank Michael Winklhofer for fruitful discussions. A.S.S. acknowledges financial support by the University of DuisburgEssen.

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Structure-Correlated Exchange Anisotropy in Oxidized Co80Ni20

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