Unique Magnetic Properties of Single Crystal γ-Fe2O3 Nanowires

LETTER pubs.acs.org/NanoLett

Unique Magnetic Properties of Single Crystal γ-Fe2O3 Nanowires Synthesized by Flame Vapor Deposition Pratap M. Rao and Xiaolin Zheng* Department of Mechanical Engineering, Stanford University, Stanford, California 94305, United States

bS Supporting Information ABSTRACT: Single crystal γ-Fe2O3 nanowires with 40
60 nm diameters were grown for the first time by singlestep atmospheric flame vapor deposition (FVD) with axial growth rates up to 5 μm/minute. Because of their superior crystallinity, these FVD γ-Fe2O3 nanowires are single magnetic domains with room temperature coercivities of 200 Oe and saturation magnetizations of 68 emu/g. KEYWORDS: Nanowire, flame synthesis, magnetic, gammaFe2O3, coercivity, curling

he ferrimagnetic material γ-Fe2O3 is of great technological importance. Nanoscale γ-Fe2O3 has found uses in diverse areas including high-density magnetic recording,1 ferrofluids,2 magnetic separation in biology,3 magnetic tweezers,4 drug delivery,5 magnetic resonance imaging,6 hyperthermia treatment of cancer,7 spintronics,8 magneto-optics,9 photocatalysis,10,11 chemical catalysis,12 magnetic recovery of catalysts,13 and chemical sensing.14,15 In particular, γ-Fe2O3 nanowires (NWs) are of great interest because their magnetic and optoelectronic properties can be tuned by controlling their diameters and aspect ratios.16,17 For example, high-aspect ratio γ-Fe2O3 NWs have larger magnetic coercivities than nanoparticles of the same volume because of a shape-induced energy barrier that hinders the thermal rotation of magnetic moments.16 Besides the shape of the NWs, their crystallinity also has a profound effect on their properties. For instance, the size of the energy barrier against moment rotation scales with crystallite volume.18,19 For a given shape and size of γ-Fe2O3 NWs, single crystal NWs have the largest possible crystallite volume, through which the energy barrier against thermal moment rotation is maximized. Therefore, it is highly desirable to synthesize single crystal γ-Fe2O3 NWs in order to take full advantage of their large aspect ratios so that their coercivities are enhanced. γ-Fe2O3 NWs and nanotubes have been synthesized by methods including electrodeposition in solution or sol
gel synthesis with the use of templates,20,21 solution synthesis without templates,22 and transformation of R-Fe2O3 NWs synthesized by solid phase thermal oxidation to γ-Fe2O3 NWs in the presence of hydrogen gas.23 In all of these methods for the synthesis of γ-Fe2O3 NWs, intermediate NWs of other iron oxide phases (R-Fe2O3, R-FeOOH, or Fe3O4) were grown first and then reduced, oxidized, or transformed to γ-Fe2O3. Such procedures are difficult to accomplish with good control over the NW crystallinity and phase purity. Furthermore, the solution methods inevitably introduce contaminants and impurities into

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the NWs, and often (always in the case of electrodeposition) lead to polycrystalline materials. Finally, there is only one report of wirelike chains of γ-Fe2O3 nanoparticles produced by chemical vapor deposition (CVD) under the presence of an orienting magnetic field.24 Thus, it is of interest to grow single crystal γ-Fe2O3 NWs and study their magnetic properties without these complicating features. Herein, we present a synthesis method that produces single crystal γ-Fe2O3 NWs with aspect ratios exceeding 100 by flame vapor deposition (FVD). To the best of our knowledge, the FVD synthesis presented here is the first demonstration of direct, single-step vapor deposition growth of γ-Fe2O3 NWs. Moreover, this FVD method is catalyst-free, atmospheric, and produces NWs at large axial growth rates. Significantly, the FVD γ-Fe2O3 NWs, due to their superior crystallinity, exhibit increased roomtemperature coercivities compared to those of similarly sized γ-Fe2O3 NWs synthesized by other methods. As shown in Figure 1, the FVD synthesis is conducted by using a 2 in.  2 in. square multielement coflow flat-flame diffusion burner (Hencken Burner) operating on CH4 and H2 fuels, with air as the oxidizer, as described in detail in our previous works on the flame synthesis of other metal oxide NWs and nanobelts (solid phase diffusion synthesis of R-Fe2O3 and CuO NWs, and vapor phase deposition synthesis of WO3 and W18O49 NWs and R-MoO3 nanobelts).25
27 The source of iron oxide vapor is a 2.5 in.  2.5 in. plain steel wire mesh (0.025 in. wire diam, 0.125 in. wire spacing, McMaster-Carr) placed 0.5 in. above the flame at a temperature of around 1550
C. The surface of this mesh is continuously oxidized by the flame product gases and evaporated to generate a steady stream of iron oxide vapors that flow downstream and deposit (in the form of NWs) onto the surface Received: March 7, 2011 Revised: April 27, 2011 Published: May 12, 2011 2390

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Figure 1. Schematic of FVD setup for Type 1 and 2 γ-Fe2O3 NW growth.

of a colder growth substrate, that is placed 3 in. above the flame, at the centerline axis of the burner. The growth substrates include iron foils (Alfa Aesar, 0.1 mm thick, 99.99% purity) and silicon (100) wafers (p-type, 1016/cm3, 0.5 mm thick). The flow rates of the fuels, CH4 and H2, are 2.4 SPLM and 9.6 SLPM, respectively, while the flow rate of the oxidizer, air (21% O2 þ 79% N2), is 28.6 SLPM, yielding an equivalence ratio of Φ ≈ 1.6 (fuel rich). Under these conditions, the postflame gas is rich in H2, H2O, CO2, and CO. The postflame region is enclosed in an openended quartz tube (3 in. diameter) to reduce air entrainment and ensure that the temperature in planes perpendicular to the flow direction is uniform. Temperatures are measured using an S-type thermocouple (Pt/Pt-10%Rh, 125 μm bead size) coated with SiO2, and the temperatures of the source mesh and growth substrates are measured in their wake. These temperatures are found to be very stable, fluctuating by only a few degrees Celsius over the duration of the growth. This FVD method offers several important advantages over conventional furnace-based CVD. First, the growth occurs at atmospheric pressure without the need for vacuum pumps or chambers. Second, the solid source materials are easily heated to temperatures significantly higher than those achievable with a CVD furnace to generate large quantities of oxide vapor and produce large nanostructure growth rates. In addition, the heat is chemically generated, allowing for simple scale-up of the growth process by increasing the burner area. Finally, the postflame gas phase composition naturally includes O2, CO2, CO, H2O, and H2, and the relative concentrations of these species (particularly the partial pressure of O2) can easily be controlled over tens of orders of magnitude by varying the fuel and oxidizer equivalence ratio fed into the flame.25 Thus, FVD offers facile control over the composition of metal oxides. The morphology, crystal structure, and composition of the FVD γ-Fe2O3 NWs were characterized by scanning electron microscopy (SEM, FEI Sirion XL30, 5 kV), transmission electron microscopy (TEM, FEI Tecnai G2 F20 X-TWIN FEG, 200 kV), X-ray diffraction (XRD, SSRL wiggler beamline 11
3, 12.7 keV, 0.15  0.15 mm spot, in transmission with MAR345 imaging plate detector at 150 mm), and X-ray photoelectron spectroscopy (XPS, SSI S-Probe, Al-kR, 1486 eV, 150  800 μm spot). SEM and XPS analyses were done on the as-grown samples. For the TEM analysis, the NWs were scraped onto carbon-coated copper grids using a razor blade, while for the XRD measurements, the NWs were removed from the growth substrate by adhesive tapes and kept on these tapes for the measurements. The background signal from the tape was removed from the XRD signal for refinement and presentation, using a separate scan of

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Figure 2. γ-Fe2O3 NWs grown by FVD for 30 min. Type 1 NWs on (a) Fe substrate, side view; (b) Fe substrate, top view; (c) Si substrate, top view. Type 2 NWs on the edges of Si substrates (d
f, side views).

the bare tape in an identical configuration (see Supporting Information). The magnetic properties of the NWs were characterized by a superconducting quantum interference device magnetometer (SQUID, Quantum Design MPMS-XL). For the magnetic measurements, the NWs were removed from the iron growth substrate (so that the magnetic properties of the NWs themselves could be measured without influence from the magnetic iron substrate) by nonmagnetic adhesive tapes and kept on these tapes for the measurements. This method of NW removal was determined to lead to much less damage to the NWs than other possible removal methods such as sonication or scraping. Subsequently, the mass of Fe in the NWs on the tapes, and thereby the mass of the γ-Fe2O3 NWs, was accurately determined by inductively coupled plasma mass spectrometry (ICP-MS, Thermo Scientific XSERIES 2) after dissolution of the NWs in 2% hydrochloric acid and deionized water to make an 8 mL solution. Thereafter, the measured mass was used to normalize the measured magnetization. Using the FVD method, we have grown both quasi-aligned NW arrays (Type 1, Figure 2, top panel) and large tree-like networks of NWs (Type 2, Figure 2, bottom panel), both of yellow/orange color when viewed with the naked eye. The Type 1 NWs reach lengths exceeding 10 μm after 30 min of growth and have uniform and large surface coverage densities (108/cm2) on the flat surfaces of the iron foils (Figure 2a and b), and much smaller surface coverage densities with less uniformity on the flat surfaces of the Si wafers (Figure 2c). For Type 1 NW growth, the substrates, with typical dimensions of 2 cm 1 cm, are assembled into sandwiched, dead-end slots with a cover slide and a gap between the substrate and cover of approximately 0.1 mm, fixed using a steel shim spacer (Figure 1). The purpose of the dead-end slot geometry is to increase the residence time of the iron oxide vapors on the substrates. Up to eight such substrates can be placed vertically (back-to-back) in the postflame region for concurrent growth of the NWs, yielding a total of order 50 μg of NWs for 30 min of growth at a substrate temperature of 650
C (although the temperature inside the dead-end slot may be considerably lower). The Type 2 treelike NW networks (Figure 2d
f) grow at the edges and corners of the iron and Si substrates with 150 μm overall length after 30 min of growth. This type of NW growth reveals the rapid nature of this flame synthesis method with axial growth rates of individual NW branches reaching 5 μm/min, which is 15 times higher than the growth rate of the Type 1 NWs. For Type 2 NWs, no special substrate geometry is required except for the presence of edges and corners, such as those present around the perimeter of a cut foil or wafer. For this type 2391

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Figure 3. TEM analysis of the FVD γ-Fe2O3 NWs. (a) TEM image of a NW with a broken tip and a portion of the root intact. (b
e) HRTEM images of the (b) Æ111æ, (c) Æ110æ, (d) Æ332æ, and (e) Æ112æ zones of the NW crystal, each showing the {220} growth planes that lie perpendicular to the NW axis, with accompanying SAD patterns in the insets. No evidence of faults, defects, or internal boundaries was found in any studied NW.

of growth, single substrates of dimensions 0.5 cm  0.5 cm are placed parallel to the flame in the postflame region (Figure 1) for 30 min at a substrate temperature of approximately 750
C, yielding of order 10 μg of NWs at the edges of the substrate. The quantity of both Type 1 and 2 NWs can be scaled up by increasing the number of concurrent growth substrates or size of the burner. The diameters of both types of NWs lie in the narrow range of 40
60 nm (from TEM images such as Figure 3a), resulting in aspect ratios exceeding 100. The NWs are slightly tapered, such that their roots have slightly larger diameters than their tips. The crystal structure of both types of NWs (XRD, Figure 4a) show a pure phase that matches very closely to that of cubic γ-Fe2O3 (space group P4332, ICDD PDFs 01-083-0112 and 00-039-1346, a = 8.347
8.351 Å). All expected peaks with I/I311 > 3% are observed, and the cubic unit cell dimension calculated from refinement of the measured pattern is 8.342 Å, which is very close to the range of cell dimensions expected for cubic γ-Fe2O3, but significantly smaller than the cell dimension expected for Fe3O4 (space group Fd3m, ICDD PDFs 04
007
2718 and 04
005
4319, a = 8.375
8.396 Å). Thus, the NWs likely consist of γ-Fe2O3, as opposed to Fe3O4. Presently, the XRD data is not of sufficient quality to determine whether the cation vacancies in the γ-Fe2O3 are partially or fully ordered over the octahedral sites.28 The chemical composition of Fe2O3 as opposed to Fe3O4 is further confirmed by XPS analysis of the Fe 2p orbitals (Figure 4b). In particular, the Fe 2p3/2 binding energy peak occurs at 710.6 eV, which is close to the characteristic peak for the 3þ ion in γ-Fe2O3 (710.829 or 710.7 eV30), and there is no signal or shoulder at smaller binding energies as would be expected for presence of the 2þ ion (708.329 or 708.5 eV30). Moreover, there exist well-resolved satellite peak structures at the higher binding energy sides of the main doublet peaks, which also indicate the absence of the 2þ ion.29,30 Finally, the XPS signal

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Figure 4. (a) Background-subtracted XRD pattern and Pawley fit, and (b) Fe 2p XPS spectrum of the FVD γ-Fe2O3 NWs.

comes from γ-Fe2O3 and not from some accidentally occurring surface R-Fe2O3 layer because the Fe 2p3/2 peak in Figure 4b is a broad maximum instead of being composed of two distinct peaks, as is characteristic of R-Fe2O3.29 Presently, it is not possible to totally exclude the possibility of carbon incorporation into the NWs during growth in the carbon-containing flame, but XPS analysis suggests that it is unlikely (see Supporting Information). Moreover, the carbon-containing fuel, methane, does not contain C
C bonds that promote soot formation. The crystal size in the NWs, determined by refinement of XRD patterns, is approximately 40 nm. This strongly suggests that the NWs are single crystals, since their physical diameters are in the range 40
60 nm, which matches the crystal size. Further evidence for the single crystal nature of the NWs is found from high-resolution TEM (HRTEM) images and selected-area diffraction (SAD) patterns (Figure 3b
e) taken from different orientations that show the NWs have cubic crystallography with a Æ110æ growth direction and seem to be perfect single crystals with no evidence of faults, defects, or internal boundaries. Therefore, on the basis of color, TEM, XRD, and XPS analyses, both Type 1 and 2 FVD NWs are shown to be single crystals of cubic γ-Fe2O3 with axes along the Æ110æ crystal direction and diameters of 40
60 nm with lengths of several micrometers. The growth mechanism of both types of NWs is likely to be direct vapor
solid growth, as opposed to vapor
liquid
solid or solid diffusion growth, based on the fact that both types of growth occur on substrates other than iron (i.e., silicon) and without the presence of catalysts. Our experiments have shown that Type 1 and 2 NWs have nearly identical magnetic properties, so we will discuss these properties without distinction. Figure 5 shows the representative magnetization versus field hysteresis loops at 300 and 30 K for a 2392

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Figure 5. Magnetization vs field hysteresis loops for a collection of randomly oriented FVD γ-Fe2O3 NWs at 300 and 30 K.

collection of FVD γ-Fe2O3 NWs with spherically random orientations. At 300 and 30 K, the collection of NWs has saturation magnetizations (Ms) of 68 and 73 emu/g, remanence to saturation magnetization (Mr/Ms) ratios of 0.34 and 0.37, and coercivities (Hc) of 220 and 380 Oe, respectively. The magnetization of the NWs does not saturate even at applied fields of 10 kOe, as is seen by the positive slope of the magnetization even at the end points of the hysteresis curves. The blocking temperature of the NWs, measured using field-cooled and zero-fieldcooled magnetization procedures, was well above room temperature (see Supporting Information). Although the packing fraction of NWs on the tapes was kept low in an effort to minimize interparticle interactions, the desirability of good signal-to-noise ratio in these magnetic measurements prevented further reductions in packing. Therefore, there are some interparticle interaction effects, but these will be quantified in the following discussion, such that the properties of individual, isolated FVD γ-Fe2O3 NWs can be elucidated. The room temperature Ms of the NWs from Figure 5 (68 emu/g) is only slightly smaller than the room temperature value of 74 emu/g for bulk γ-Fe2O3, which is further evidence for the purity of the material.19 The measured Ms is smaller than the bulk value and the slope of the magnetization is positive even at large applied fields because NWs, which are finite, nonellipsoidal particles, cannot be saturated by a homogeneous field, such as that applied in this measurement.31 The ensemble measurement in Figure 5 also sheds light on the remanent state of the FVD γ-Fe2O3 NWs. Mr is maximized for uniformly magnetized single domain structures, which have Mr/ Ms = 1 when magnetized axially. The individual FVD γ-Fe2O3 NWs should be nearly perfect single-domain structures at the remanent state because their diameters are much smaller than the critical size below which domain wall formation becomes energetically unfavorable (166 nm for spherical particles of γ-Fe2O3, and much larger for particles with considerable shape anisotropy).18 In fact, based on the reduced diameter S = 2.67 and the large aspect ratio (S = d/d0, where d = 40 nm is the average NW diameter determined by XRD, and d0 = 15 nm is twice the exchange length in γ-Fe2O3),32 the isolated NWs are expected to have an “in plane flower” remanent state, which is a highly uniform single domain state of magnetization in which almost all of the magnetic moments point along the NW axis and Mr/Ms ≈ 1.33,34 On the other hand, because of geometry, a collection of noninteracting single domain particles with spherically random distribution should have a low temperature Mr/Ms

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ratio of 0.5 (or slightly larger, due to intrinsic cubic magnetocrystalline anisotropy35).36 Moreover, when interparticle interactions exist, the Mr/Ms ratio will deviate from the value of 0.5. Making the assumption that the NWs are single domain structures, we can estimate the Mr/Ms ratio of the NW ensemble using the results of micromagnetic simulations and effective-medium theory.37 To this end, we calculate a dimensionless parameter λ = μ0Ms2/6K (where the per volume magnetic anisotropy energy K = 1/2 μ0HcMs, and μ0 = 4π  10
7 J/mA2 is the permeability of free space), which compares the strength of the interparticle magnetostatic interactions to the particle anisotropy energy. For these γ-Fe2O3 NWs at 30 K with λ = 4.25 (log10 λ = 0.63), the simulations and theory predict a value of Mr/Ms ≈ 0.35, which is very close to the measured value of Mr/Ms = 0.37.37,38 Thus, our assumption is correct, and the NWs are single domain structures. This reduction of Mr/Ms from the noninteracting value of Mr/Ms = 0.5 occurs because the NWs form multiparticle flux closure domains to reduce the total magnetostatic energy of the ensemble.37 Finally, the room temperature values of Ms and Mr/Ms (Figure 5) are smaller than the corresponding low-temperature values simply due to thermal effects. In addition, Figure 5 shows the coercivity of the ensemble of FVD γ-Fe2O3 NWs. The 380 Oe coercivity of these FVD γ-Fe2O3 NWs at 30 K is a consequence of noncoherent magnetization reversal. The coherent magnetization reversal coercivity of single, isolated γ-Fe2O3 NWs should be 2320 Oe (2260 Oe due to demagnetization shape anisotropy energy caused by the large aspect ratio, and 60 Oe due to the uniaxial component of intrinsic magnetocrystalline energy).32,39,40 If the magnetization is reversed by the coherent rotation of moments, the lowtemperature coercivity of this collection of γ-Fe2O3 NWs with spherically random orientations would be 1110 Oe.36 However, since the NW diameters (40
60 nm) are considerably larger than d0 = 15 nm, which is twice the exchange length in γFe2O3,32 their magnetization reversal will not occur by coherent rotation of the magnetic moments.31 Rather, the reversal in these NWs is expected to occur by noncoherent reversal modes, which are lower energy paths of reversal than coherent rotation, leading to reduced coercivity. Most likely, the reversal occurs by curling, which is the noncoherent rotation of moments away from the NW axis direction in the plane perpendicular to the radius at each point, in an amount that is a function of radius only41 because curling has the lowest energy barrier among all reversal modes for NWs in this size range.32 To verify whether the reversal does in fact occur by curling, we compared the coercivity predicted by the micromagnetic curling model to our measured coercivity of 380 Oe at 30 K. For these γ-Fe2O3 NWs with reduced diameter S = 2.67, the curling model predicts a coercivity that is 0.169 times that of coherent rotation (2320 Oe), or approximately 390 Oe, for a collection of noninteracting NWs with spherically random orientation at low temperature.42 This coercivity is only slightly larger than the measured value of 380 Oe at 30 K and strongly suggests that the reversal of these NWs occurs by curling. The effect of interparticle interactions accounts for most of the difference between the theoretical and measured values; for an estimated NW packing fraction of 10%, the measured coercivity should be around 15 Oe smaller than the noninteracting value.43 Thus, these FVD γ-Fe2O3 NWs conform remarkably well to the micromagnetic curling model, despite the fact that they have finite (albeit large) lengths, while the curling model assumes infinitely long cylinders.41 Finally, the room-temperature coercivity of 220 2393

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Nano Letters Oe (Figure 5) is smaller than the low-temperature coercivity simply due to thermal effects. Using the curling model and the ensemble measurements, we may estimate with confidence that the axial coercivity of single, isolated FVD γ-Fe2O3 NWs will be 350 Oe at 30 K and 200 Oe at 300 K.40,42 Furthermore, the per volume magnetic anisotropy energy (K = 1/2 μ0HcMs) calculated from the estimated low-temperature single-NW coercivity (Hc = 350 Oe = 2.8  104 A/m) and saturation magnetization (Ms = 73 emu/g = 3.6  105 A/m), is approximately 6.2  103 J/m3, which for the NW volume (V > 1  10
21 m3) results in a ratio of magnetic anisotropy energy to thermal energy KV/kbT > 1500 at room temperature. Since the magnetization reversal mechanism is highly dependent on the structure of the material, the fact that our FVD γFe2O3 NWs reverse their magnetization by curling confirms that their shapes are highly regular without internal defects such as pores. The curling mode cannot occur in the presence of certain irregularities of shape or internal defects. For example, different reversal mechanisms were observed in the much-studied elongated γ-Fe2O3 particles that were used for decades as the recording medium in magnetic tapes.1 In one detailed study, the coercivities of individual (mostly single-crystal) γ-Fe2O3 particles (diameter, 50
90 nm; aspect ratio, ∼10) were found to range widely from 400 to 1100 Oe at room temperature.32 The particles with coercivities at the lower end of the range had highly irregular “peanut” shapes and were reported to reverse their magnetization by fanning,44 while particles with coercivities at the higher end of the range had regular spindle shapes but contained internal pores and were reported to reverse their magnetization by buckling.41 Furthermore, the coercivities of the FVD NWs can be compared to those of γ-Fe2O3 NWs obtained by the transformation of R-Fe2O3 NWs23 with diameters of 50
90 nm and similarly large aspect ratios. The ensemble coercivity of the transformed NWs was 360 Oe at 30 K, which is very similar to that of the FVD NWs (380 Oe), indicating that both NWs have similar per volume energy barriers for resisting magnetic moment rotation. However, the ensemble coercivity of the transformed NWs was zero at 300 K (superparamagnetic behavior) in contrast to that of the FVD NWs (220 Oe), indicating that the total per NW energy barrier against moment rotation is much smaller in the case of the transformed NWs. The smaller energy barrier implies smaller crystallite volume, which suggests that the transformation from R-Fe2O3 to γ-Fe2O3 may be incomplete. Superparamagnetic behavior at room temperature was also observed for γ-Fe2O3 NWs synthesized by the following various solution-based methods: (1) γ-Fe2O3 NW21 and nanotube20 arrays (both with diameter of 200 nm) synthesized by templateassisted electrodeposition and sol
gel techniques, respectively, and (2) γ-Fe2O3 NWs (diameter, 20 nm) grown in solution without the use of templates.22 Studies have shown that singlecrystal γ-Fe2O3 nanoparticles of 20 nm diameter have roomtemperature coercivities greater than 50 Oe,45 and even polycrystalline γ-Fe2O3 particles with crystallite sizes as small as 5 nm have some room-temperature coercivity.19 Thus, the solutiongrown γ-Fe2O3 NWs and nanotubes with zero coercivity at room temperature likely have crystallite sizes less than 5 nm and are highly polycrystalline. Finally, the ensemble room temperature coercivity of the FVD NWs is about twice that of wirelike chains of 50 nm γ-Fe2O3 nanoparticles synthesized by CVD24 (120 Oe). The lower coercivity of the wire-like chain of nanoparticles can be attributed to the weak magnetic exchange coupling

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between the axially arranged crystals of the nanoparticle chains, while strong exchange coupling exists in the FVD NWs. These comparisons show that the room temperature coercivity of the FVD γ-Fe2O3 NWs is much larger than those of γ-Fe2O3 NWs synthesized by any other reported method. These distinct magnetic properties of our γ-Fe2O3 NWs, grown directly from vapor in a flame, are a consequence of their superior crystallinity. In summary, a single-step, rapid, atmospheric, catalyst-free flame synthesis method was used to synthesize pure, single crystal γ-Fe2O3 NWs with Æ110æ growth directions and 40
60 nm diameters on iron and silicon substrates with axial growth rates up to 5 μm/min. To the best of our knowledge, this FVD synthesis is the first demonstration of the direct growth of γ-Fe2O3 NWs from vapor and produces NWs of superior crystallinity. As a result of the improved crystallinity, these FVD γ-Fe2O3 NWs exhibit enhanced magnetic properties compared to γ-Fe2O3 NWs synthesized by other methods. Specifically, these NWs are single domain structures with room-temperature coercivities of 200 Oe, compared to the zero or small room-temperature coercivities previously reported for γ-Fe2O3 NWs, nanotubes, or NW-like structures synthesized by other methods. These NWs likely reverse their magnetization by the incoherent process of curling.

’ ASSOCIATED CONTENT

bS

Supporting Information. (1) Details of X-ray diffraction data collection and pattern refinement, (2) XPS quantification of carbon, and (3) FC and ZFC magnetization versus temperature measurements. This material is available free of charge via the Internet at http://pubs.acs.org.

’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected].

’ ACKNOWLEDGMENT P.M.R. sincerely thanks the Link Foundation for its support through the Link Foundation Energy Fellowship. X.L.Z. sincerely thanks PECASE for support of this work. Portions of this research were carried out at the Stanford Synchrotron Radiation Lightsource, a Directorate of SLAC National Accelerator Laboratory and an Office of Science User Facility operated for the U.S. Department of Energy Office of Science by Stanford University. The authors would like to thank the following people for their generous help with measurements: Maxwell Shapiro and Dr. Ian Fisher (SQUID), Mingliang Zhang and Dr. Shan Wang (AGM), and Matt Bibee and Dr. Apurva Mehta (Synchrotron XRD, refinement). ’ REFERENCES (1) Dronskowski, R. Adv. Funct. Mater. 2001, 11 (1), 27–29. (2) Holm, C.; Weis, J. J. Curr. Opin. Colloid Interface Sci. 2005, 10 (3
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