Topotaxial Phase Transformation in Cobalt Doped Iron Oxide Core

Subscriber access provided by Fudan University

Article

Topotaxial phase transformation in cobalt doped iron oxide core/shell hard magnetic nanoparticles Alberto López-Ortega, Elisabetta Lottini, Giovanni Bertoni, César de Julián Fernández, and Claudio Sangregorio Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.6b04768 • Publication Date (Web): 10 Jan 2017 Downloaded from http://pubs.acs.org on January 13, 2017

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a free service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are accessible to all readers and citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

Chemistry of Materials is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Chemistry of Materials

Topotaxial phase transformation in cobalt doped iron oxide core/shell hard magnetic nanoparticles Alberto López-Ortega†‡*, Elisabetta Lottini†, Giovanni Bertoni§, César de Julián Fernández‡, Claudio Sangregorio⊥* †

INSTM and Dipartimento di Chimica “U. Schiff”, Università degli Studi di Firenze, I-50019 Firenze, Italy CIC nanoGUNE, E-20018 Donostia-San Sebastian, Spain. § CNR-IMEM, Parco Area delle Scienze 37/A, I-43124 Parma, Italy. ⊥ INSTM and CNR-ICCOM, I-50019 Firenze, Italy. ‡

ABSTRACT Core-shell nanoparticles based on a CoxFe1-xO rock-salt core, and on a shell corresponding to cubic spinel CoxFe3-xO4, have been systematically annealed in order to completely oxidize and generate the fully ferrimagnetic cobalt ferrite structure. The annealing has been performed through a solventmediated process at high temperatures to avoid interparticle aggregation, usually observed in classical annealing methods. We carefully describe the oxidative process occurred during the initial shell passivation and in the following O2 mediated oxidation. It has been found that the rock-salt to spinel transformation occurs via topotaxial growth over the (200)RS//(400)S and (220)RS//(440)S planes shared between the two structures. This chemical transformation depends on the amount of divalent cobalt atoms present in the oxide structures. Within this respect, the solvent-mediated annealing process permits the release of a small amount of divalent cations, allowing the stoichiometry rearrangement required to form the spinel phase. The growth occurs through a topotaxial process which involves the formation of a mosaic texture of small spinel subdomains, separated by antiphase boundaries, into the well-defined rock-salt structure along the NPs. The existence of antiphase boundaries gives rise to the presence of exchange bias phenomena even for completely oxidized nanoparticles. The exchange bias effect increases the energy product of these nanocomposites, making this approach appealing for the 1 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 2 of 30

realization of a novel class of free rare-earth permanent magnets. INTRODUCTION Magnetic nanoparticles (NPs) have novel fundamental properties arising from their nanoscale dimension, which can be exploited in a large number of applications such as biomedicine,1,2 catalysis,3,4 magnetic recording,5 and permanent magnets.6,7 Concurrently, many different synthetic approaches have been developed in order to obtain magnetic structures with tailored properties, combining the requirement of large scale production with fine control on the morphological and structural properties of the NPs.8–12 Specifically, the high boiling point solvent thermal decomposition approach has established as one of the most reproducible methodologies to obtain highly crystalline NPs with a precise control on their size, morphology and properties.13 Many different materials from pure metals to oxide based ceramic structures have been synthesized through classic thermal decomposition approaches using carbonyls14 or acetates/acetylacetonates13,15 complexes as metal precursors. More recently, the decomposition of a Fe3+-oleate complex in high boiling point solvents has been proposed as an effective alternative route to synthesize highly monodisperse iron oxide NPs.16 However, even if this methodology has been extended to other transition metal oxides such as manganese17 or cobalt,18 it has been found that the long organic oleate chain of the precursor strongly affects the final structure of the as-synthesized NPs. Indeed, the synthetic approach often leads to the formation of core/shell (CS) NPs rather than single phase oxides.19,20 As a matter of fact, the decomposition of the carbon-based ligand into CO and H2 species forces the reduction of the Fe3+ ions to Fe2+ favoring the growth of the rock-salt, wustite (Fe1-xO),structure in place of the expected magnetite (Fe3O4) or maghemite (γ-Fe2O3), spinel phase.21 The following partial surface oxidation to the spinel structure generates therefore a rather complex structure, corresponding to a Fe1-xO/Fe3O4 CS architecture. Interestingly, many authors have reported this effect in the literature, with special emphasis on the fascinating magnetic properties of the system due to their bi-magnetic character, i.e., 2 ACS Paragon Plus Environment

Page 3 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


antiferromagnetism (AFM) of the Fe1-xO core and ferrimagnetism (FiM) of the Fe3O4 or γ-Fe2O3 core 19,20,22,23

Recently, Wetterskog et al.24 have shown that a subsequent oxidation step to transform CS NPs

into single phase spinel structures introduces structural defects leading to the formation of antiphase boundaries, which can explain their anomalous magnetic properties. In this work, we have used a mixed Fe3+,Co2+-oleate complex to synthesize 10 nm FexCo1-xO-CoFe2O4 CS NPs, and we have studied their following transformation in pure cobalt ferrite by solvent-mediated annealing. We carefully describe the oxidative process occurred during the initial shell passivation and the following O2 mediated oxidation. We explain the morphological, structural and magnetic changes induced by these two processes in terms of topotaxial growth with the formation of a mosaic texture in the final spinel structure separated by antiphase boundaries. Moreover, we show the strong effect of the inclusion of Co2+ ions into the iron spinel structure. The high reduction potential of the Co2+ ions and the cobalt-to-iron ratio play dominant role into the oxidation process, forcing the partial release of Co2+ ions into the solution, a mechanism allowed by the solvent annealing process.. The simultaneous presence of high anisotropy and exchange bias makes these hybrid nanomaterials promising to be employed as a novel class of free rare-earth magnets.25

RESULTS As-prepared NPs (sample M0) were synthesized following a previously reported procedure.16,25 The resulting NPs present spherical shape with an average particle diameter of 11(1) nm. Particle size histograms were fitted using a Gaussian distribution, showing a unique size population with standard deviation less than 10% (see Fig. 1a). The resultant NPs have a CS structure, which, as we have discussed in a previous report,20,25 comprises a core region with the typical structure of the rock-salt phase CoxFe1-xO, and a shell which corresponds to the cubic spinel CoxFe3-xO4.20 3 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 4 of 30

In order to better understand the formation and stabilization of the CS structure and on their subsequent complete oxidation, the as-prepared sample (M0) has been subjected to a systematic solvent-mediated annealing process. The method permits heat treatments of an ensemble of NPs at high temperatures (i.e., 300 ºC) avoiding the interparticle aggregation usually observed in typical annealing methods.26 To this aim, sample M0 has been dispersed in 1-octadecene, and heated up to 300 ºC under a constant air bubbling for two different times, i.e., 5 min (M1) and 15 min (M2), in order to completely oxidize the CS structure. Larger dwelling times (≥ 30 min) led to the dissolution of the NPs, probably due to the high temperature used and to the presence of traces of oleic acid from the previous synthetic step. Figure 1b-c depicts TEM images form M1 and M2 NPs. No particle growth occurs during the heating process, the size distribution being the same of M0, i.e., 11 (1) nm. Moreover, the NPs do not suffer of agglomeration as observed in standard oven annealings.26 However, HRTEM images revealed a slightly modification in the morphology, the annealed NPs exhibiting a more faceted structure (see Fig. 1d-f). M1 and M2 have indeed preferential faceting over the (100) and (110) planes as typically observed in NPs with cubic crystallographic structure.27 Powder XRD diffraction patterns for M0, M1, and M2 are shown in Figure 2 and Figure SI1. The pattern of M0 confirms the presence of two crystallographic phases, being characterized by the typical peaks of the rock-salt (RS) phase with a minor contribution from the cubic spinel (S) structure. A clear structural evolution is observed upon annealing, although for both oxidized samples the pattern does not match the one expected for a completely oxidized structure. Indeed, on one hand, the strong increase in the relative intensity between the (311)S and (111)RS at ca. 2θ = 36º, the narrowing of (311)S, and the increase in intensity of the (511)S peak at 2θ = 57º are clear evidences of the growth of the spinel phase. On the other hand, the pattern is still dominated by the (200)RS and (220)RS peaks at 2θ = 42º and 2θ = 61º, suggesting that the rock-salt crystal structure is preserved during the process. 4 ACS Paragon Plus Environment

Page 5 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Although the full width at half maximum (FWHM) of the diffractions peaks of the annealed samples is unaltered with respect to M0 (1.05º), a shift towards larger angles is observed upon oxidation, indicating a reduction of the cell parameters for both phases (see Table 1). Rietveld analysis confirmed a clear phase evolution during the annealing process. Firstly, the calculated cell parameters for rock-salt and spinel structures in M0 are rather different from those expected for bulk phases. Both parameters are strongly influenced by the extremely confined character of the CS system,25 where the shell and the core suffer an expansion and a contraction, respectively. Conversely, after annealing the cell parameter of the spinel phase decreases significantly, reaching the typical value of cobalt ferrite NPs,7,28 while only a small abatement is observed for that of the rock-salt, to reach a low mismatch between the two phases (0.1%). The relative weight fraction of the spinel phase increases from 67% in M0 to 80% and 82% in M1 and M2, respectively, indicating that the annealing process does not allow the complete oxidation of the as-prepared NPs. Finally, it should be noted that the dwelling time has no effect on the NP structure since all the parameters of both phases are very similar for M1 and M2. In order to investigate the local composition of the NPs, EEL analysis was performed (Figure 3).29 Elemental quantification of M0 confirms the expected CS structure with a non-homogeneous distribution of iron, cobalt, and oxygen atoms along the NP diameter (see Fig. 3a-c). In particular, the NP consists of a Co0.4Fe0.6O core 6.5 nm in diameter, surrounded by a 2 nm CoFe2O4 shell, separated by a relatively sharp interface. The formation of a CS structure is confirmed by EEL mapping of iron ions at the two different oxidation states on a single NP (Fig. 3d): while Fe2+ ions are only placed in the core, Fe3+ are mainly located into the shell region. Conversely, after the annealing process, EEL analysis revealed the loss of the CS structure and the oxidation towards the spinel structure. In fact, the metal (Co + Fe) and O percentage in M1 and M2 are constant along the NP diameter and close to the value expected for a pure spinel structure (M3O4) while, EEL mapping displays a uniform distribution of Fe2+ and Fe3+ ions in the whole NP. In addition, the Co percentage in M1 is considerably lower than 5 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 6 of 30

in M0, suggesting a partial loss of this metal ion during the wet-annealing process. To confirm this data, the total amount of cobalt and iron in a single NP was calculated by integrating the EELS signal for several NPs: a 25% decrease of the total Co amount was observed in the annealed samples. XMCD spectra at Fe and Co L2,3 edges recorded at low temperatures in samples M0 and M1 allowed us to further study the oxidative process using a non-local experimental technique, thus providing an information averaged over a large number NPs. The Co profile is the same for the two samples, equal to that expected for Co2+ ions in octahedral coordination.30,31 Conversely, at the Fe edge, while M0 sample shows the two peaks characteristic of Fe2+ and Fe3+ in octahedral sites, after annealing the spectrum resembles that expected for the spinel structure, where an extra peak appears corresponding to the Fe3+ in tetrahedral coordination (see Fig. 4).10,32 In order to go further in depth with the structural characterization, the local structure of the NPs was analysed on HRTEM images by fast Fourier transform (FFT), Bragg filtering images and geometric phase analysis (GPA)33,34 (see Fig. 5 and Fig. 6). FFT images of M0 and M1 show the spots expected for spinel and rock-salt phases (see Fig. S2a,b).25,32 The contributions of the two phases can be separately studied by considering the corresponding reflections. The one at 0.30 nm distance can be indexed exclusively as the (220)S spinel planes. The ones at d = 0.21 nm and 0.15 nm contain the contributions from both rock-salt and spinel structure (i.e., (200)RS/(400)S and (220)RS/(440)S, respectively). We have simulated the expected contributions from the two phases to these reflections (considering approximately the thickness of the core and the shell in the NPs form EELS data) and we found a ratio of ~ 3 between rock-salt and spinel intensities (mostly due to the higher dynamic scattering factor of these reflections in the rock salt-structure). The corresponding diffractions are shown in Figure S2c,d. As a result, they can be used to get insight on the rock-salt structure of the core. For M0, the (220)S Bragg filtering image suggests a non-continuous distribution of the spinel phase, as it is expected for a shell morphology. Indeed, the image depicts a defective structure with a 6 ACS Paragon Plus Environment

Page 7 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


considerable number of defects. Conversely, both (200)RS/(400)S and (220)RS/(440)S images display a well-defined structure in all the NP without the presence of any defects on their periodicity. This welldefined structure of the rock-salt phase can be interpreted as due to the topotaxial growth of the shell phase over the initial rock-salt core. Indeed, GPA images confirm the presence of a uniform structure for the rock-salt phase, without significant lattice deformation (less than 2 %) throughout the NP.24 However, the (220)S outline shows two clearly differentiated regions, where a roughly constant deformation is observed from the particle surface to the core-shell interface and a negligible deformation region in the center of the particle (see Fig. 6a). This result confirms the CS architecture where the highly defective spinel phases is located in the shell region. Indeed, the deformation profile show a marked sign change on the phase-deformation profile at the interface between both phases strongly correlated with the formation of antiphase boundaries. On the other hand, similar results are obtained after the oxidation process (sample M1); (200)RS/(400)S and (220)RS/(440)S planes still present a well-defined structure along the whole NP even if the amount of the spinel phase is increasing; conversely, (220)S fringes are now spread all over the particles, showing a non-continuous structure with several defects similar to M0. However, in this case the deformation profile depicts a continuous sign change (i.e., antiphase boundaries) along the annealed particle. Main dilatation profiles revealed the loss of the CS structure with the (220)S profile showing with a randomly distributed deformation along the whole particle. In fact, the annealing process oxidizes partially the rock-salt core allowing the growth of the spinel phase in small subdomains with a defective non-continuous structure. Conversely, (200)RS/(400)S planes ascribed to the rock-salt structure still depict low deformation values along the NP. Figure 7a displays the temperature dependence of the zero field cooled and field cooled (ZFC-FC) magnetizations of M0, M1, and M2. For all samples the characteristic blocking process of single domain NPs is observed. The blocking temperature (TB) increases with the annealing time, being 190, 7 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 8 of 30

220, and 230 K for M0, M1, and M2, respectively (see Table 2). The dependence of TB in this type of AFM-FiM bi-magnetic structures is rather complex to interpret. Indeed, for as-synthesized sample (M0), TB can be strongly affected by the exchange coupling between the core and the shell.35 This magnetic interaction, associated with the formation of an uniaxial magnetic anisotropy, in principle can lead to a net increase of TB well above the values commonly observed for the single FiM material with similar volume.36–38 Conversely, the obtained TB for M0 perfectly fits that of non-interacting CoxFe3xO4

NPs of 7-8 nm diameter (i.e. the same volume of the FiM shell NP).7,39,40 This result corroborates

the idea that the exchange coupling do not play a dominant role in the thermal stability of these NPs. Probably, the low magnetic field applied during the measurement is not large enough to completely align the interface spins and/or the AFM core is not a single domain.35 On the other hand, the reduction of the AFM core size after the annealing can give raise to a a strong modification of the magnetic properties and/or progressive loss of its magnetic order.25 In addition, the presence of antiphase boundaries in the annealed sample may also induce an extra anisotropy, although probably much weaker than the pure AFM/FM coupling.41,42 Indeed, we found TB scales fairly well with the amount of FiM phase, as predicted by the simple theory for superparamagnetic relaxation, where TB = KVFiM/25kB (K is the effective anisotropy constant and kB is the Boltzmann’s constant).43 In addition, in order to better study the oxidation process, magnetization curves at 30 kOe were recorded (Fig. 7b, bottom panel): for sample M0 the characteristic magnetic transition of the antiferromagnetic (AFM) rock-salt core is observed at TN = 235 K;25,44 conversely, the same transition is not present for annealed NPs, indicating the reduction of the AFM phase amount and/or the loss of the AFM structure after the oxidation process. Hysteresis loops were also measured at 10 K after a 120 kOe FC process. The loop of M0 displays the typical features expected for exchange coupled CS AFM/FiM NPs:25 a large exchange bias, HE, (19 kOe for M0) and a high coercive field, HC, (8.2 kOe), while due to the presence of the AFM phase, a low high field magnetization is observed (19 emu/g, 8 ACS Paragon Plus Environment

Page 9 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure S3). The oxidative process induces a strong modification of these features: MS increases till 38 emu/g, while HC and HE decrease reaching ca. 17 and 3.2 - 3.4 kOe, respectively for both M1 and M2 (see Figure S3). The increase on the magnetization is related to the larger FiM volume in annealed NPs; however, the value is still much lower than that expected for a cobalt ferrite (i.e., 80 - 90 emu/g). Moreover, the persistence of exchange coupling after the annealing could arise from the residual rocksalt structure and/or the presence of antiphase boundaries created during the spinel growth.24

DISCUSSION The rock-salt to spinel chemical transformation has been described in the literature to occur due to the well-known mediated diffusion and growth mechanism.24,45 According to this description, the oxygen deposition onto the rock-salt surface creates a potential gradient from the surface to the inner core.24,45 The invariant oxygen sublattice structure shared by the two phases and the small number of oxygen vacancies of the cation deficient, non-stoichiometric metal mono-oxide,46,47 only permits the cationic diffusion to maintain the electroneutrality of the whole structure.45 As a consequence, the phase transformation involves a change in the cation density and concomitantly in the metal and oxygen percentages, as well as structural rearrangements in the cation sublattice. Moreover, phase transformation in the first stage normally results in the creation of an interface between the transformed and the initial material. This process requires an extra energy which depends on how the structures match one to the other, and on the shape transformation between them (i.e., strain energy).48 However, due to the reduced lattice misfit between the two oxygen sublattices, the interface only involves a change in the distribution of the cations within the two structures, maintaining a simple topotaxial orientation relationship.49 Therefore, the spinel growth on the surface of the NPs, already observed in sample M0, takes place by oriented nucleation (i.e., topotaxial growth) on the (200)RS//(400)S and 9 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 10 of 30

(220)RS//(440)S planes. It should be noted that both planes have equivalent structure, where cations are intercalated in regular planes at similar distances. Indeed, oriented nucleation is the most favorable situation from both thermodynamic and structural points of view, due to the close similarity of the structures. In a second step, the reaction continues with the displacement of the rock-salt/spinel interface, preserving the two phases orientation during the reaction.50Indeed, the continuous structure in the (200)RS/(400)S Bragg filtered images and the lack of deformations at the core/shell interface observed by GPA analysis implies the formation of a homogenous long-range ordered arrangement of oxide and metal ions, corroborating, thus, the topotaxial growth of the spinel phase over the rock-salt core;48 Moreover, it has to be noted that the cell parameter compression/relaxation of the core/shell regions, characteristic of nanometric CS structures, permits to decrease the interface mismatch (i.e., elastic strains) between wustite-like and spinel phases.25,51 In fact, the lattice mismatch estimated from the XRD analysis presents values much smaller than expected, i.e., strains of ~ 2% (see Table 1),24 avoiding the progressive lattice relaxation from the core to the surface usually observed in structures with bigger shell51 and in epitaxial film.52 Conversely, the additional crystallographic planes of the spinel phase which are not present in the rock-salt structure are constrained only in the shell region, creating dislocations and only a clear antiphase boundaries at the core and shell interface.24 It should be noted that the rather uniform strain deformation in the shell region for the (220)S planes in M0 suggests that the oxidation initially occurs through the formation of a continuous, well-defined structure rather than through the predicted formation of small subdomains.24,53 The following oxidation (samples M1 and M2) occurs through a similar vacancy-assisted cation migration mechanism,47,54 which in fact is strongly dependent on the oxygen activity, ion diffusion constants, structural defects, and vacancy concentration.47,55,56 The oxidation process should lead to the characteristic particle size increase due to the oxygen deposition.47 Indeed, even if the mean diameter of the particles do not change with the annealing process, the formation of faceted particles results in a 10 ACS Paragon Plus Environment

Page 11 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


slightly increased volume with respect to the as-prepared ones (almost spherical). In principle, as it has been observed in the literature for similar systems, the temperature used in the annealing process should be high enough to completely oxidize the CS NP through the displacement of the spinel/rocksalt interface.24 However, even if EEL analysis suggests a large transformation into the spinel structure, the XRD patterns of annealed NPs still show the coexistence of both phases. The large amount of divalent cobalt atoms in the iron oxide structure and their high reduction potential (E0Fe(III)/Fe(II) = 0.77 V and E0Co(III)/Co(II) = 1.82 V) is crucial to avoid the complete formation of the spinel phase. As it has been demonstrated by XMCD spectra the annealing process permits the oxidation of the Fe2+ to Fe3+ and a cation rearrangement from octahedral to tetrahedral positions, meanwhile cobalt atoms remain unaltered in their divalent form in octahedral cavities.30,31 Therefore, taking into account the different cation valence and site occupancy of the rock-salt and inverse spinel structures, the formation of the latter phase alone maintaining the initial stoichiometry is not permitted. Indeed, while rock-salt phase is formed by an equivalent number of divalent cations and oxygen atoms (all octahedral sites occupied), in the inverse spinel only half of octahedral and 1/8 of tetrahedral sites are filled (for a complete inverse spinel 1/4 of octahedral positions are occupied by divalent ions while trivalent ions are located in 1/4 of octahedral and 1/8 of tetrahedral holes). Assuming a complete Fe2+ oxidation in a Co0.4Fe0.6O stoichiometry the amount of divalent cobalt ions (2/5 of total cations) is too large to stabilize the spinel structure (only able to contain up to 1/3 of divalent cations); therefore, the NP is forced to release this species in order to transform into the spinel structure.57 Indeed, this effect is demonstrated by EEL spectra which show the cobalt amount reduces when passing from M0 to M1. Moreover, the progressive growth of the oxidized phase occurs by the displacement of the rock-salt/spinel interface along the (200)RS//(400)S and (220)RS//(440)S planes which maintain an assemblage similar to M0, with a non-defective structure along the whole NP. However, with the impossibility of creating a stoichiometric spinel structure, the constrained oxidation of the pure spinel, e.g., planes (220)S, occurs 11 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 12 of 30

through the concomitant increase of the elastic strain, forming dislocations that move and assemble as the reaction progresses. The constrained oxidation process has been supported by the resemblance of crystal structures before and after the annealing process. Indeed, while the initial passivation of the M0 NPs leads to a continuous, well-defined shell structure, the subsequent oxidation occurs in a nonhomogenous manner into the NPs, forming small subdomains with the presence of a high number of defects and antiphase boundaries. The presence of this large number of dislocations creates a high number of crystalline boundaries bringing about the formation of a mosaic texture, in a similar way as it was proposed for the topotaxial oxidation of Ni(OH)2 crystals.50 The limited structural coherence of this defective, non-continuous spinel crystal structure with respect to XRD measurements, strongly affects the observed diffraction profiles, which do not resemble those expected for a pure cubic spinel structure. Consequently, it can be concluded that the oxidative growth of the spinel phase occurs through a topotaxial process involving the formation of a mosaic texture of small spinel subdomains into the well-defined rock-salt structure along the NP. The magnetic properties of the samples agree well with the proposed mechanism. The increase of TB and MS are direct evidences of the increase of the FiM volume, and thus of the spinel phase. Moreover, the annealed samples do not display the characteristic ordering transition at TN of the AFM core, which in fact could be related to the strong reduction in size below the correlation length (~ 2 nm) of the AFM phase and/or its complete oxidation to the spinel phase. The formation of antiphase boundaries after the growth of the spinel phase over the rock-salt structure is supported by the FC hysteresis loops of annealed samples. In fact, first of all the magnetizations after the annealing process do not still reach the values expected for 10 nm cobalt ferrite NP (88 and 80 emu/g at 2.5 and 300 K, respectively);7 second, if on one hand the reduction of HC is expected due to the loss of the extra magnetic anisotropy associated to the exchange coupling with the AFM phase, on the other hand large values of HE are still present. Wetterskog et al.24 demonstrated that the presence of antiphase boundaries created during the 12 ACS Paragon Plus Environment

Page 13 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


spinel growth in FeO/Fe3O4 NPs gives rise to exchange bias even for completely oxidized NPs. At the antiphase boundaries, indeed, the cations in the FiM cobalt ferrite structures have a different local strcutral ordering with rescet the corresponding bulk or highly crystalline materials.41 The cation disorder at the boundary results in changes of the electronic and magnetic interactions across the antiphase boundaries interface. Therefore depending on the type of antiphase boundaries several possible combinations of different super-exchange pathway can be found, possibly giving raise to canted spin structures and exchange coupled regions into the annealed NPs.42 These effects can explain both the reduced magnetic moment of the annealed samples and the presence of exchange bias properties. The appearance of exchange bias properties in the annealed samples makes these nanosystems a promising material to build up a novel class of free rare-earth magnets, as previously proposed in the literature.25,58,59 The performance of a material as permanent magnet is normally quantified by the socalled maximum energy product, BHmax.60 We previously demonstrated that the FC procedure causes, for CS NPs similar to those here investigated, a relative increase of BHmax, BHmaxFC/BHmaxZFC of 7 times.25 In this work we show that this value can be further doubled through the annealing process. Importantly, unlike conventional AFM/FM systems, where the low TN of the AFM material limits the exploitation of exchange bias to temperature lower than 300K, in this case it can be easily made effective above room temperature by simply increasing the NPs size.

CONCLUSIONS In conclusion, we have presented the solvent-mediated annealing process of cobalt-doped iron oxide CS NPs. The method permits heat treatments of ensemble of NPs at high temperatures (i.e., 300 ºC) avoiding interparticle aggregation, as usually observed in standard annealing methods. Interestingly, it 13 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 14 of 30

has been found that the rock-salt to spinel oxidative phase transformation occurs via topotaxial growth over the (200)RS//(400)S and (220)RS//(440)S planes shared between the rock-salt and spinel structures. This chemical transformation depends on the amount of divalent cobalt atoms present in the oxide structures. Within this respect, the solvent-mediated annealing process permits the release of a small amount of divalent cations allowing the stoichiometry rearrangement required to form the spinel phase, whose growth occurs through a topotaxial process through the formation of a mosaic texture of small spinel subdomains into the well-defined rock-salt structure along the NPs. Furthermore, annealed samples display exchange bias arising from the formation of spinel subdomains and antiphase boundaries. The combination of these two effects together with the increase of the magnetization leads to a significant increase of the energy stored in the material, suggesting solvent-mediated annealing process can be a powerful strategy to improve the performance of RE- free permanent magnets, which is currently a largely investigated research area.

EXPERIMENTAL SECTION The synthesis was carried out using standard airless procedures and commercially available reagents: 1-octadecene (ODE, 90%), docosane (DCE, 99%), ethanol (EtOH, >99.8%), hexane (Hx, >95%), oleic acid (OA, 90%), sodium oleate (NaOl, >97.0%), iron(III) chloride hexahydrate (FeCl3·6H2O, >98%), cobalt(II) chloride hexahydrate (CoCl2·6H2O, >98%). All starting materials were purchased from Sigma-Aldrich, except sodium oleate that was acquired from TCI America, and used without further purification. Monodisperse spherical NPs were synthesized through thermal decomposition of mixed metal-oleate complex in high-boiling solvent containing oleic acid as stabilizing surfactant. The mixed metal (Co2+Fe3+)-oleate was prepared mixing 4 mmol of FeCl3·6H2O, 2 mmol of CoCl2·6H2O, 16 mmol of Sodium Oleate (NaOl), 10 mL of H2O, 10 mL of Ethanol and 20 mL of Hexane; then, the 14 ACS Paragon Plus Environment

Page 15 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


mixture was refluxed for 4 h. After removal of the aqueous phase, the organic one was heated to 90 °C under vacuum to remove residual hexane and ethanol. To synthesize 11 nm iron-cobalt oxide NPs, 1.5 g of the mixed metal-oleate complex and 0.15 g of oleic acid were dissolved in 10 g of 1-octadecene (ODE) and magnetically stirred for 1 h under N2 flow. The mixture was heated up to 320 °C at a rate of 3 °C/min and maintained at this temperature for 2 h. Finally, NPs were washed by several cycles of coagulation with ethanol and 2-propanol, centrifugation at 5000 rpm, disposal of supernatant solution and re-dispersion in hexane. About a quarter of obtained NPs were further oxidized by dispersing 50 mg of NPs in 20 g of ODE and heating to 300°C for 5 minutes, bubbling air in the mixture.

Low-resolution transmission electron microscopy (TEM) images were obtained using a Philips CM12 microscope with a LaB6 filament operated at 100 kV. High resolution TEM (HRTEM) and electron energy loss (EEL) spectra were acquired at 200 kV on a JEOL JEM-2200FS equipped with a Ω filter. The NPs were dispersed in hexane and then placed drop wise onto a carbon supported grid. The particles size histograms were obtained by averaging manually measured diameters of more than 300 particles from TEM images. The acquisition conditions of the EEL spectrum image maps are described elsewhere.25 Geometrical phase analysis (GPA) was done using the FRWRtools plugin61 for Digital Micrograph (Gatan, Inc.).The simulations of electron diffractions where done in the multislice approximation using the xHREM™ software (HREM Research Inc.). The determination of cobalt and iron concentration in the samples was performed using a Rigaku ZSX Primus II X-ray fluorescence spectrometer (XRF). The microstructure of the NPs was investigated by X-ray powder diffraction (XRD) using a Bruker New D8 ADVANCE ECO diffractometer with Cu Kα radiation. The measurements were carried out in the range 25-70°, with a step size of 0.03° and a collection time of 1.5 s. Quantitative analysis of the XRD data was performed with a full pattern fitting procedure based on the fundamental parameter approach (Rietveld method) using MAUD software.62 XMCD 15 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 16 of 30

measurements were performed on dried CS nanoparticles spread onto carbon tape at the CiPo beamline of the ELETTRA synchrotron (Trieste, Italy). XMCD spectra were recorded at the Fe L2,3 edges using total electron yield (TEY) mode at 10 K. The XMCD signal was normalized by the area of the x-ray absorption spectra after correcting for the background. The magnetic properties of the NPs were measured on tightly packed powdered samples using vibrating sample mode magnetometer with 120 kOe (MagLab VSM12T-Oxford) and 90 kOe (VSM, Quantum Design PPMS) maximum field. Magnetization versus temperature measurements were performed in zero-field cooled (ZFC) and field cooled (FC) conditions with 50 Oe or 30 kOe probe fields. Hysteresis loops were measured in ZFC and FC conditions after cooling from RT to 10 K with a 120 kOe applied field. Magnetization data are reported for the amount of inorganic material as determined by ICP analysis.

ASSOCIATED CONTENT Supporting Information. HRTEM image and its respective indexed FFT image for sample M0. Simulated transmission electron diffractions for rock-salt and spinel phases. ZFC and FC hysteresis loops at 10 K for M0, M1, and M2 samples. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected]; [email protected] Author Contributions EL, ALO, CJF and CS conceived the idea. EL and ALO synthesized the nanoparticles and performed 16 ACS Paragon Plus Environment

Page 17 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


the measurements. GB performed electron microscopy measurements. CJF, EL, AL and CS performed and analysed the XMCD measurements. ALO, GB and CS wrote the manuscript. All authors contributed to revise the manuscript. Note The authors declare no competing financial interest.

ACKNOWLEDGMENTS This work is supported by the EU-FP7 through NANOPYME Project (No. 310516). ALO acknowledges the Juan de la Cierva Program (MINECO IJCI-2014-21530). The authors thank the ELETTRA Lightsource for experiments at the Circular Polarization beamline, and Nicola Zema, Stefano Turchini and Daniele Catone, for the supporting during the measurements.

17 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 18 of 30

REFERENCES

(1)

Lee, J.-H. H.; Jang, J.-T. T.; Choi, J.-S. S.; Moon, S. H.; Noh, S.-H. H.; Kim, J.-G. G. J.-W. W.; Kim, J.-G. G. J.-W. W.; Kim, I.-S. S.; Park, K. I.; Cheon, J. W. Exchange-Coupled Magnetic Nanoparticles for Efficient Heat Induction. Nat. Nanotechnol. 2011, 6, 418–422.

(2)

Pankhurst, Q. A.; Connolly, J.; Jones, S. K.; Dobson, J. Applications of Magnetic Nanoparticles in Biomedicine. J. Phys. D. Appl. Phys. 2003, 36, R167–R181.

(3)

Dalpozzo, R. Magnetic Nanoparticle Supports for Asymmetric Catalysts. Green Chem. 2015, 17, 3671–3686.

(4)

Kainz, Q. M.; Reiser, O. Polymer- and Dendrimer-Coated Magnetic Nanoparticles as Versatile Supports for Catalysts, Scavengers, and Reagents. Acc. Chem. Res. 2014, 47, 667–677.

(5)

Sun, S. Monodisperse FePt Nanoparticles and Ferromagnetic FePt Nanocrystal Superlattices. Science. 2000, 287, 1989–1992.

(6)

López-Ortega, A.; Estrader, M.; Salazar-Alvarez, G.; Roca, A. G.; Nogués, J. Applications of Exchange Coupled Bi-Magnetic Hard/soft and Soft/hard Magnetic Core/shell Nanoparticles. Phys. Rep. 2015, 553, 1–32.

(7)

López-Ortega, A.; Lottini, E.; de Julián Fernández, C.; Sangregorio, C. Exploring the Magnetic Properties of Cobalt-Ferrite Nanoparticles for the Development of a Rare-Earth-Free Permanent Magnet. Chem. Mater. 2015, 27, 4048–4056.

(8)

Willard, M. A.; Kurihara, L. K.; Carpenter, E. E.; Calvin, S.; Harris, V. G. Chemically Prepared Magnetic Nanoparticles. International Materials Reviews. 2004, 125–170.

(9)

Cozzoli, P. D. Advanced Wet-Chemical Synthetic Approaches to Inorganic Nanostructures; Cozzoli, P. D., Ed.; Transworld Research Network, 2008.

(10)

Juhin, A.; López-Ortega, A.; Sikora, M.; Carvallo, C.; Estrader, M.; Estradé, S.; Peiró, F.; Baró, M. D.; Sainctavit, P.; Glatzel, P.; Nogués, J. Direct Evidence for an Interdiffused Intermediate Layer in Bi-Magnetic Core-Shell Nanoparticles. Nanoscale 2014, 6, 11911–11920.

(11)

López-Ortega, A.; Roca, A. G.; Torruella, P.; Petrecca, M.; Estradé, S.; Peiró, F.; Puntes, V.; Nogués, J. Galvanic Replacement onto Complex Metal-Oxide Nanoparticles: Impact of Water or Other Oxidizers in the Formation of Either Fully Dense Onion-like or Multicomponent Hollow MnOx /FeOx Structures. Chem. Mater. 2016, 28, 8025–8031.

(12)

Pellicer, E.; Cabo, M.; López-Ortega, A.; Estrader, M.; Yedra, L.; Estradé, S.; Peiró, F.; Saghi, Z.; Midgley, P.; Rossinyol, E.; Golosovsky, I. V.; Mayoral, A.; Prades, J. D.; Suriñach, S.; Baró, M. D.; Sort, J.; Nogués, J. Controlled 3D-Coating of the Pores of Highly Ordered Mesoporous Antiferromagnetic Co3O4 Replicas with Ferrimagnetic FexCo3−xO4 Nanolayers. Nanoscale 2013, 5, 5561.

(13)

Park, J.; Joo, J.; Kwon, S. G.; Jang, Y.; Hyeon, T. Synthesis of Monodisperse Spherical Nanocrystals. Angew. Chem. Int. Ed. Engl. 2007, 46, 4630–4660.

(14)

Peng, S.; Wang, C.; Xie, J.; Sun, S. Synthesis and Stabilization of Monodisperse Fe Nanoparticles. J. Am. Chem. Soc. 2006, 128, 10676–10677. 18 ACS Paragon Plus Environment

Page 19 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(15)

Ullah, R.; Dutta, J. Photocatalytic Degradation of Organic Dyes with Manganese-Doped ZnO Nanoparticles. J. Hazard. Mater. 2008, 156, 194–200.

(16)

Park, J.; An, K.; Hwang, Y.; Park, J.-G.; Noh, H.-J.; Kim, J.-Y.; Park, J.-H.; Hwang, N.-M.; Hyeon, T. Ultra-Large-Scale Syntheses of Monodisperse Nanocrystals. Nat. Mater. 2004, 3, 891– 895.

(17)

Shin, J.; Anisur, R. M.; Ko, M. K.; Im, G. H.; Lee, J. H.; Lee, I. S. Hollow Manganese Oxide Nanoparticles as Multifunctional Agents for Magnetic Resonance Imaging and Drug Delivery. Angew. Chem. Int. Ed. 2009, 48, 321–324.

(18)

Zhang, Y.; Zhu, J.; Song, X.; Zhong, X. Controlling the Synthesis of CoO Nanocrystals with Various Morphologies. J. Phys. Chem. C 2008, 112, 5322–5327.

(19)

Kavich, D. W.; Dickerson, J.; Mahajan, S.; Hasan, S.; Park, J. H. Exchange Bias of Singly Inverted FeO/Fe3O4 Core-Shell Nanocrystals. Physical Review B. 2008, 174414.

(20)

Estrader, M.; López-Ortega, A.; Golosovsky, I. V; Estradé, S.; Roca, A. G.; Salazar-Alvarez, G.; López-Conesa, L.; Tobia, D.; Winkler, E.; Ardisson, J. D.; Macedo, W. A. A.; Morphis, A.; Vasilakaki, M.; Trohidou, K. N.; Gukasov, A.; Mirebeau, I.; Makarova, O. L.; Zysler, R. D.; Peiró, F.; Baró, M. D.; Bergström, L.; Nogués, J. Origin of the Large Dispersion of Magnetic Properties in Nanostructured Oxides: FexO/Fe3O4 Nanoparticles as a Case Study. Nanoscale 2015, 7, 3002–3015.

(21)

Bronstein, L. M.; Huang, X.; Retrum, J.; Schmucker, A.; Pink, M.; Stein, B. D.; Dragnea, B. Influence of Iron Oleate Complex Structure on Iron Oxide Nanoparticle Formation. Chem. Mater. 2007, 19, 3624–3632.

(22)

Sun, X. L.; Huls, N. F.; Sigdel, A.; Sun, S. H. Tuning Exchange Bias in Core/Shell FeO/Fe3O4 Nanoparticles. Nano Lett. 2012, 12, 246–251.

(23)

Gomes, J. de A.; Sousa, M. H.; Tourinho, F. A.; Aquino, R.; da Silva, G. J.; Depeyrot, J.; Dubois, E.; Perzynski, R. Synthesis of Core−Shell Ferrite Nanoparticles for Ferrofluids: Chemical and Magnetic Analysis. J. Phys. Chem. C 2008, 112, 6220–6227.

(24)

Wetterskog, E.; Tai, C.-W. W.; Grins, J.; Bergström, L.; Salazar-Alvarez, G. Anomalous Magnetic Properties of Nanoparticles Arising from Defect Structures: Topotaxial Oxidation of Fe 1– X O|Fe 3−δ O 4 Core|Shell Nanocubes to Single-Phase Particles. ACS Nano 2013, 7, 7132–7144.

(25)

Lottini, E.; López-Ortega, A.; Bertoni, G.; Turner, S.; Meledina, M.; Tendeloo, G. Van; de Julián Fernández, C.; Sangregorio, C. Strongly Exchange Coupled Core|shell Nanoparticles with High Magnetic Anisotropy: A Strategy towards Rare Earth -Free Permanent Magnets. Chem. Mater. 2016, 28, 4214–4222.

(26)

Quesada, A.; Granados-Miralles, C.; López-Ortega, A.; Erokhin, S.; Lottini, E.; Pedrosa, J.; Bollero, A.; Aragón, A. M.; Rubio-Marcos, F.; Stingaciu, M.; Bertoni, G.; de Julián Fernández, C.; Sangregorio, C.; Fernández, J. F.; Berkov, D.; Christensen, M. Energy Product Enhancement in Imperfectly Exchange-Coupled Nanocomposite Magnets. Adv. Electron. Mater. 2016, 2, 1500365.

(27)

Baaziz, W.; Pichon, B. P.; Lefevre, C.; Ulhaq-Bouillet, C.; Greneche, J.-M.; Toumi, M.; Mhiri, T.; Bégin-Colin, S. High Exchange Bias in Fe 3−δ O 4 @CoO Core Shell Nanoparticles Synthesized by a One-Pot Seed-Mediated Growth Method. J. Phys. Chem. C 2013, 117, 11436– 11443. 19 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 20 of 30

(28)

Fantechi, E.; Campo, G.; Carta, D.; Corrias, A.; de Julián Fernández, C.; Gatteschi, D.; Innocenti, C.; Pineider, F.; Rugi, F.; Sangregorio, C. Exploring the Effect of Co Doping in Fine Maghemite Nanoparticles. J. Phys. Chem. C 2012, 116, 8261–8270.

(29)

Yedra, L.; Xuriguera, E.; Estrader, M.; López-Ortega, A.; Baró, M. D.; Nogués, J.; Roldan, M.; Varela, M.; Estradé, S.; Peiró, F. Oxide Wizard: An EELS Application to Characterize the White Lines of Transition Metal Edges. Microsc. Microanal. 2014, 20, 698–705.

(30)

Oh, S. H.; Shin, R. H.; Lefèvre, C.; Thomasson, A.; Roulland, F.; Shin, Y.; Kim, D.-H.; Kim, J.Y.; Demchenko, A.; Leuvrey, C.; Mény, C.; Jo, W.; Viart, N. Incorporation of Cobalt Ions into Magnetoelectric Gallium Ferrite Epitaxial Films: Tuning of Conductivity and Magnetization. RSC Adv. 2015, 5, 34265–34271.

(31)

Hochepied, J. .; Sainctavit, P.; Pileni, M. . X-Ray Absorption Spectra and X-Ray Magnetic Circular Dichroism Studies at Fe and Co L2,3 Edges of Mixed Cobalt–zinc Ferrite Nanoparticles: Cationic Repartition, Magnetic Structure and Hysteresis Cycles. J. Magn. Magn. Mater. 2001, 231, 315–322.

(32)

López-Ortega, A.; Estrader, M.; Salazar-Alvarez, G.; Estradé, S.; Golosovsky, I. V.; Dumas, R. K.; Keavney, D. J.; Vasilakaki, M.; Trohidou, K. N.; Sort, J.; Peiró, F.; Suriñach, S.; Baró, M. D.; Nogués, J. Strongly Exchange Coupled Inverse Ferrimagnetic Soft/hard, MnxFe(3-x)O4/FexMn(3x)O4, Core/shell Heterostructured Nanoparticles. Nanoscale 2012, 4, 5138–5147.

(33)

Li, L.; Gan, Z.; McCartney, M. R.; Liang, H.; Yu, H.; Gao, Y.; Wang, J.; Smith, D. J. Atomic Configurations at InAs Partial Dislocation Cores Associated with Z-Shape Faulted Dipoles. Sci. Rep. 2013, 3, 3229.

(34)

Hÿtch, M. J.; Snoeck, E.; Kilaas, R. Quantitative Measurement of Displacement and Strain Fields from HREM Micrographs. Ultramicroscopy 1998, 74, 131–146.

(35)

Skumryev, V.; Stoyanov, S.; Zhang, Y.; Hadjipanayis, G.; Givord, D.; Nogués, J. Beating the Superparamagnetic Limit with Exchange Bias. Nature 2003, 423, 850–853.

(36)

Nogués, J.; Sort, J.; Langlais, V.; Skumryev, V.; Suriñach, S.; Muñoz, J. S.; Baró, M. D. Exchange Bias in Nanostructures. Phys. Rep. 2005, 422, 65–117.

(37)

Vasilakaki, M.; Trohidou, K. N.; Nogués, J. Enhanced Magnetic Properties in AntiferromagneticCore/ferrimagnetic-Shell Nanoparticles. Sci. Rep. 2015, 5, 9609.

(38)

De Toro, J. A.; Marques, D. P.; Muñiz, P.; Skumryev, V.; Sort, J.; Givord, D.; Nogués, J. High Temperature Magnetic Stabilization of Cobalt Nanoparticles by an Antiferromagnetic Proximity Effect. Phys. Rev. Lett. 2015, 115, 57201.

(39)

Vázquez-Vázquez, C.; López-Quintela, M. A.; Buján-Núñez, M. C.; Rivas, J. Finite Size and Surface Effects on the Magnetic Properties of Cobalt Ferrite Nanoparticles. J. Nanoparticle Res. 2011, 13, 1663–1676.

(40)

Chesnel, K.; Trevino, M.; Cai, Y.; Hancock, J. M.; Smith, S. J.; Harrison, R. G. Particle Size Effects on the Magnetic Behaviour of 5 to 11 Nm Fe 3 O 4 Nanoparticles Coated with Oleic Acid. J. Phys. Conf. Ser. 2014, 521, 12004.

(41)

Arora, S. K.; Sofin, R. G. S.; Nolan, A.; Shvets, I. V. Antiphase Boundaries Induced Exchange Coupling in Epitaxial Fe3O4 Thin Films. J. Magn. Magn. Mater. 2005, 286, 463–467.

(42)

Margulies, D. T.; Parker, F. T.; Spada, F. E.; Goldman, R. S.; Li, J.; Sinclair, R.; Berkowitz, A. E. Anomalous Moment and Anisotropy Behavior in Fe3O4 Films. Phys. Rev. B 1996, 53, 9175– 9187. 20 ACS Paragon Plus Environment

Page 21 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(43)

Bean, C. P.; Livingston, J. D. Superparamagnetism. J. Appl. Phys. 1959, 30, S120.

(44)

Mallick, P.; Mishra, N. C. Evolution of Structure, Microstructure, Electrical and Magnetic Properties of Nickel Oxide (NiO) with Transition Metal Ion Doping. American Journal of Materials Science. Scientific & Academic Publishing 2012, pp 66–71.

(45)

Maruyama, T.; Ueda, M. Void Formation Induced by the Divergence of the Diffusive Ionic Fluxes in Metal Oxides Under Chemical Potential Gradients. J. Korean Ceram. Soc. 2009, 47, 8–18.

(46)

Martin, M. Impurity Diffusion of Iron in Cobalt Oxide. Solid State Ionics 1988, 28–30, 1230– 1234.

(47)

Backhaus-Ricoult, M.; Dieckmann, R. Defects and Cation Diffusion in Magnetite (VII): Diffusion Controlled Formation of Magnetite During Reactions in the Iron-Oxygen System. Berichte der Bunsengesellschaft für Phys. Chemie 1986, 90, 690–698.

(48)

Bhadeshia, H. K. D. H. Diffusional Formation of Ferrite in Iron and Its Alloys. Scr. Metall. 1987, 21, 1605–1609.

(49)

Summerfelt, S. R.; Carter, C. B. The Movement of the Spinel-NiO Interface in Thin Films. Ultramicroscopy 1989, 30, 150–156.

(50)

Figlarz, M. Topotaxy, Nucleation and Growth. Solid State Ionics 1990, 43, 143–170.

(51)

Salazar-Alvarez, G.; Sort, J.; Suriñach, S.; Baró, M. D.; Nogués, J. Synthesis and SizeDependent Exchange Bias in Inverted Core-Shell MnO|Mn3O4 Nanoparticles. J. Am. Chem. Soc. 2007, 129, 9102–9108.

(52)

Bean, J. C. GexSi1−x/Si Strained-Layer Superlattice Grown by Molecular Beam Epitaxy. J. Vac. Sci. Technol. A Vacuum, Surfaces, Film. 1984, 2, 436.

(53)

Fung, K. K.; Qin, B.; Zhang, X. X. Passivation of α-Fe Nanoparticle by Epitaxial γ-Fe2O3 Shell. Mater. Sci. Eng. A 2000, 286, 135–138.

(54)

Yu, X. L.; Jiang, Z. Y.; Yang, D. J.; Wei, D. Bin; Yang, Q. Precipitation Behavior of Magnetite in Oxide Scale during Cooling of Microalloyed Low Carbon Steel. Adv. Mater. Res. 2012, 572, 249–254.

(55)

Sun, L.; Marrocchelli, D.; Yildiz, B. Edge Dislocation Slows down Oxide Ion Diffusion in Doped CeO2 by Segregation of Charged Defects. Nat. Commun. 2015, 6, 6294.

(56)

Tinkler, S.; Subramanian, R.; Dieckmann, R. Defects and Transport in the Solid Solution (Co,Fe)1 − ∆ O at 1200°C—II. Cation Tracer Diffusion and Electrical Conductivity. J. Phys. Chem. Solids 1994, 55, 273–286.

(57)

Hazen, R. M.; Jeanloz, R. Wüstite (Fe 1-X O): A Review of Its Defect Structure and Physical Properties. Rev. Geophys. Sp. Phys. 1984, 22, 37–46.

(58)

Sort, J.; Nogués, J.; Suriñach, S.; Muñoz, J. S.; Baró, M. D.; Chappel, E.; Dupont, F.; Chouteau, G. Coercivity and Squareness Enhancement in Ball-Milled Hard Magnetic–antiferromagnetic Composites. Appl. Phys. Lett. 2001, 79, 1142.

(59)

Sort, J.; Suriñach, S.; Muñoz, J. S.; Baró, M. D.; Nogués, J.; Chouteau, G.; Skumryev, V.; Hadjipanayis, G. C. Improving the Energy Product of Hard Magnetic Materials. Phys. Rev. B 2002, 65, 174420.

(60)

Coey, J. M. D. Magnetism and Magnetic Materials; Cambridge University Press: Cambridge, 2010. 21 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 22 of 30

(61)

https://www.physik.hu-berlin.de/en/sem/software/software_frwrtools (2016).

(62)

Lutterotti, L. MAUD program http://www.ing.unitn.it/~maud/ (2016).

TABLES

Table1. Structural parameters obtained from Rietveld evaluation of XRD patterns of as prepared and oxidized NPs. Rock-salt phase

Spinel phase Cell Parameter (nm)

Crystal Size (nm)

w%

Lattice mismatch (%)

Sample

Cell Parameter (nm)

Crystal Size (nm)

M0

0.422(1)

19(2)

33(2)

0.842(1)

5

67

0.1(1)

M1

0.419(1)

16(2)

20(2)

0.838(1)

7

80

0.0(1)

M2

0.420(1)

15(2)

18(2)

0.838(1)

7

82

0.1(1)

w%

Table2. Magnetic properties at LT of as prepare and oxidized NPs. MT ZFC Sample TB TN HC HE MS HC (K) (K) (kOe) (kOe) (emu/g) (kOe) 190 234 13.0 0 18.9 19.2 M0 220 -----15.5 0 37.8 17.4 M1 230 -----15.2 0 37.8 17 M2


FC HE (kOe) 8.7 3.4 3.2

MS (emu/g) 19 38 39

Page 23 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


FIGURES

Figure 1. LR-TEM (a-c) and HR-TEM (d-f) images of M0 (a,d); M1 (b,e) and M2 (c,f) CS NPs; the corresponding particle size histograms are reported in the insets of the LR-TEM mages.



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 24 of 30

Figure 2. Powder XRD patterns for M0, M1 and M2 samples (indexed planes are labeled with RS and S for rock-salt and spinel structure, respectively).


Page 25 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 3. (a) Iron, cobalt and oxygen elemental quantification for samples M0, M1, and M2 along the radius of a NP; (b) Radial distribution of metal atoms (Co + Fe, squares) and oxygen (O, circles) percentage for M0 (full symbols) and M1 (empty symbols); (c) Fe2+/Fe3+ EEL mapping for M0 and M1 samples, where green and red colors refers to Fe3+ and Fe2+, respectively.



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 26 of 30

Figure 4. (a) Cobalt and (b) iron L2,3 edges XMCD spectra for samples M0 (black curve) and M1 (red curve), respectively.


Page 27 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 5. HRTEM images from [001] oriented particles in samples M0 and M1, and their respective Bragg filtered maps for the (220)S, (200)RS/(400)S and (220)RS/(440)S diffraction spots; large red squares are enlargement of selected areas (small red squares) to highlight dislocation defects in (220)S map.



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 28 of 30

Figure 6. Phase maps (Φ) and amplitude maps (A) obtained from GPA analysis on selected g reflections. Phase is plotted in -π (black) +π (white) and amplitude from 0 (black) to maximum (white) values. Corresponding lattice deformation profiles for samples M0 and M1 calculated from the dilatation maps. A considerably higher deformation is found on (220)S planes rather than on (200)RS 28 ACS Paragon Plus Environment

Page 29 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


planes.

Figure 7. (a) 50 Oe ZFC-FC curves for samples M0, M1, and M2 (upper panel) and 3T FC curves for samples M0 and M1 (bottom panel). (b) Hysteresis loops at 10 K for M0, M1 and M2. 29 ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

TABLE OF CONTENTS GRAPHIC


Page 30 of 30

Topotaxial Phase Transformation in Cobalt Doped Iron Oxide Core

Recommend Documents