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Probing the Role of Dopant Oxidation State in the Magnetism of Diluted Magnetic Oxides using Fe-Doped InO and SnO Nanocrystals 2

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Natalie S Garnet, Vahid Ghodsi, Lisa N. Hutfluss, Penghui Yin, Manu Hegde, and Pavle V. Radovanovic J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.6b09480 • Publication Date (Web): 03 Jan 2017 Downloaded from http://pubs.acs.org on January 8, 2017

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The Journal of Physical Chemistry C 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.

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The Journal of Physical Chemistry

Probing the Role of Dopant Oxidation State in the Magnetism of Diluted Magnetic Oxides using FeDoped In2O3 and SnO2 Nanocrystals Natalie S. Garnet, Vahid Ghodsi, Lisa N. Hutfluss, Penghui Yin, Manu Hegde and Pavle V. Radovanovic* Department of Chemistry, University of Waterloo, 200 University Avenue West, Waterloo, ON N2L 3G1, Canada.

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ABSTRACT

Investigation of the origin of high-Curie temperature ferromagnetism in diluted magnetic oxides has become one of the focal points of research on solid-state magnetism. While several possible mechanisms have been proposed theoretically, broader experimental evidence is still lacking. Here we report a comparative study of the electronic structure and magnetic properties of colloidal Fe-doped In2O3 and SnO2 nanocrystals, as building blocks for grain boundary-rich diluted magnetic oxide films. The dopant ions in both nanocrystal host lattices are principally in 3+ oxidation state, with possibly a minor presence of Fe2+ in In2O3, and no conclusive evidence of the presence of Fe2+ in SnO2 nanocrystals. Subsequently, we found that Fe-doped In2O3 nanocrystalline films exhibit only minor ferromagnetic ordering (with the magnetic moment of less than ca. 0.1 µB/Fe) and decreasing saturation magnetization with increasing doping concentration at room temperature. The saturation magnetic moment of Fe-doped SnO2 nanocrystalline films is insignificant or below the detection limit. These results contrast previous findings for analogous Mn-doped nanocrystals, which contain mixed oxidation states (Mn2+ and Mn3+), and exhibit a robust ferromagnetism at room temperature. The correlation between the mixed dopant oxidation states and the observed magnetic properties implies that ferromagnetism in these systems is of a Stoner type, enabled by electron transfer between dopant ions and the local defect states arising from the grain boundaries within a nanocrystalline film. These results suggest the prospect of probing and manipulating ferromagnetism in non-magnetic oxides by simultaneous control of the transition metal dopant oxidation states and extended structural defects.


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INTRODUCTION Following the first report of room-temperature ferromagnetism in cobalt-doped TiO2,1 a flurry of research has been launched since the beginning of the century to understand the mechanism of long range dilute magnetic ordering, and expand the range of available diluted magnetic oxides (DMOs).2-5 In analogy to diluted magnetic semiconductors (DMSs),6 most notably Mn-doped GaAs,7-11 this class of materials has been seen as potential spin injectors into non-magnetic semiconductors,12 which can enable the use of electron spin for non-volatile information manipulation and storage, as the basis for spin-electronics (spintronics).3,13-15 Unlike GaAs-based diluted magnetic semiconductors, DMOs could operate at room temperature while exhibiting a high degree of transparency desirable for complex multifunctional devices.2,3 In the course of the studies of DMOs, the mechanism of dilute magnetic ordering has emerged as one of the more intriguing questions in solid-state chemistry and physics. Traditional descriptions of exchange interaction in magnetic oxides, including double exchange and superexchange,16 cannot generally account for high-Curie temperature (TC) long-range dilute magnetic ordering, since the doping concentrations are generally well below the necessary percolation threshold.

Analogously to III-V DMSs, carrier-mediated dopant exchange

interactions have initially been proposed,17-21 although in some cases no correlation has been found between magnetic ordering of dopant ions and charge carrier concentration.22 This discrepancy has prompted investigations into both extrinsic origin of ferromagnetism23-25 (i.e. dopant-related secondary phase formation) and alternative mechanisms of magnetic interactions in solid state.26-29 A number of studies of high-TC ferromagnetic DMO samples prepared under relevant conditions convincingly ruled out the presence of secondary phases.30 Furthermore, ferromagnetism has been detected even in high surface area oxide nanostructures prepared in the


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absence of transition metal dopants (“d0-ferromagnetism”),31-35 suggesting that there may be multiple mechanisms of dilute ferromagnetism, which could potentially coexist. Different models of dilute ferromagnetism in oxides have subsequently been proposed, with varying degrees of quantitative treatment.27,29,36,37 The common feature of the models that have gained substantial traction is the importance of structural defects in mediating ferromagnetic ordering.30,38-40 According to the bound magnetic polaron model, ferromagnetic ordering of dopant ions is mediated by shallow donor electrons, generated by structural defects, which form bound magnetic polarons.27 Overlapping of these polarons within the host lattice leads to the formation of a spin-split donor impurity band. High-TC ferromagnetism is enabled when empty majority- or minority-spin dopant 3d states lie at the Fermi level in the impurity band. This model is consistent with the behavior of kinetically stable dopant ions, such as Cr3+, in structurally defective transparent metal oxide films containing extended interfacial structural defects (i.e., grain boundaries).41 Another model relying on defects at the surfaces of nanostructures is the charge transfer (CT) ferromagnetism model.29 In this model, ferromagnetism is attained by spin-splitting of the defect-derived electronic band, enabled by the coexistence of dopant ions in two different oxidation states. For example, electron transfer from a dopant ion with the lower oxidation state to the local defect density of states may raise the Fermi level to a peak in the local density of states, leading to spin-splitting of the defect band states.

Unlike the bound magnetic polaron model, which involves the Heisenberg-type

exchange, CT ferromagnetism is of a Stoner type, and the role of dopant ions is not to carry a magnetic moment but to serve as a reservoir for charge transfer. This mechanism is also transferrable to non-magnetic oxides containing donor/acceptor surface states acting as a charge reservoir. The presence of interfacial structural defects is essential in this model as well, since


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the largest fraction of the magnetic moment corresponds to defect-rich areas (i.e., nanocrystal interfaces and grain boundaries) where the local defect density of states is the highest. We have previously demonstrated strong ferromagnetic ordering in Mn-doped In2O3 and SnO2 nanocrystalline films prepared from colloidal nanocrystals (NCs) as building blocks.42,43 It has been found that in both nanocrystal host lattices dopant ions coexist as Mn2+ and Mn3+, and the overall behavior of these systems is consistent with CT ferromagnetism. However, to more deeply investigate the possibility of the CT ferromagnetism model, suitable experiments need to be designed, including the investigations of a variety of dopant ions exhibiting a different degree of mixed valence states in a given host lattice. Fe3+ is isoelectronic with Mn2+, which allows for a direct comparison of the magnetic properties of the analogous Fe-doped and Mn-doped metal oxide host lattices. Moreover, Mn3+ has a higher tendency to be reduced to 2+ oxidation state compared to Fe3+ according to the standard aqueous reduction potentials (+1.51 V and +0.77 V for Mn3+ and Fe3+, respectively). In this work we report the synthesis and detailed investigation of iron-doped In2O3 (Fe:In2O3) and SnO2 (Fe:SnO2) colloidal NCs, and show that the majority of iron dopants exist in 3+ oxidation state, with possibly a minor fraction of Fe2+. The resulting magnetic properties directly examine one of the criteria for CT ferromagnetism and show the close inter-relationship between mixed oxidation state and observed magnetic moment. The largest net magnetic moment measured for Fe-doped In2O3 nanocrystalline films is well below 0.1 µB per dopant, in contrast with > 0.3 µB/dopant for the analogous Mn-doped In2O3 films. The results of this work confirm the possibility of CT ferromagnetism as an alternative mechanism of ferromagnetism in oxides, and demonstrate the ability to use redox potentials of transition metal ions to tune magnetic


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properties of DMOs. In that context, this study contributes to explaining “phantom” ferromagnetism in this class of materials.

EXPERIMENTAL SECTION Materials. All reagents and solvents are commercially available, and were used as received without further purification. Indium acetylacetonate (In(acac)3, 98 %) and tin(IV) chloride pentahydrate (SnCl4·5H2O, 98 %) were purchased from STREM Chemicals. Iron(III) chloride (FeCl3, 97 %) and iron(III) chloride hexahydrate (FeCl3·6H2O, 99.9 %) were purchased from Sigma-Aldrich and Baker Chemical Co., respectively. Oleylamine (70 %), oleic acid (90 %), trioctylphosphine oxide (TOPO, 90 %), dodecylamine (DDA, 98 %), 1,4-dioxane (99 %), ammonium hydroxide (NH4OH, 28.0-30.0 %), and hexane (HPLC grade) were purchased from Sigma-Aldrich. Toluene (99.98 %) was purchased from EMD Chemicals. Preparation of Fe-Doped In2O3 NCs. The synthesis and work-up methods were adapted from our previously reported work.41,42,44 In a 100 mL three-neck round bottom flask, 2 mmol of In(acac)3 and varied amounts (5 to 15 mol %) of FeCl3 were combined with 24 mmol of oleylamine. For cubic phase Fe-doped In2O3 NCs, the content was degassed and subsequently heated to 230 °C, after which the reaction mixture was refluxed under argon atmosphere with constant stirring for 1 h. For rhombohedral phase Fe-doped In2O3 NCs the reaction mixture was heated to 210 °C. The reaction mixture was then slowly cooled to room temperature and NCs were precipitated with ethanol. Obtained Fe:In2O3 NCs were rinsed by the addition of ethanol, followed by centrifugation at 3000 rpm for 5 minutes; this process was repeated for a total of 3 rinses. The NCs were then combined with an excess amount of TOPO and heated at 90 °C for 1 h with constant stirring, followed by precipitation with ethanol and centrifugation. The TOPO


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treatment was performed three times to ensure NC surfaces are free of surface-bound dopant ions and easily dispersed in nonpolar organic solvents.45 Use of iron acetylacetonate (Fe(acac)3) precursor also afforded Fe:In2O3 NCs. Treated NCs were suspended in hexane or toluene for analysis and measurements. All samples were handled in an identical fashion and kept under conditions free of contamination by magnetic materials and tools. Preparation of Fe-Doped SnO2 NCs. Fe-doped SnO2 NCs were synthesized by a sol-gel method.43 Briefly, 1.4 g of SnCl4·5H2O and varying amounts of FeCl3·6H2O (from 5 to 15 mol %) were added to 20 mL of deionized water and the mixture was stirred until precursors were dissolved. After cooling the reaction mixture in an ice bath for 15 mins, the particles were nucleated through the slow dropwise addition of concentrated NH4OH (30 %) until reaching pH 6. The reaction beaker was left for 3 h for the nucleated particles to settle, followed by triple washing with deionized water. A pipette-full of NH4OH was then added to the suspension in 15 min intervals until it became fully transparent. This suspension was then refluxed for 15 h at 90 °C. Upon cooling the suspension to room temperature, the resulting NCs were extracted by the addition of 1,4-dioxane, collected by centrifuging, and washed with ethanol 3 times. The precipitated NCs were resuspended in an excess amount of melted DDA and heated at 120 °C for 30 min resulting in a clear suspension. Such suspensions were cooled, precipitated and washed with ethanol. In the final step, DDA-capped NCs were treated with TOPO at 140 °C for 1 hour, and afterwards washed with ethanol 3 times. TOPO treatment was repeated 2 more times to ensure the removal of the dopant ions from the NC surface.45 Finally, NCs were suspended in a small amount of hexane for optical measurements and the preparation of nanocrystalline films. Structural Characterization. Powder X-ray diffraction patterns were collected with an INEL powder diffractometer. This instrument is equipped with a position-sensitive detector and


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monochromatic Cu Kα radiation (λ = 1.5418 Å). The obtained NC powders were dried, crushed, and loaded onto an aluminum holder. Transmission electron microscopy (TEM) imaging, energy dispersive X-ray spectroscopy (EDX) and electron energy loss spectroscopy (EELS) measurements were performed with a JEOL-2010F microscope operating at 200 kV. Specimens were prepared using standard copper grids with lacey Formvar/carbon support films purchased from Ted Pella Inc. The average nanocrystal sizes and size distributions were determined based on the measurements of at least 150 NCs using Gatan Digital Micrograph software. Spectroscopic Measurements and Analysis. Optical absorption spectra were collected with a Varian Cary 5000 UV-vis-NIR spectrophotometer. Absorption spectra were collected using 1 cm path-length quartz cuvettes for room-temperature measurements of colloidal NCs, and strain-free quartz substrates for low-temperature measurements of deposited NCs. Magnetic circular dichroism (MCD) spectroscopy measurements were performed in a Faraday configuration with a Jasco J-815 spectropolarimeter. The NCs deposited on strain-free quartz substrates were mounted in a variable magnetic field (0-7 T) magneto-optical cryostat (SM4000-8, Oxford Instruments) equipped with a variable temperature insert (1.5–300 K). X-ray absorption spectroscopy (XAS) measurements were performed at the Canadian Light Source (CLS). Iron Kedge X-ray absorption near-edge structure (XANES) and extended X-ray absorption fine structure (EXAFS) spectra were collected at the Hard X-ray MicroAnalysis (HXMA, 06ID-1) beamline. An iron foil was used to internally calibrate the energy of the Fe K-edge (7112 eV). The spectra were collected in a fluorescence mode using a 30-element germanium detector. Least-squares curve fitting of the EXAFS spectra was performed over the k range 2.9 – 11.6 Å-1, using hematite Fe2O3 crystal structure as the model in the FEFF 7.0 program (see Supporting Information for more detail).46


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Magnetization Measurements. For magnetization measurements of nanocrystalline films colloidal NCs were spin-coated on sapphire substrates previously sonicated in aqua regia, followed by mild annealing at 300 °C for 1 min. The spin coating process was repeated until the mass of the nanocrystalline film reached at least 1 mg. Magnetization data were generated using a Physical Property Measurement System (PPMS, Quantum Design) in ACMS mode. The magnetic hysteresis loops were recorded at 300 K and magnetic fields of 0-5000 Oe.

RESULTS AND DISCUSSION Typical XRD patterns of Fe-doped In2O3 and SnO2 NCs, prepared with different concentrations of Fe3+ precursor, are shown in the Figure 1a and b, respectively. Iron-doped In2O3 NCs have body-centered cubic (bcc) crystal structure characteristic for bixbyite-type In2O3 NCs. With increasing doping concentration, the XRD reflections shift to higher 2θ values, indicating shrinking of the NC host lattice due to the substitution of smaller Fe3+ for In3+ cations. Synthesis of Fe:In2O3 NCs at low temperatures (< 210 °C) and low Fe3+ precursor concentration (< ca. 5 %) results in the stabilization of corundum-type In2O3 NCs having rhombohedral (rh) crystal structure (Figure S1 in Supporting Information). Considering the size-phase relationship of In2O3,44,47-49 it can be inferred that NC growth is sufficiently hindered by the combination of iron dopant and low temperature conditions to maintain NC sizes below ca. 5 nm, and thus facilitate retention of the rh-In2O3 phase.44 With higher-temperature synthesis (ca. 230 oC and higher), all samples contain only bcc-In2O3 NCs, regardless of the doping concentration. This implies that higher temperatures promote NC growth beyond the critical size, despite possible effect of dopant ion on the NC growth.44,50 Interestingly, however, it was found that higher dopant concentrations (i.e., 10 % and 15 %) at lower temperatures resulted in bcc-phase NCs,


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further suggesting that the combination of low doping concentration and low temperature results in the stabilization of the rh-In2O3 NCs, while an increase in either of those parameters will lead to the formation of bcc-In2O3. The most stable phase of Fe2O3 is hematite, α-Fe2O3, which has a rhombohedral (corundum-type) structure.51 Thus, in analogy to Cr3+ and Mn3+ dopants,44 it would be anticipated that rh-In2O3 would be the favored host lattice for incorporating Fe3+ dopants due to the structural similarity between the two materials. Moreover, β-Fe2O3, which has body-centered cubic structure, is a metastable phase of Fe2O3, so it is unexpected that Fe3+ would have higher solubility in the bcc-In2O3 host lattice. However, it is possible that the stability of Fe2O3 polymorphs reverses when the material is reduced below the critical size, similarly to In2O3,47 explaining the observed trend in incorporation of Fe3+ dopants in In2O3. XRD patterns of Fe:SnO2 NCs (Figure 1b) agree well with the pattern of rutile SnO2 (cassiterite). The peaks are significantly broader relative to those for In2O3, suggesting much smaller NC sizes.

Figure 1. XRD patterns of (a) Fe:In2O3 and (b) Fe:SnO2 NCs synthesized with different dopant to host NC precursor ratio, as indicated in the graphs. The sticks in (a) and (b) represent the bulk patterns of bcc-In2O3 and SnO2, respectively. The insets show the corresponding unit cells,


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where pink (a) and violet (b) spheres are metal sites, and red spheres are oxygen sites (oxygen ions coordinating In3+ d-sites in (a) are removed for clarity). Typical TEM images of Fe:bcc-In2O3 NCs synthesized at 230 °C with 5 % Fe3+ precursor concentration are shown in Figure 2a and b. The obtained NCs have nearly ideal spherical shape, an average diameter of 8.6 nm, and the doping concentration of 1.8 %. The lattice spacing of 2.92 Å corresponds to the {222} lattice plane (Figure 2b). Nanocrystals synthesized at 210 °C with 5 % Fe3+ precursor in the reaction mixture are shown in Figure 2c and d. These NCs are significantly smaller than NCs shown in Figure 2a, as expected, but retain quasi-spherical shape and high crystallinity. The lattice spacing determined for the NC shown in Figure 2d (2.87 Å) corresponds to the {104} plane of rh-In2O3, with a particle diameter of 3.8 nm, much smaller than the critical size for phase transformation, demonstrating the presence of rh- and bcc-In2O3 NCs in the low and high temperature regimes, respectively. At 5 % dopant precursor concentration and increasing reaction temperature from 210 °C to 230 °C, the average NC size increases from ca. 3 to 9.5 nm (Figure S2 in Supporting Information); this corresponds to the anticipated change in phase from rh- to bcc-In2O3. TEM imaging additionally revealed the formation of relatively large nanoflower structures of Fe:bcc-In2O3, when high concentrations of FeCl3 dopant precursor were used (Figure S3 in Supporting Information). Similar NC selfassembly was observed in the case of other doped In2O3 NCs,42,52 which was associated with dipole moment formation due to unequal charging of the NC surfaces by anions in solution, limited surface protection, and overall surface stabilization.53 Fe:SnO2 NCs have even smaller average size than rh-In2O3 NCs (Figure 2e), but are also highly crystalline, as evidenced from the high-resolution TEM image in Figure 2f. Hydrothermal synthesis results in the formation of sub-


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3 nm diameter Fe:SnO2 nanowires, likely achieved by oriented attachment of pre-formed NCs (Figure S4 in Supporting Information).43

Figure 2. TEM images of (a,b) 1.8 % Fe:bcc-In2O3 NCs, (c,d) 1.3 % Fe:rh-In2O3 NCs, and (e,f) 2.5 % Fe:SnO2 NCs. (b), (d), and (f) show lattice-resolved TEM images and the corresponding lattice spacing of individual NCs. The uniformity of dopant distribution throughout the samples and within individual NCs was probed using elemental mapping and line scan analysis. Typical STEM image and the corresponding EDX In and Fe elemental maps of 1.8 % Fe:In2O3 NCs are shown in Figure 3a-c. The spatial distribution of Fe mirrors that of In without any Fe-rich regions that would indicate the obvious presence of secondary phases. However, this observation does not confirm that


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dopant ions are distributed uniformly within individual NCs. To probe the distribution of Fe within single NCs we performed EELS line scan on selected individual NCs (Figure 3d). EELS has a lower detection limit than EDX, allowing for analytical measurements on single NCs without their drift or degradation under electron beam in the microscope. The line scan profiles of Fe and In are very well matched (Figure 3e), indicating that Fe dopant ions are randomly distributed throughout the NC, and are not segregated on NC surface. Similar results were obtained for higher doping concentration. The spatial distribution of Fe and In in 7.9 % Fe:In2O3 NCs is nearly identical throughout the sample (Figure 3f-j), as well as within individual nanoflowers (Figure 3i and j, and Figure S5 in Supporting Information). The results in Figure 3 demonstrate that the distribution of Fe dopants is homogeneous at both the ensemble and single NC levels.


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Figure 3. (a) STEM image and the corresponding (b) In and (c) Fe EDX elemental maps of 1.8 % Fe-doped In2O3 NCs. The designated area in (a) indicates the region analyzed by EDX. (d) STEM image and (e) corresponding In (green) and Fe (yellow) EELS line scans of a single NC in (a). Line in (d) indicates the location of the line scan. (f) STEM image and the corresponding (g) In and (h) Fe EDX elemental maps of 7.9 % Fe-doped In2O3 NCs. (i) STEM image and (j) corresponding In (green) and Fe (yellow) EDX line scans of a single nanoflower in (f). Line in (i) indicates the location of the line scan. To elucidate the oxidation state of iron dopant ions in In2O3 and SnO2 NCs we performed XAS measurements at the Fe K-edge. XANES spectra of Fe-doped In2O3 and SnO2 NCs are shown together with those for Fe3+ and Fe2+ reference compounds (hematite and olivine, respectively) in Figure 4a (see also Figure S6 in Supporting Information). Qualitative inspection of the spectra in Figure 4a reveals that all samples and Fe3+ reference have very similar positions of the spectral edge and nearly identical transitions, suggesting a dominant, if not exclusive, presence of Fe3+ in both In2O3 and SnO2, consistent with previous reports.54


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Figure 4. (a) Fe K-edge XANES spectra of typical 1.8 % Fe:bcc-In2O3, 1.3 % Fe:rh-In2O3, and 2.5 % Fe:SnO2 NCs, as indicated in the graph. The spectra of reference compounds are shown for comparison. (b) Pre-edge features of XANES spectra in (a). (c) Fitting of the pre-edge XANES spectra of typical Fe-doped In2O3 and SnO2 NCs to determine the pre-edge centroid position. (d) Dependence of the pre-edge centroid position on the Fe dopant oxidation state. In addition to the K-edge energy, pre-edge features in XANES spectra of transition metal absorbers are indicative of the local site symmetry and orbital occupancy. In the case of first row TM ions, the pre-edge feature is the result of 1s → 3d electronic transition.55 The 1s → 3d transition is formally dipole forbidden; however, symmetry-induced mixing of 3d and 4p states allows this transition to gain intensity. Alternatively, the pre-edge feature may also result from 1s → 3d quadrupole transitions. Figure 4b shows the pre-edge transitions for the XANES spectra in Figure 4a. The position and overall band shape of the pre-edge peaks for Fe doped in In2O3 and SnO2 NCs are very similar to those for the hematite reference, confirming that most of the dopant ions are in 3+ oxidation state. For the Fe K-edge in particular, analysis of the centroid of the pre-edge peaks has proven to be the most consistent way to quantify iron oxidation states and local site symmetry using XANES spectra.56 Fittings of the pre-edge peaks for Fe-doped bccIn2O3 (top) and SnO2 (bottom) NCs used for subsequent determination of the centroid position are shown in Figure 4c (see Figure S7 in Supporting Information for fitting the pre-edge peaks for other samples and reference compounds). The centroid, or intensity-weighted average of the pre-edge energy positions, is correlated with the oxidation state of the transition metal absorber, a higher energy centroid indicating a higher oxidation state. It should be noted that the coordination number of metal absorber also influences the pre-edge spectral feature and therefore should be considered when estimating the fractions of different redox states in the material. The


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Fe pre-K-edge centroid position for different Fe-doped NC samples and reference compounds is plotted as a function of iron oxidation state in Figure 4d. The solid line is connecting the points indicating the centroid positions for reference compounds, and as such serves as a calibration line. The points for all samples are clustered around the centroid position corresponding to hematite (Fe3+) reference, with the lowest estimated oxidation state for 1.8 % Fe:bcc-In2O3 NCs (ca. 2.9). These results suggest that the upper limit of the fraction of Fe2+ in these samples is ca. 10 %. This is in stark contrast with Mn-doped In2O3 and SnO2 NCs, which contain comparable fractions of Mn2+ and Mn3+ oxidation states.42 Quantitative analysis also confirmed that total pre-edge peak area for Fe:rh-In2O3 NC sample is somewhat greater than that for Fe:bcc-In2O3 NC samples, implying that the dopant sites are more distorted in the rh-In2O3 phase. Assuming that Fe3+ is substitutionally incorporated and thus in a six-coordinate environment in all samples, this observation may allude to the specific sites in which Fe3+ resides within the bcc-In2O3 NCs. The rh-In2O3 phase contains only one type of In3+ site, in which the oxygen atoms are arranged in a trigonal antiprism around In3+, and the symmetry of these sites is C3v.57-59 Within the bcc-In2O3 lattice there are two types of In3+ sites, the b-sites are of C3i symmetry, while the d-sites are highly distorted having C2 symmetry.60 Since Fe3+ must be incorporated into the C3v sites in the rh-In2O3 lattice and the pre-edge peak area indicates that this is the most disordered sample, it is reasonable to presume that Fe3+ incorporates into the more-symmetric b-sites in bcc-In2O3 NCs. If Fe3+ were incorporated into the d-sites, it would be expected that the pre-edge total area for Fe3+ in bcc-In2O3 NCs would be greater than that for Fe3+ in rh-In2O3 NCs, since the symmetry is lower. This result is indeed consistent with the results for other transition metal dopants including Mn2+/3+ and Cr3+.41,42,44


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Figure 5. (a,b) Fourier filtered EXAFS spectra (a) and pseudo-radial distribution functions (b) of Fe:In2O3 NCs having different doping concentrations (as indicated in the graph). (c,d) Fourier filtered EXAFS spectra (c) and pseudo-radial distribution functions (d) of Fe:SnO2 NCs having different doping concentrations (as indicated in the graph). Table 1. Structural Parameters for the First Shell (Fe-O) Obtained from Least-Squares Curve Fitting of Fe K-edge EXAFS Spectra of Fe-doped In2O3 and SnO2 Nanocrystals Sample

N

R (Å)a

σ2 (Å2) b

1.3% Fe:rh-In2O3

2.6

1.96

0.0027

1.8% Fe:bcc-In2O3

3.0

1.99

0.0024

7.9% Fe:bcc-In2O3

3.9

1.97

0.0026

2.5% Fe:SnO2

3.9

1.97

0.0016

13.5% Fe:SnO2

4.1

1.95

0.0017

a

R±0.02 Å; b Debye-Waller factor (σ2 ±0.002 Å2)


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To gain deeper insight into the local environment of Fe3+ dopants we also performed the analysis of the EXAFS spectra of Fe3+ in In2O3 and SnO2 NCs. Figure 5a and b shows k2weighted Fe K-edge Fourier-filtered EXAFS spectra and the corresponding pseudo-radial distribution functions, respectively, for the first coordination shell (Fe-O) for different Fe:bccIn2O3 NC samples. The EXAFS oscillation amplitude and the Fourier transform magnitude increase with increasing doping concentration, as expected. When the amount of Fe increases in each sample, the first shell (Fe-O) interactions are more pronounced. The fitting of the spectra was performed using single-scattering theory approach (dashed traces). The fits are in excellent agreement with the derived spectra, and the resulting structural parameters are given in Table 1. The Fe-O bond distances (R) obtained from EXAFS analysis were found to be 1.95 – 1.99 Å, shorter than the In−O distances in In2O3 (2.26 Å). This difference is due to the smaller ionic radius of Fe3+ (0.65 Å) relative to In3+ (0.80 Å). The average Fe-O distances slightly decrease with increasing doping concentration for a given NC structure, owing to the progressive shrinking of the host lattice with increasing substitutional incorporation of a smaller cation, in agreement with the XRD data in Figure 1. Similar results were obtained for Fe:SnO2 NCs (Figure 5c and d and Table 1). The Fe3+ coordination numbers obtained from the fitting were found to be between 2.6 and 3.9 for Fe:In2O3, and 3.9 – 4.1 for Fe:SnO2 NCs (Table 1). This reduction of the coordination number from the expected value (six) can be attributed to the internal structural disorder and a significant fraction of substitutional ions near the NC surface, which decrease the overall EXAFS amplitude of the Fe-O interaction.61 Taken together, the results of EXAFS analysis also suggests substitutional incorporation of Fe3+ in In2O3 and SnO2 NCs.


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Figure 6. (a) Absorption (top) and variable-field MCD (bottom) spectra of 5.6 % Fe:In2O3 NCs collected at 5 K. Inset: Magnified charge transfer transitions of Fe3+ dopants. (b) Magnetic fielddependent MCD intensity of 5.6 % Fe:In2O3 NCs recorded at different wavelengths, as indicated in the graph (the intensity at 400 nm is multiplied by a factor of 10 for clarity). Dashed lines represent saturation magnetization expected for Fe3+ (S=5/2) in In2O3 lattice. (c) The same as in (a) for 9.6 % Fe:SnO2 NCs. (d) The same as in (b) for 9.6 % Fe:SnO2 NCs. MCD spectroscopy measures the difference between the absorption of left and right circularly polarized light for an external magnetic field oriented parallel to the direction of propagation of the circularly polarized beam. Owing to the magnetic field dependence, MCD spectroscopy allows for molecular and electronic state-specific investigation of the electronic structure and magnetic interactions of transition metal ions in various matrices. Figure 6a and c shows absorption (top) and MCD (bottom) spectra of typical Fe:In2O3 and Fe:SnO2 NC samples, respectively, for different magnetic fields from 0 to 7 T, in 1 T increments. Both MCD spectra


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are dominated by the intense broad band peaking at ca. 305 nm for Fe:In2O3 and 320 nm for Fe:SnO2 NCs. This band is more symmetric in the case of Fe:SnO2, which also features a narrower peak on the high energy side. In addition, both spectra have a set of much weaker and more structured features in the range from 375 nm to 500 nm, although with the opposite signs (insets in Figure 6a and c). High-spin Fe3+ is a d5 system, generally exhibiting a very low intensity of the d-d transitions.62,63 The absorption bands at wavelengths shorter than 500 nm have been readily assigned to ligand-to-metal charge transfer (LMCT) transitions.62,64 Given that Fe3+ is the main form of dopant ions in both NC host lattices, we assign the series of bands in Figure 6a and c to LMCT transitions. Based on the general assignments of LMCT transitions in six-coordinate Fe3+ complexes, in the ideal Oh notation the lowest-lying bands between 420 and 450 nm may be tentatively assigned to symmetry-split t1gt2g, the bands/shoulders at 350-420 nm to t1u, t2ut2g , and the strongest bands at high energy (ca. 320 nm) to t1g, t1u, t2ueg transitions.62 The blue side of the main MCD band for Fe:In2O3 NCs, as well as the aforementioned narrow band at ca. 290 nm for Fe:SnO2 NCs coincide with the corresponding NC band gap transitions, suggesting dopant-induced splitting of the NC band states. Figure 6b and d plots the monochromatic MCD intensities as a function of external magnetic field for Fe:In2O3 and Fe:SnO2 NCs, respectively. The field-dependence of MCD intensity shows the same behavior at different wavelengths, suggesting the same origin for these transitions.64 This dependence is markedly different from that expected for S=5/2 spin only behavior characteristic for Fe3+ calculated using the reported g-value (dashed line),65 further confirming that the observed spectral bands are not due to intra-ionic ligand-field transitions,66 and/or there is a significant degree of spin-orbit coupling. Similar discrepancy is also found for magnetic


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susceptibility measurements of free-standing NCs. The results of MCD measurements confirm the Fe3+ nature of dopant ions, and indicate magnetization of the NC host lattices. The key question in the context of the electronic structure of iron dopants discussed above is how it relates to the overall magnetic properties, from which general conclusions about the mechanism of dilute magnetic ordering in transparent metal oxides could emerge. Figure 7a shows the room-temperature magnetic hysteresis loops for typical Fe-doped In2O3 (blue line) and SnO2 (green line) nanocrystalline films prepared from colloidal NCs as building blocks. These curves are compared to the magnetization data for a Mn:In2O3 and Mn:SnO2 nanocrystalline films previously investigated (dashed lines). Saturation magnetization per dopant ion for these four types of nanocrystalline films is plotted as a function of doping concentration in Figure 7b. For all samples, the maximum value of ms is measured at low doping levels, followed by a gradual decrease in ms with increasing doping concentration. While the magnetization of Fe:SnO2 is generally smaller than that for Fe:In2O3 nanocrystalline films for similar doping concentrations, both Fe:In2O3 and Fe:SnO2 samples have significantly lower ms values than the analogous Mn-doped films. It has been previously proposed that long-range magnetic ordering in different transition metal dopants in oxide nanocrystalline films and aggregates arises from the existence of defects formed at the NC interfaces, which play a key role in mediating magnetic ordering of dopant ions.30 Importantly, oxygen vacancies, as electron donors, have been proposed to be a key component of these grain boundaries.67 According to the magnetic polaron model,27 ferromagnetic exchange is mediated by shallow donor electrons that in conjunction with magnetic dopants form bound magnetic polarons. These polarons can then overlap to create a spin-split impurity band. For early and late 3d transition metal dopant ions, high TC occurs when


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empty majority- and minority-spin states, respectively, lie at the Fermi level in the impurity band.

Figure 7. (a) Magnetic hysteresis loops for Fe:In2O3 (blue trace) and Fe:SnO2 (green trace) nanocrystalline films prepared from 1.8 % Fe:In2O3 NCs having an average diameter of 8.6 nm and 2.5 % Fe:SnO2 NCs having an average diameter of 2.5 nm, respectively. The analogous data for Mn:In2O3 (dashed black trace) and Mn:SnO2 (dashed red trace) having similar doping concentrations are shown for comparison. (b) Saturation magnetization dependence on doping concentration for Fe:In2O3, Fe:SnO2, Mn:In2O3, and Mn:SnO2 nanocrystalline films. This model is not applicable for some high-TC DMOs containing mid-3d series transition metal dopants. In some of these systems it has been found that dopant ions exist in a mixed valence state. One of the notable examples particularly relevant for this work is Mn-doped In2O3 and SnO2 nanocrystalline films, in which manganese dopants coexist in 2+ and 3+ oxidation states.42,43 These systems exhibit robust ferromagnetism/superparamagnetism at room temperature despite showing no signature of magnetization associated directly with manganese.42 The observed magnetization in these and related systems has been described by CT ferromagnetism, in which charge transfer from the reduced form of dopant ions to the local density of states associated with extended structural defects leads to the Stoner-type defect band


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splitting.29 The role of dopant ions in this model is not to carry a magnetic moment, but to provide a reservoir of charge carriers. Similarly to the bound magnetic polaron model, high-TC CT ferromagnetism requires percolation of structural defects, which usually occurs at the NC interfaces. Mn2+ and Fe3+ are isoelectronic, and are expected to experience similar magnetic properties. According to the magnetic polaron model, d5 dopants should exhibit very low magnetic moment owing to the energy position of the minority and majority 3d bands.27 However, in contrast to Fe-doped In2O3 and SnO2 nanocrystalline films, identically prepared Mn-doped In2O3 and SnO2 samples show robust ferromagnetism or superparamagnetism at room temperature. The most striking difference between these two sets of DMO systems is the presence of mixed valence states in the case of Mn-doped NCs, consistent with the hypothesis of CT ferromagnetism. Another stipulation of the CT ferromagnetism is the lack of obvious coercivity in the M vs H measurements due to the absence of spin-orbit interaction as the primary source of magnetocrystalline anisotropy, which is also observed in our measurements for both Fe:In2O3 and Mn:In2O3 nanocrystalline films (Figure 7a). It is possible that other mechanisms of ferromagnetic ordering have a minor contribution to the observed behavior, including magnetic polaron and RKKY mechanisms. It has been suggested theoretically that at small and intermediate distances between Mn dopants in CdSe NCs the dopant centers are ferromagnetically ordered, while at larger distances their ordering is antiferromagnetic.68 This result is in agreement with RKKY mechanism, and is consistent with a slight increase in magnetic moment observed for Mn:In2O3 for doping concentrations approaching 20 %. The results of the present work contrast the findings by Feng et al.69 who investigated the correlation between the grain size and the magnetic properties of Fe and Mn-doped In2O3 films prepared by pulsed laser deposition. They found that Fe:In2O3 is ferromagnetic with the magnetic


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moment increasing with increasing grain size, which has been associated with the incorporation of Fe within the grains. On the other hand, for the Mn:In2O3 films, the magnetization was the largest for the samples having the smallest grain size, suggesting the localization of dopants in the grain boundaries. Although, the importance of the dopant electronic structure was not discussed in much detail, it should be emphasized that the authors found the evidence of the mixed valence state or iron (Fe2+ and Fe3+) in their samples. In our work both Fe and Mn are internally incorporated and the average sizes of host NCs are well-controlled, eliminating the dependence of the magnetic properties on the grain size and dopant location. The correlation between the dopant oxidation states and ferromagnetism, characteristic for the CT model is, therefore, fully consistent for internally Fe-doped In2O3 nanocrystalline films prepared by different methods. The ability to control the oxidation states of specific dopants by a sample preparation method provides yet another way of controlling the magnetic properties of DMOs. It is also instructive to discuss the difference between the magnetization of Fe-doped In2O3 and SnO2 samples in Figure 7. The Fe:SnO2 NCs are significantly smaller than Fe:In2O3 NCs, indicating larger surface/interface area in Fe:SnO2 nanocrystalline films, and a higher concentration of extended interfacial defects. Although these grain boundaries are important for CT ferromagnetism, Fe:SnO2 nanocrystalline films have extremely low magnetic moments, in many cases below the detection limit. Taken together, these results underscore that the presence of grain boundaries is not a sufficient condition for ferromagnetism in non-magnetic oxides. High TC ferromagnetism in DMOs also requires the presence of transition metal dopants or other defects which can electronically interact with the grain boundaries. Finally, it should be noted that our results confirm that magnetization in nanocrystalline films is not coming from the presence of secondary phases. Iron(III) oxide (hematite) is ferromagnetic, and with increasing


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doping concentration any tendency of phase segregation is expected to result in enhanced ferromagnetism. This is, however, exactly opposite from what we observed in this work, confirming that Fe3+ ions remain internally doped in nanocrystalline films. This conclusion is also in agreement with our previous results on Mn-doped In2O3 and SnO2.42,43

CONCLUSIONS In summary, we investigated the electronic structure and magnetic properties of Fe-doped In2O3 and SnO2 nanocrystalline films prepared from colloidal NCs as building blocks. We demonstrated that dopant ions exist in 3+ oxidation state, with a possibility of a minor presence of Fe2+ (up to 10 %). The measured magnetization is negligible relative to similarly prepared Mn-doped In2O3 and SnO2 samples in which Mn dopants exist in a mixed valence state (2+ and 3+). The increase in the concentration of extended interfacial defects associated with smaller SnO2 NCs was, alone, not enough to induce ferromagnetism in Fe:SnO2 films. Together with the other evidence in this work, these findings are consistent with the proposed CT ferromagnetism, as an emerging form of magnetism in defective solid-state materials. The results of this work also demonstrate the possibility of achieving, probing, and manipulating ferromagnetism in nonmagnetic oxides by simultaneous control of the electronic structure of transition metal dopant ions and extended defect states using chemical means.

ASSOCIATED CONTENT Supporting Information. EXAFS analysis, XRD pattern of Fe:rh-In2O3 NCs, NC size distribution histograms of Fe:In2O3 NCs, additional TEM images of Fe-doped In2O3 and SnO2


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NCs, EDX and EELS dopant distribution analysis, additional Fe K-edge XANES spectra of Fedoped In2O3 and SnO2 NCs, XANES pre-edge analysis of Fe-doped In2O3 and SnO2 NCs and reference compounds. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] Notes The authors declare no competing financial interest. ACKNOWLEDGEMENTS This work was supported by the Natural Sciences and Engineering Research Council of Canada (RGPIN-2015-04032), Canada Foundation for Innovation (CFI-LOF 204782), and the University of Waterloo (UW-Bordeaux Collaborative Research Grant). The research described in this article was partly performed at Canadian Light Source (CLS), which is supported by the NSERC, NRC, CIHR, the Province of Saskatchewan, Western Economic Diversification Canada, and the University of Saskatchewan. We thank Dr. Ning Chen (CLS) for his assistance at the HXMA beamline. M.H. acknowledges CLS for a Graduate Travel Awards. P.V.R. thanks Canada Research Chairs program for generous support.


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