Fabrication, Formation Mechanism, and Magnetic Properties of Metal

Dec 16, 2010 - Department of Physics and Astronomy, UniVersity of Delaware, Newark, Delaware 19716, United States;. Department of Materials Science an...
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J. Phys. Chem. C 2011, 115, 373–378

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Fabrication, Formation Mechanism, and Magnetic Properties of Metal Oxide Nanotubes via Electrospinning and Thermal Treatment Xing Chen,† Karl M. Unruh,† Chaoying Ni,‡ Bakhtyar Ali,‡ Zaicheng Sun,§ Qi Lu,† Joseph Deitzel,| and John Q. Xiao*,† Department of Physics and Astronomy, UniVersity of Delaware, Newark, Delaware 19716, United States; Department of Materials Science and Engineering, UniVersity of Delaware, Newark, Delaware 19716, United States; NSF Center for Micro-Engineered Materials, Department of Chemical and Nuclear Engineering, UniVersity of New Mexico, Albuquerque, New Mexico 87131, United States; and Center for Composite Materials, UniVersity of Delaware, Newark, Delaware 19716, United States ReceiVed: August 30, 2010; ReVised Manuscript ReceiVed: NoVember 29, 2010

A simple procedure has been developed for preparing high aspect ratio nanotubes of R-Fe2O3 and Co3O4 with diameters less than 100 nm and wall thicknesses less than 25 nm based on an appropriate heat treatment of electrospun polymeric fibers containing Fe(III) and Co(II) ions. The transformation of the as-prepared nanofibers to the final nanotube structure has been studied by scanning and transmission electron microscopy as well as X-ray diffraction, differential scanning calorimetry/thermogravimetric, and X-ray photoelectron spectroscopy measurements. These measurements and comprehensive analysis have led to a semiquantitative picture of a new nanotube formation mechanism. On the basis of the principles established in this article, it is foreseeable that many other oxide nanotubes could be designed and fabricated, opening a broad avenue to investigate electrical, chemical, mechanical, and magnetic properties. In this particular case, we have shown that magnetic properties are very different between R-Fe2O3 nanofibers and nanotubes, and they are distinctly different from their bulk counterpart. 1. Introduction The ability to manipulate and control the structure of matter at the nanoscale has made many new classes of materials available for the study of fundamental physical processes and potential applications. The discovery of C nanotubes in 1991,1 for example, was soon followed by experimental and theoretical studies that established the unique and technologically useful properties of these materials.2 Subsequent studies demonstrated that nanotubes with interesting and useful properties could also be prepared from materials other than C.3 Many of these materials and fabrication methods have been summarized in a recent review article.4 More recently, an electrospinning apparatus based on a two-capillary coaxial spinneret has been used to prepare TiO2 nanotubes with relatively large diameters, and a single capillary approach has been used for the preparation of ZnO and CeO2 nanotubes.5-7 In this article we describe a simple process for preparing high aspect ratio R-Fe2O3 and Co3O4 nanotubes with diameters less than 100 nm from heat-treated Fe(III) and Co(II) containing polymeric nanofibers (a US patent based on the process described in this paper was filed in April 2009). A standard electrospinning apparatus was used to prepare the precursor nanofibers because of its simplicity, versatility, low cost, controllability, and scale-up potential8-13 and because electrospun precursors had been successfully used to prepare metal oxide nanofibers.14-16 As work progressed, however, it soon became clear that control and optimization of the nanotube * Corresponding author. E-mail: [email protected]. † Department of Physics and Astronomy, University of Delaware. ‡ Department of Materials Science and Engineering, University of Delaware. § University of New Mexico. | Center for Composite Materials, University of Delaware.

formation process would require a systematic study of the thermal evolution in the structural and chemical properties of the as-spun nanofibers. These studies have allowed us to identify a set of necessary conditions that must be satisfied in order to reliably and reproducibly produce R-Fe2O3 and Co3O4 nanotubes and which may also prove useful in guiding the choice of precursors and heat treatments for the preparation of other nanotube materials as well. 2. Experimental Section The electrospinning apparatus used in this work consisted of a single stainless steel capillary tube maintained at a potential of 12 kV with respect to a ground plate positioned about 10 cm from the capillary tip. A syringe pump allowed the material in the capillary to be replenished at a controlled rate of 0.4 mL/h as it was extracted by electrical forces from the Taylor cone at the open end of the capillary. The entire setup was placed in a glovebox with a controlled relative humidity of about 15% in order to promote the rapid evaporation of solvent from the fluid jet and to ensure the formation of a solid fiber prior to impact on the ground plate. The electrospun fibers were prepared from solutions consisting of 0.5 g of Fe(NO3)3 · 9H2O (98+% metals basis purity) or Co(NO3)2 · 6H2O (98+% metals basis purity) and 0.5 g of 1.3 × 106 molecular weight poly(vinylpyrrolidone) (PVP) dissolved in 5 mL of N,N-dimethylformamide (DMF) and 5 mL of 2-propanol. All chemicals were purchased from Alfa Aesar and were used as-received without further purification. The structure and chemical composition of the as-prepared and heat-treated fibers were studied by scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray diffraction (XRD), and X-ray photoelectron spectroscopy (XPS) measurements. Fiber diameter statistics were determined

10.1021/jp1082533  2011 American Chemical Society Published on Web 12/16/2010

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Figure 1. (a) SEM image of as-spun Fe(III)/PVP nanofibers, (b) SEM images of R-Fe2O3 nanotubes, (c) TEM images of R-Fe2O3 nanotubes, and (d) TEM images of Co3O4 nanotubes. The nanotubes were obtained from their respective precursor nanofibers after a heat treatment in air at 500 °C for 1 h.

by averaging the measured diameters of about 100 randomly chosen individual fibers in each of 2-3 separate SEM images. The XRD data were analyzed by fitting the background subtracted diffraction peaks to line shapes consisting of a sum of symmetric Pearson VII functions constrained to model the Cu KR1/KR2 components of the peak profile with wavelengths of 0.154 06 and 0.154 44 nm, respectively. Particle size estimates were obtained from the fitted line widths using the Scherrer formula after a quadrature correction for the instrumental line broadening. Best fit lattice parameters were obtained from the fitted peak positions using the CelRef software package.17 The thermal evolution of the samples was studied using a power compensated DSC and a heat flux DSC/thermogravimetric analyzer (TGA). The temperature and heat flow scales of the DSC and DSC/TGA were calibrated using the melting points of high purity In (Tm ) 156.60 °C) and Zn (Tm ) 419.53 °C) and to the latent heat of fusion of Zn (H ) 108.06 J/g).18 In each case the heating rate was 20 °C/min, and the purge gas was air. The XPS measurements were carried out using monochromatic Al KR radiation with the energy scale corrected for possible sample charging based on a C 1s binding energy of 284.6 eV. The resulting spectra were analyzed using the CasaXPS software package (Casa Software Ltd., U.K.). The magnetic properties of the Fe oxide nanotubes were studied using a vibrating sample magnetometer (VSM). All of the magnetic measurements were carried out at room temperature in applied magnetic fields up to 1 T. 3. Results and Discussion Figure 1a shows a typical low-magnification SEM image of as-spun Fe(III) containing PVP nanofibers (similar images of

the as-spun Co(II) containing PVP nanofibers were essentially indistinguishable). Most of the fiber diameters shown in this image fall between about 100 and 150 nm. Figure 1b-d shows higher magnification SEM and TEM images of the R-Fe2O3 and Co3O4 nanotubes obtained after thermal treatment in ambient air at 500 °C for 1 h. The TEM images clearly indicate that tubular structures with inner diameters between about 30 and 50 nm and wall thicknesses less than about 25 nm were formed in both cases. XRD measurements (see Figure S1 in Supporting Information) carried out on the heat-treated fibers revealed the formation of the single phase oxides R-Fe2O3 and Co3O4 with lattice parameters of a ) 0.5036(3) nm and c ) 1.375(3) nm for the Fe oxide and a ) 0.8088(3) nm for the Co oxide, in reasonable agreement with the reported bulk values of a ) 0.5034 nm and c ) 1.3747 nm and a ) 0.8065 nm, respectively.19 These measurements also yielded a Scherrer particle size of about 30 nm for the Fe oxide and about 40 nm for the Co oxide nanotubes roughly consistent with the TEM observations. In order to clarify the nanotube formation process, power compensated DSC and heat flux DSC/TGA measurements were carried out on the as-spun Fe(III) containing nanofibers. Both calorimetric measurement techniques exhibit two well-resolved exothermic signals, and the power compensated DSC result is shown in Figure 2a. The first peak centered at about 300 °C corresponds to an energy release of about 10.5 kJ/g and the second, centered at about 450 °C, corresponds to an energy release of about 1.1 kJ/g (these signals will be referred to as the first and second exotherms in the following). In addition to the survey scans, an as-spun nanofiber sample was scanned to a temperature of 385 °C, which is approximately intermediate

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Figure 2. (a) Heat flow (black), normalized fiber mass (red), and calculated fiber volume (blue) as a function of the temperature DSC heat-treated Fe(III) samples. The horizontal dashed blue line has been inserted to emphasize that the fiber volume remained essentially unchanged about 400 °C. (b) Heat flow (black) and normalized mass (red) as a function of the temperature for sample heat-treated in a non-heat-compensated TGA. Overheating was observed as indicated by the forward tilting in temperature scale of the heat flow curve.

Figure 3. SEM images of Fe(III) samples heat treated at various temperatures showing a decrease in fiber diameter up to about 400 °C. Fiber diameters remain relatively unchanged at higher temperatures.

between the two exotherms, rapidly cooled to room temperature, and then rescanned to a final temperature of 500 °C (see Figure S2 in Supporting Information). While no evidence of the first exothermic signal was observed in the second scan, the 450 °C exotherm was still clearly present. This observation indicates that from a kinetic standpoint the thermal processes corresponding to the two exothermic signals can be easily separated. The normalized evolution in the sample mass from the TGA measurements has been included in Figure 2a for comparison with the corresponding calorimetric signal. On the basis of the DSC survey data, a second set of Fe(III) samples was prepared and scanned to a series of temperatures associated with the various thermal events seen in Figure 2a. SEM images of these samples (see Figure 3) were used to estimate the normalized fiber volumes (assuming the nanofibers were solid cylinders and their length remained unchanged on thermal treatment) shown in Figure 2a. A comparison of the measured mass loss with the calculated volume reduction indicates that both quantities remain closely associated up to the completion of the first exotherm at about 400 °C. Above 400 °C, the fiber volume remains relatively constant while the mass loss continues. A change in the microstructure of the heat-

treated nanofibers also occurs at a temperature of about 400 °C, as shown in the TEM images in Figure 4. Below 400 °C the nanofiber images show no core/shell contrast, suggesting a uniform and homogeneous composition across the fiber diameter. At about 400 °C, however, core/shell contrast characteristic of the formation of an oxide shell begins to develop and becomes increasingly prominent at higher temperatures. The importance of the observation that the appearance of an oxide shell signals a minimum in the fiber diameter will be discussed at greater length below. In order to obtain a more detailed picture of the chemical evolution from precursor nanofibers to metal oxide nanotubes, X-ray photoelectron spectroscopy (XPS) spectra in the region of the C 1s, N 1s, and O 1s electronic states were obtained from an as-spun Fe(III) sample as well as Fe(III) samples heat treated at 255, 325, 385, and 475 °C. The results of these measurements are shown in Figure 5. In the as-spun fibers, the presence of N-O/O-N signals at 406.5/532.3 eV arise from the nitrate in Fe(NO3)3, and the small O-Fe signal at 529.3 eV may be attributed to the association of Fe(III) ions with the carbonyl oxygen in PVP as previously suggested.20 The C-C, C-N/N-C, and CdO/OdC XPS signals at 284.6, 285.8/399.5,

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Figure 4. TEM images of Fe(III) samples heat treated in DSC to the indicated temperatures. Core/shell contrast indicates the start of tube formation.

Figure 5. XPS spectra of as-spun and heat-treated Fe(III) samples near the C 1s (left), N 1s (middle), and O 1s (right) binding energies. The temperature indicated on the left column applies to all three frames in the same row. Note the presence of an unidentified peak at about 535.6 eV in the 325 °C heat-treated O 1s spectrum.

and 288.0/531.3 eV, respectively, are expected for PVP.21 The association between the Fe(III) ions and the carbonyl oxygen is probably also responsible for the observed shift in the CdO/OdC binding energy to a value about 0.4 eV greater than that of pure PVP as has been similarly observed in previous IR measurements.20 Between room temperature and 325 °C the disappearance of the N-O/O-N signal indicates the decomposition of nitrate ions which is responsible for the initial mass loss. The intensity increase of the O-Fe signal after 325 °C

indicates the formation of Fe2O3. Between 255 and 385 °C the increase in the CdO/C-N intensity ratio followed by a decrease is most likely the result of intermediaries formed during the decomposition of PVP. This process has previously been reported to begin with the breaking of the C-N bond connecting the pyrrolidone ring to the backbone of the polymer followed (in an oxidizing environment) by the oxidation of the resulting polyene and the eventual formation of CO2 and H2O.22,23 The large heat release in the DSC measurements is consistent with the mechanism of PVP oxidation. At 475 °C, the absence of N-C/C-N signals signifies the complete disassociation of the pyrrolidone ring from the polymer backbone while the small CdO/OdC signal indicates the continued presence of residual decomposition products. In this picture the mass loss observed at temperatures above about 255 °C is due to the volatilization of the PVP decomposition products, also consistent with the decrease and eventual disappearance of the C-O signal. It is worth noting that we were not able to identify a reasonable candidate for the XPS signal at 535.6 eV (gray line) in the sample heat treated at the intermediate temperature of 325 °C. The results of the SEM, TEM, DSC, TGA, and XPS measurements described above clarify the transformation process from the as-spun metal ion/polymer nanofibers to a metal oxide nanotube structure. When the as-spun nanofibers are heated in air, the presence of nitrate ions facilitates the initiation of the PVP oxidation. For comparison, the decomposition of pure PVP in air has generally been observed to take place at higher temperature and with a smaller heat release than that in the first exotherm observed in our DSC measurements, supporting the important role of the nitrate ion as an oxidant in the present case.22,24 As thermal decomposition and the volatilization of the decomposition products progresses, metal oxide clusters start to form on the fiber surface where metal ions are in contact with ambient air. (By way of comparison, it is worth noting that when an as-spun Fe(III) sample was heat treated in argon, pure metal formed instead of metal oxide.25) These metal oxides form small aggregates as the fiber diameter continues to decrease

Metal Oxide Nanotubes until they assemble into a rigid shell. The close association between the mass loss and fiber diameter reduction over this temperature range indicates that the fibers remain solid. This also suggests that the degradation of PVP into volatile decomposition products occurs at the surface where Fe2O3 nanocrystals accumulate and aggregate, eventually leading to the formation of a rigid R-Fe2O3 network. Once the oxide shell forms, the fiber diameter stops decreasing, while the removal of additional PVP drives the remaining metal ions toward the surface. PVP is completely removed at around 475 °C where the formation of oxide nanotubes is completed. As shown in the TEM images in Figure 4, the core-shell structure initially forms between 385 and 410 °C and continues to develop until the nanotubes form at 475 °C. The results of the measurements described above strongly suggest that a necessary condition for nanotube formation is that an oxide shell must form before the removal of PVP is complete (i.e., prior to the second exothermic heat flow peak arising from the final decomposition of PVP in a DSC measurement). This observation provides a useful guide in designing an appropriate heat treatment for reliably and reproducibly converting as-spun fibers to nanotubes. In particular, the heating rate and thermal anchoring of samples must be designed to minimize sample overheating due to the considerable heat produced during the exothermic oxidation of PVP. If the sample heating rate is too rapid and/or if the thermal anchoring of the sample to its environment is insufficient, sample overheating can lead to the overlap of the oxide shell formation and polymer removal processes. In these cases we have observed poor quality nanotubes or, in extreme cases, oxide nanofibers rather than nanotubes. Sample overheating can be directly observed in a (non-powercompensated) DSC as a forward tilt in the DSC signal as a function of the programmed furnace temperature. This effect is shown in Figure 2b. The single forward tilting exothermic DSC signal indicates that the sample temperature is greater than the programmed furnace temperature, and the corresponding abrupt mass loss suggests that the decomposition of PVP and oxide formation occurs at essentially the same time. Samples with very poor quality in terms of nanotube morphology formed under these conditions (see TEM images in Figure S3) reinforces the importance of a well-controlled heat treatment. Even in the absence of sample overheating, it may not be possible to form nanotubes if the kinetic or thermodynamic processes leading to oxide formation occur at temperatures above the PVP removal temperature. For example, it has previously been reported that the formation temperature of NiO following the thermal decomposition of Ni(NO3)2 · 6H2O is about 100 °C higher than the formation temperature of Fe2O3 from Fe(NO3)3 · 9H2O.26 The fact that we have not been able to form NiO nanotubes, even with a well-controlled heat treatment, is consistent with the proposed nanotube formation mechanism. The nanotube formation process described above differs in several important ways from the mechanism described in ref 7. We clearly identified, by revealing the sample microstructures at different temperatures along the DSC-TGA curve, the importance of the oxide formation before the complete removal of the polymer carrier in order to form the oxide tubes. The argument in ref 7 that the accumulation of metal ions on the surface of nanofibers was due to the evaporation of the precursor solvent was not observed in our experiment. In fact, the TGDTA measurement in ref 7 closely resembles to Figure 2B in our experiment, where a sharp exothermal peak was ac-

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Figure 6. Room temperature hysteresis loops of R-Fe2O3 nanofibers and nantubes.

companied by a big mass loss at around 387 °C and reliable nanotube formation is not expected in this condition. Bulk R-Fe2O3 is a canted antiferromagnetic material at room temperature with a small net magnetic moment of 0.002 µB (or 0.4 emu/g)27 and Ne¨el temperature of 948 K.28 The moments in the a-b plane are aligned and canted along the c-axis due to the spin-orbital coupling promoted by the low symmetry of the cation sites. In comparison with bulk R-Fe2O3, it can be seen from the room temperature hysteresis loops shown in Figure 6 that solid nanofibers and nanotubes of R-Fe2O3 exhibit significantly larger magnetizations of about 0.75 and 1.50 emu/g, respectively, at an applied field of 1 T. These magnetization values are similar to those previously reported for R-Fe2O3 nanorods (15-25 nm diameter) and uniform long straight chains of nanoparticles (19.8 nm average diameter), respectively,29,30 but larger than those of individual nanoparticles which are typically less than about 0.5 emu/g.31-35 Unlike conventional magnetic nanowires,36 the hysteresis loops measured with the applied field parallel and perpendicular to the fiber plane are essentially identical, indicating the absence of a net or effective shape anisotropy in these materials. The ferromagnetic behaviors in nanotubes are distinctively different from that of nanofibers, which may suggest a magnetic interaction between the spins on the inner and outer surfaces of the nanotubes. Recent work by Dr. Nielsch’s group and by Han et al. revealed interesting magnetic properties of ferrimagnetic and ferromagnetic nanotubes, respectively.37,38 Notably, two distinct magnetic reversal mode was observed in the magnetic nanotubes. Very different magnetic properties are generally expected in antiferromagnetic, ferrimagentic, and ferromagnetic materials. We currently do not have sufficient data to verify if those two reversal modes apply to antiferromagnetic nanotubes, and more work is underway to further explore their magnetic properties. 4. Conclusions In this article we have described a novel process for the preparation of metal oxide nanotubes based on an appropriate heat treatment of electrospun polymer/metal-salt precursor nanofibers. A systematic series of measurements have identified several distinct physical and chemical processes leading to the formation of the desired metal oxide nanotubes. These processes include the decomposition of nitrate ions which leads to the partial decomposition and volatilization of the polymer component in as-spun nanofibers, the formation of metal oxide aggregates on the nanofiber surface and their aggregation into a continuous shell, and the removal of remaining polymer. Precursor selection and thermal treatment conditions are dis-

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cussed based on the nanotube formation mechanism, which could provide useful guidance for nanotube fabrication of other materials. Ferromagnetic-like behavior has been observed in both R-Fe2O3 nanofibers and nanotubes and is more pronounced in the case of nanotubes, due most likely to the interaction among spins on the surfaces. Acknowledgment. This work was supported by ONR Grant N00014-08-1-0423. Any opinions, findings, and conclusions or recommendations expressed in this material are those of the author(s) and do not necessarily reflect the views of the Office of Naval Research. Supporting Information Available: XRD patterns of R-Fe2O3 and Co3O4 nanotubes, DSC scans of sample heated to 385 °C, then cooled down, and heated again to 500 °C, TEM image of overheated R-Fe2O3 sample. This material is available free of charge via the Internet at http://pubs.acs.org. References and Notes (1) Iijima, S. Nature 1991, 354, 56–58. (2) Jorio, A.; Dresselhaus, G.; Dresselhaus, M. S. Carbon Nanotubes: AdVanced Topics in the Synthesis, Structure, Properties and Applications; Springer-Verlag: Berlin, 2008. (3) Tenne, R.; Remsˇkar, M.; Enyashin, A.; Seifert, G. Carbon Nanotubes 2008, 631–671. (4) Fan, H. J.; Go¨sele, U.; Zacharias, M. Small 2007, 3, 1660–1671. (5) Li, D.; Xia, Y. Nano Lett. 2004, 4, 933–938. (6) Chen, W. S.; Huang, D. A.; Chen, H. C.; Shie, T. Y.; Hsieh, C. H.; Liao, J. D.; Kuo, C. Cryst. Growth Des. 2009, 9, 4070–4077. (7) Cui, Q.; Dong, X.; Wang, J.; Li, M. J. Rare Earths 2008, 26, 664– 669. (8) Bognitzki, M.; Czado, W.; Frese, T.; Schaper, A.; Hellwig, M.; Steinhart, M.; Greiner, A.; Wendorff, J. H. AdV. Mater. 2001, 13, 70–72. (9) Teo, W. E.; Ramakrishna, S. Nanotechnology 2006, 17, R89–R106. (10) Fridrikh, S. V.; Yu, J. H.; Brenner, M. P.; Rutledge, G. C. Phys. ReV. Lett. 2003, 90, 4. (11) Li, D.; Wang, Y. L.; Xia, Y. Nano Lett. 2003, 3, 1167–1171. (12) Zhou, F. L.; Gong, R. H.; Porat, I. J. Appl. Polym. Sci. 2010, 115, 2591–2598. (13) Qiu, Y. J.; Yu, J.; Rafique, J.; Yin, J.; Bai, X. D.; Wang, E. J. Phys. Chem. C 2009, 113, 11228–11234. (14) Guan, H. Y.; Shao, C. L.; Wen, S. B.; Chen, B.; Gong, J.; Yang, X. H. Inorg. Chem. Commun. 2003, 6, 1302–1303. (15) Li, D.; Xia, Y. Nano Lett. 2003, 3, 555–560.

Chen et al. (16) Li, D.; Herricks, T.; Xia, Y. Appl. Phys. Lett. 2003, 83, 4586– 4588. (17) CELREF is a component of the LMGP suite of programs for the interpretation of X-ray experiments developed by Jean Laugier and Bernard Bochu (ENSP/Laboratoire des Mate´riaux et du Ge´nie Physique, France) and can be obtained at http://www.inpg.fr/LMGP. (18) Gmelin, E.; Sarge, S. M. Pure Appl. Chem. 1995, 67 (11), 1789– 1800. (19) Villars, P.; Calvert, L. D.; Pearson, W. B. Pearson’s Handbook of Crystallographic Data for Intermetallic Phases; ASM International: Materials Park, OH, 1991. (20) Liu, M.; Yan, X.; Liu, H.; Yu, W. React. Funct. Polym. 2000, 44, 55–64. (21) Beamso, G.; Briggs, D. High Resolution XPS of Organic Polymers: The Scienta ESCA300 Database; American Chemical Society: Washington, DC, 1993; Vol. 70. (22) Peniche, C.; Zaldı´var, D.; Pazos, M.; Pa´z, S.; Bulay, A.; Roma´n, J. S. J. Appl. Polym. Sci. 2003, 50, 485–493. (23) Bianco, G.; Soldi, M. S.; Pinheiro, E. A.; Pires, A. T. N.; Gehlen, M. H.; Soldi, V. Polym. Degrad. Stab. 2003, 80, 567–574. (24) Du, Y. K.; Yang, P.; Mou, Z. G.; Hua, N. P.; Jiang, L. J. Appl. Polym. Sci. 2005, 99, 23–26. (25) Chen, X. Presented at the 11th Joint MMM-Intermag Conference, Washington, DC, Jan 2010. (26) Elmasry, M. A. A.; Gaber, A.; Khater, E. M. H. J. Therm. Anal. Calorim. 1998, 52, 489–495. (27) Flanders, P. J.; Remeika, J. P. Philos. Mag. 1965, 11, 1271–1288. (28) Morrish, A. H. Canted Antiferromagnetism: Hematite; World Scientific Publishing Co. Pte. Ltd.: Singapore, 1994. (29) Tang, B.; Wang, G.; Zhuo, L.; Ge, J.; Cui, L. Inorg. Chem. 2006, 45, 5196–5200. (30) Meng, L. R.; Chen, W.; Chen, C.; Zhou, H.; Peng, Q.; Li, Y. Cryst. Growth Des. 2010, 10, 479–482. (31) Yin, W.; Chen, X.; Cao, M.; Hu, C.; Wei, B. J. Phys. Chem. C 2009, 113, 15897–15903. (32) Jacob, J.; Abdul Khadar, M. J. Magn. Magn. Mater. 2010, 322, 614–621. (33) Zeng, S.; Tang, K.; Li, T.; Liang, Z.; Wang, D.; Wang, Y.; Qi, Y.; Zhou, W. J. Phys. Chem. C 2008, 112, 4836–4843. (34) Jing, Z.; Wu, S.; Zhang, S.; Huang, W.; Huang, Mater. Res. Bull. 2004, 39, 2057–2064. (35) Jing, Z. H.; Wu, S. H. Mater. Chem. Phys. 2005, 92, 600–603. (36) Kou, X.; Fan, X.; Zhu, H.; Xiao, J. Q. Appl. Phys. Lett. 2009, 94, 112509–3. (37) Bachmann, J.; Escrig, J.; Pitzschel, K.; Montero, J.; Jing, J.; Gorlitz, D.; Altbir, D.; Nielsch, K. J. Appl. Phys. 2009, 105, 07B521. (38) Han, X.; Shamaila, S.; Sharif, R.; Chen, J.; Liu, H.; Liu, D. AdV. Mater. 2009, 21, 4619–4624.

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