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polymer composites
ARTICLE IN PRESS
Journal of Magnetism and Magnetic Materials 280 (2004) 412–418
Magnetic behavior of iron and iron-oxide nanoparticle/polymer composites C. Bakera, S. Ismat Shaha,b,*, S.K. Hasanainc a
Department of Materials Science and Engineering, University of Delaware, 208 Dupont Hall, Newark, DE 19716, USA b Department of Physics and Astronomy, University of Delaware,Newark, DE, USA c Department of Physics, Quaid-i-Azam University, Islamabad, Pakistan Received 15 January 2004; received in revised form 25 March 2004 Available online 20 April 2004
Abstract An inert gas condensation technique was used to prepare nanometer-sized particles of metallic iron and iron oxide. The particles were passivated by the controlled oxidation of the particle surface leading to an Fe-oxide shell-Fe core structure. Nanoparticle–polymer composites were obtained by spin casting mixtures of nanoparticles and polymethylmethacrylate films. The magnetic properties of the nanoparticles compressed into pellets and dispersed in the composites were both studied. The particles were observed to exhibit increased coercivity and exchange bias. The exchange bias was observed to increase with oxide shell thickness. The magnetism in the nanoparticle composites was studied as a function of nanoparticle loading. It was observed that when the particles were dispersed into the nanocomposite the coercivity was increased, suggesting a heightened anisotropy barrier. Similarly, the magnetic relaxation results indicate that the composites exhibit significantly reduced relaxations through the entire temperature range, as compared to the compressed pellet. This observation supports the possibility of heightened anisotropy barriers due to reduced dipolar interactions. r 2004 Elsevier B.V. All rights reserved. PACS: 75.50.Tt; 75.30.Et; 75.75.+a Keywords: Inert gas condensation; Iron; Iron-oxide; Polymer nanocomposites
1. Introduction In the single domain region, nanoparticles may exhibit phenomena such as enhanced coercivity, *Corresponding author. Department of Materials Science and Engineering, University of Delaware, 208 Dupont Hall, Newark, DE 19716, USA Tel.: +1-302-831-1618; fax: +1-302831-4545. E-mail address: [email protected] (S. Ismat Shah).
reduced saturation magnetization, and superparamagnetism. These phenomena, among others, have made these materials the subject of intense study for technological applications as well as theoretical research. An interesting aspect of magnetic nanoparticle research has been borne out of the need to passivate the particles. The high surface energy of metal nanoparticles often requires that the surface be passivated with a protective layer of another species. Deposition of
0304-8853/$ - see front matter r 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jmmm.2004.03.037
ARTICLE IN PRESS C. Baker et al. / Journal of Magnetism and Magnetic Materials 280 (2004) 412–418
an oxide layer on the nanoparticle is often used to attain passivation. The oxide passivation layer generally formed on iron is either g-Fe2O3 or Fe3O4, both of which are ferrimagnetic species. Similarly, Co nanoparticles are often passivated with CoO, which is antiferromagnetic (AFM) [1]. Materials with a ferromagnetic (FM)/antiferromagnetic (AFM) or ferrimagnetic interface have been found to exhibit an exchange bias when field cooled (FC) below the Ne! el temperature of the AFM component (see review by Berkowitz and Takano [2]). The exchange bias is represented by a shift in the magnetization hysteresis loop in the direction of the cooling field, and is generally defined as Heb ¼ jH1FC þ H2FC j=2 [1]. Exchange bias has been reported to be the result of exchange anisotropy between the layers. Meiklejohn and Bean first discovered this phenomenon in experiments with Co/CoO core/shell nanoparticles [3]. We have synthesized iron-oxide passivated iron nanoparticles using an inert gas condensation (IGC) technique. The iron nanoparticles were synthesized with varying diameters, and oxide passivation layer thickness. In this report, we investigate the oxide passivation layer effects on the magnetic properties of the iron nanoparticles. The iron nanoparticles were also uniformly dispersed into spin cast polymethylmethacrylate (PMMA) films in varying concentrations. The coercivity and thermoremanent magnetization of the samples were studied for these films and compared to that of the particles compressed into pellets. We have observed that effects, such as exchange bias, are the result of an enhancedanisotropy effect occurring in the oxide passivation layer. The anisotropy barrier is however affected by inter-particle dipolar interactions, which tend to reduce the barrier. It is postulated that as the particles are dispersed into the polymer matrix, dipolar interactions are reduced, and thus anisotropy barriers are heightened in the samples.
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high-purity Fe wire toward a resistively heated Al2O3-coated W boat where it melted and evaporated. The resulting Fe vapor was entrained and cooled in a flowing stream of He gas where individual Fe atoms coalesced into small particles. These particles were deposited on a filter from which they were occasionally dislodged and collected in a removable container mounted to the bottom of the vacuum chamber [6]. Fe/Fe-oxide core/shell nanoparticles were prepared by slowly exposing the collected Fe nanoparticles to atmospheric oxygen. The average particle size was 1372 nm. Oxide shell thickness was controllable by controlling the exposure rate after particle growth. The nanoparticles were dispersed by sonication into a solution of PMMA powder and acetone. The solution was then spin cast until it hardened into a thin film. The particles were dispersed in three concentrations of 1.5, 1.0, and 0.5 mg of Fe in 100 mg of PMMA powder. The structure of the Fe/Fe-oxide samples was studied by X-ray diffraction (XRD) and transmission electron microscopy (TEM). In the former case, diffraction patterns were obtained over a wide angular range, from 2y ¼ 10 to 110 . The diffraction peaks were used in conjunction with the Scherrer formula, to estimate the mean particle size (after subtracting the measured natural linewidth in quadrature) [5]. TEM samples were prepared by suspending a small amount of the Fe nanoparticles in acetone, placing a small droplet of the suspension onto a TEM grid, and allowing the solvent to evaporate. Magnetic measurements were carried out in a DC extraction magnetometer in a Physical Properties Measurement System apparatus by Quantum Design Corporation. Magnetization hysteresis loops were measured to fields of 2 T. Magnetic relaxation measurements were performed after cooling in a field of 5000 Oe from 300 K to the appropriate temperature at a rate of 2 K/min. The field was removed after waiting for 5 min at the desired temperature and the moment was measure for 40 min.
2. Experimental 3. Results and discussion All of the Fe nanoparticles studied in this work were prepared by an ICG technique [4,5]. The procedure consisted of continuously feeding a
XRD patterns for all samples indicated diffraction peaks from both a-Fe and a spinel phase
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Fe-oxide. In Fig. 1, XRD patterns for three samples are displayed. The designation Ir (integrated intensity ratio) refers to the normalized ratio of iron-oxide shell g-Fe2O3 (3 1 1) peak (or Fe3O4) intensity to Fe (1 1 0) core peak intensity, for example Ir 0.49 is a sample with oxide shell peak intensity almost half of that of the Fe core peak intensity. The Fe peaks are indexed. The broad Fe-oxide (spinel) peaks correspond to either g-Fe2O3 or Fe3O4 where the predominant peak is the (3 1 1) peak found at 2yB35 : The peak positions for the two oxide phases are situated too close to explicitly determine which phase is present by this technique [7]. The intensity of these broad peaks varies as the oxide shell thickness is varied. All of the Fe core particle sizes are comparable in this work. The distribution in core
size for any sample is greater than the error in XRD peak broadening measurements. A TEM lattice image for a typical sample is given in Fig. 1b showing a darker core corresponding to a-Fe surrounded by the light shell structure of the spinel phase Fe oxide. The Fe core is single crystal, as confirmed by the high-resolution lattice imaging. The difference in the shades is related to the nonuniform thickness of the sample. The oxide shell is polycrystalline, as evident by the various orientations of the lattice planes. The shell in the figure is approximately 2–3 nm thick. Fig. 2 illustrates the effects of cooling sample Ir 0.49 in a magnetic field. In Fig. 2a, the room temperature magnetization loop is given. The coercivity was found to be 440 Oe; therefore, the sample does not exhibit superparamagnetism. The
Fig. 1. (a) XRD patterns for three Fe/Fe-oxide samples. The a-Fe peaks are indexed. (b) TEM micrograph of an Fe-oxide-coated Fe nanoparticle.
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sample was then cooled in a field of 2 T, from 300 to 5 K. Fig. 2b shows the magnetization loop after the FC. It is evident that the coercivity has increased considerably and a loop shift or exchange bias of 277 Oe is present. The increased coercivity is the result of increased anisotropy energy barriers with decreasing temperatures. The loop shift occurs as the sample is cooled below a blocking temperature, Tb; characteristic of the oxide layers. The effect occurs when the sample is cooled from above Tb of the ferrimagnetic shell, where the moments are oriented at random, and below TC of the FM core, where the moments are aligned with the field. As the sample is cooled below the Tb ; the ferrimagnetic moments begin to align with the field and consequently with the FM moments and undergo an exchange interaction with the moments of the FM core. If the field is reversed, the FM moments tend to rotate, but their reversal is prevented due to their coupling with the ferrimagnetic moments that encounter a large anisotropy in the oxide layers. In this case, there is a large energy barrier for the moments to overcome if they are to align with the direction opposite to that of the cooling field direction. The result is a loop shift [8]. In Fig. 3, the exchange bias as a function of temperature is plotted for three samples. Exchange bias is observed as the temperature is lowered below B80 K, which may also correspond to a Tb for the ferrimagnetic shell. It is evident that the exchange bias increases as the temperature is decreased for each sample. It is also evident that the exchange bias increases with oxide content in the samples. Sample Ir 0.49 has the greatest exchange bias at 5 K, which corresponds to the highest ratio of oxide to Fe peak intensity. This suggests that the thicker oxide layer has heightened anisotropy energy. This is similar to the finite size volume effect first presented by Parker et al. [9] from experiments performed on g-Fe2O3 nanoparticles. A similar dependence, i.e., increasing exchange bias with oxide thickness, was also found by Del Bianco et al. [10] for oxide passivated Fe nanoparticles of 6–15 nm. They also observed values comparable to our results of the exchange bias. In contrast, Peng et al. [11] studied samples
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that were both intentionally passivated in situ and samples that were not intentionally passivated but did acquire an oxide passivation layer upon removal from the growth chamber. They found the exchange bias to be as high as 1200 Oe after FC to 5 K for a 9 nm diameter sample of in situ passivated particles. This value is B4 times greater than the highest value found in our samples. However, a sample in which the only surface oxidation was due to the exposure to the atmosphere was found to have a value for exchange bias similar to our samples. This suggests that the intentionally oxidized samples in their work may have had considerably more oxide coating than our samples, and that the exchange bias increases dramatically for large amounts of oxidation in samples smaller than 10 nm. Fig. 4 shows a Scanning Electron Micrograph (SEM) of a Fe/PMMA sample. Although the majority of the particles are situated below the surface of the film, and thus are not visible in the image, some appear on the surface. It can also be seen that some agglomeration of the particles has occurred. The coercivity, Hc vs. temperature for sample Ir 0.41, both packed into a disc and dispersed in a PMMA film, is given in Fig. 5a. For all
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temperatures, Hc is lower for the sample packed into a disc than when it is dispersed in the nanocomposite film. Suppressed coercivity due to inter-particle interaction has been observed in a Monte Carlo modeling study [12]. The lowered Hc value for the packed sample may be the result of a lowered anisotropy barrier due to dipolar interactions. When the average inter-particle distances are increased, as is the case when they are dispersed in the film, the anisotropy barrier effectively increases. In Fig 5b, Heb for Fe dispersed in PMMA films for three concentrations is given. It is seen that at 5 K the exchange bias decreases with the increase in concentration of the Fe in the films.
This is also the result of the inter-particle dipolar interactions effectively increasing in strength and becoming dominant over the intra-particle exchange interaction. Thermoremanent magnetic relaxations for sample Ir 0.41, both pressed into a disc and dispersed in a PMMA film, are plotted in Fig. 6. The samples were cooled in a field of 5000 Oe and the field was then removed. The relaxation is measured at 6 and 20 K. Initially, the moments in the sample are aligned with the field, but as the field is removed a number of moments attain thermal energy to overcome the anisotropy barrier energy and arrange themselves at random. The sample that was compressed into a disc has a much faster relaxation than the sample dispersed into the polymer. This is due to the lowered anisotropy barrier for the particles pressed into the pellet. The magnetization (m) vs. time (t) for the powder pressed into a pellet can be modeled with a logarithmic function as t m ¼ m0 1 S log 1 þ ; ð1Þ t where S is the magnetic viscosity, m0 is the initial value of the moment (at t ¼ 0), and t is a time constant obtained from the fit (typically about 40 s). The magnetic relaxation being represented by a logarithmic function is common for a system of particles with a distribution of barrier energies. Here the magnetic viscosity S; which is the rate of relaxation of the moments, is related to the distribution of the energy barriers in the sample [13].
Fig. 4. SEM micrograph of Fe/Fe-oxide nanoparticles dispersed in a PMMA film.
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Fig. 5. (a) Coercivity Hc for sample Ir 0.41 both pressed into a disc and dispersed in a PMMA film. (b) Exchange bias Heb as a function of temperature for sample Ir 0.41 dispersed in varying concentrations in PMMA nanocomposites.
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Fig. 6. Theromoremanent relaxation in sample Ir 0.41 pressed into a pellet (left) and dispersed into a polymer, (right) at 6 K, (top) and at 20 K (bottom).
S is determined from the fit of the data to Eq. (1) to be equal to 0.0008 at 6 K and 0.0025 at 20 K. The temperature-dependent increase in relaxation is expected since the available energy (kT) for excitation of the moments over the anisotropy barrier increases with temperature. The relaxation rate depends on the ratios of the energies
R¼
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ð2Þ
where K is the anisotropy constant for the material, V is the particle volume, and k is Boltzman’s constant. However, it can also be noted that there is virtually no relaxation in the Fe/PMMA sample at both temperatures. This is due to the heightened anisotropy constant K for the case where the particles are separated too far for dipolar interactions to be effective. It is only at 40 K where relaxation begins to occur in the Fe/ PMMA samples. In this case, the magnetic viscosity has a value of 0.0034.
4. Conclusion Fe-oxide/Fe core shell nanoparticles were synthesized in varying sizes. The exchange bias of the particles was found to depend on the amount of Fe-oxide coating the samples. This effect occurs as a result of a heightened anisotropy barrier that results from the ferrimagnetic oxide shell and the exchange interaction with the Fe core. The nanoparticles were dispersed into PMMA films in varying concentrations. It is observed that interparticle effects such as dipolar interaction are reduced as the particles are separated from each other. The result is an increase in coercivity, and a much slower relaxation rate in thermoremanent magnetization experiments.
Acknowledgements The authors would like to acknowledge Professor K.M. Unruh for the use of his XRD apparatus and help with the diffraction analysis. The authors
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would also like to acknowledge the NSF INT0138151 US-Pakistan Cooperative Research Grant.
References [1] D.L. Peng, K. Sumiyama, T. Hihara, S. Yamamuro, T.J. Konno, Phys. Rev. B 61 (2000) 3103. [2] A.E. Berkowitz, K. Takano, J. Magn. Magn. Mater. 200 (1999) 552. [3] W.H. Meiklejohn, C.P. Bean, Phys. Rev. 102 (1956) 1413; W.H. Meiklejohn, C.P. Bean, Phys. Rev. 105 (1957) 904. [4] C.G. Granqvist, R.A. Buhrman, J. Appl. Phys. 47 (1976) 2200. [5] C. Baker, S. Ismat Shah, S.K. Hasanain, B. Ali, L. Shah, G. Li, T. Ekiert, K.M. Unruh, Mater. Res. Soc. Symp. Proc. 746 (2003) 201.
[6] H. Shen, B. Gunther, H. Sch.afer, Z. Li, Zh. Qi, Scr. Metall. Mater. 32 (1995) 1677. [7] E. Bonetti, L. Savini, A. Deriu, G. Albanese, J. Moya, J. Magn. Magn. Mater. 262 (2003) 132. [8] B. Martinez, X. Obradors, Ll. Balcells, A. Rouanet, C. Monty, Phys. Rev. Lett. 80 (1998) 181. [9] F.T. Parker, M.W. Foster, D.T. Margulies, A.E. Berkowitz, Phys. Rev. B 47 (1993) 7885. [10] L. Del Bianco, D. Fiorani, A.M. Testa, E. Bonetti, L. Savini, S. Signoretti, Phys. Rev B 66 (2002) 174418. [11] D.L. Peng, T. Hihara, K. Sumiyama, H. Morikawa, J. Appl. Phys. 92 (2002) 3075. [12] D. Kechrakos, K.N. Trohidou, J. Magn. Magn. Mater. 262 (2003) 107. [13] K.D. Humfield, A.K. Giri, S.A. Majetich, E.L. Venturini, IEEE Trans. 37 (2001) 2194.
Journal of Magnetism and Magnetic Materials 280 (2004) 412–418
Magnetic behavior of iron and iron-oxide nanoparticle/polymer composites C. Bakera, S. Ismat Shaha,b,*, S.K. Hasanainc a
Department of Materials Science and Engineering, University of Delaware, 208 Dupont Hall, Newark, DE 19716, USA b Department of Physics and Astronomy, University of Delaware,Newark, DE, USA c Department of Physics, Quaid-i-Azam University, Islamabad, Pakistan Received 15 January 2004; received in revised form 25 March 2004 Available online 20 April 2004
Abstract An inert gas condensation technique was used to prepare nanometer-sized particles of metallic iron and iron oxide. The particles were passivated by the controlled oxidation of the particle surface leading to an Fe-oxide shell-Fe core structure. Nanoparticle–polymer composites were obtained by spin casting mixtures of nanoparticles and polymethylmethacrylate films. The magnetic properties of the nanoparticles compressed into pellets and dispersed in the composites were both studied. The particles were observed to exhibit increased coercivity and exchange bias. The exchange bias was observed to increase with oxide shell thickness. The magnetism in the nanoparticle composites was studied as a function of nanoparticle loading. It was observed that when the particles were dispersed into the nanocomposite the coercivity was increased, suggesting a heightened anisotropy barrier. Similarly, the magnetic relaxation results indicate that the composites exhibit significantly reduced relaxations through the entire temperature range, as compared to the compressed pellet. This observation supports the possibility of heightened anisotropy barriers due to reduced dipolar interactions. r 2004 Elsevier B.V. All rights reserved. PACS: 75.50.Tt; 75.30.Et; 75.75.+a Keywords: Inert gas condensation; Iron; Iron-oxide; Polymer nanocomposites
1. Introduction In the single domain region, nanoparticles may exhibit phenomena such as enhanced coercivity, *Corresponding author. Department of Materials Science and Engineering, University of Delaware, 208 Dupont Hall, Newark, DE 19716, USA Tel.: +1-302-831-1618; fax: +1-302831-4545. E-mail address: [email protected] (S. Ismat Shah).
reduced saturation magnetization, and superparamagnetism. These phenomena, among others, have made these materials the subject of intense study for technological applications as well as theoretical research. An interesting aspect of magnetic nanoparticle research has been borne out of the need to passivate the particles. The high surface energy of metal nanoparticles often requires that the surface be passivated with a protective layer of another species. Deposition of
0304-8853/$ - see front matter r 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jmmm.2004.03.037
ARTICLE IN PRESS C. Baker et al. / Journal of Magnetism and Magnetic Materials 280 (2004) 412–418
an oxide layer on the nanoparticle is often used to attain passivation. The oxide passivation layer generally formed on iron is either g-Fe2O3 or Fe3O4, both of which are ferrimagnetic species. Similarly, Co nanoparticles are often passivated with CoO, which is antiferromagnetic (AFM) [1]. Materials with a ferromagnetic (FM)/antiferromagnetic (AFM) or ferrimagnetic interface have been found to exhibit an exchange bias when field cooled (FC) below the Ne! el temperature of the AFM component (see review by Berkowitz and Takano [2]). The exchange bias is represented by a shift in the magnetization hysteresis loop in the direction of the cooling field, and is generally defined as Heb ¼ jH1FC þ H2FC j=2 [1]. Exchange bias has been reported to be the result of exchange anisotropy between the layers. Meiklejohn and Bean first discovered this phenomenon in experiments with Co/CoO core/shell nanoparticles [3]. We have synthesized iron-oxide passivated iron nanoparticles using an inert gas condensation (IGC) technique. The iron nanoparticles were synthesized with varying diameters, and oxide passivation layer thickness. In this report, we investigate the oxide passivation layer effects on the magnetic properties of the iron nanoparticles. The iron nanoparticles were also uniformly dispersed into spin cast polymethylmethacrylate (PMMA) films in varying concentrations. The coercivity and thermoremanent magnetization of the samples were studied for these films and compared to that of the particles compressed into pellets. We have observed that effects, such as exchange bias, are the result of an enhancedanisotropy effect occurring in the oxide passivation layer. The anisotropy barrier is however affected by inter-particle dipolar interactions, which tend to reduce the barrier. It is postulated that as the particles are dispersed into the polymer matrix, dipolar interactions are reduced, and thus anisotropy barriers are heightened in the samples.
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high-purity Fe wire toward a resistively heated Al2O3-coated W boat where it melted and evaporated. The resulting Fe vapor was entrained and cooled in a flowing stream of He gas where individual Fe atoms coalesced into small particles. These particles were deposited on a filter from which they were occasionally dislodged and collected in a removable container mounted to the bottom of the vacuum chamber [6]. Fe/Fe-oxide core/shell nanoparticles were prepared by slowly exposing the collected Fe nanoparticles to atmospheric oxygen. The average particle size was 1372 nm. Oxide shell thickness was controllable by controlling the exposure rate after particle growth. The nanoparticles were dispersed by sonication into a solution of PMMA powder and acetone. The solution was then spin cast until it hardened into a thin film. The particles were dispersed in three concentrations of 1.5, 1.0, and 0.5 mg of Fe in 100 mg of PMMA powder. The structure of the Fe/Fe-oxide samples was studied by X-ray diffraction (XRD) and transmission electron microscopy (TEM). In the former case, diffraction patterns were obtained over a wide angular range, from 2y ¼ 10 to 110 . The diffraction peaks were used in conjunction with the Scherrer formula, to estimate the mean particle size (after subtracting the measured natural linewidth in quadrature) [5]. TEM samples were prepared by suspending a small amount of the Fe nanoparticles in acetone, placing a small droplet of the suspension onto a TEM grid, and allowing the solvent to evaporate. Magnetic measurements were carried out in a DC extraction magnetometer in a Physical Properties Measurement System apparatus by Quantum Design Corporation. Magnetization hysteresis loops were measured to fields of 2 T. Magnetic relaxation measurements were performed after cooling in a field of 5000 Oe from 300 K to the appropriate temperature at a rate of 2 K/min. The field was removed after waiting for 5 min at the desired temperature and the moment was measure for 40 min.
2. Experimental 3. Results and discussion All of the Fe nanoparticles studied in this work were prepared by an ICG technique [4,5]. The procedure consisted of continuously feeding a
XRD patterns for all samples indicated diffraction peaks from both a-Fe and a spinel phase
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Fe-oxide. In Fig. 1, XRD patterns for three samples are displayed. The designation Ir (integrated intensity ratio) refers to the normalized ratio of iron-oxide shell g-Fe2O3 (3 1 1) peak (or Fe3O4) intensity to Fe (1 1 0) core peak intensity, for example Ir 0.49 is a sample with oxide shell peak intensity almost half of that of the Fe core peak intensity. The Fe peaks are indexed. The broad Fe-oxide (spinel) peaks correspond to either g-Fe2O3 or Fe3O4 where the predominant peak is the (3 1 1) peak found at 2yB35 : The peak positions for the two oxide phases are situated too close to explicitly determine which phase is present by this technique [7]. The intensity of these broad peaks varies as the oxide shell thickness is varied. All of the Fe core particle sizes are comparable in this work. The distribution in core
size for any sample is greater than the error in XRD peak broadening measurements. A TEM lattice image for a typical sample is given in Fig. 1b showing a darker core corresponding to a-Fe surrounded by the light shell structure of the spinel phase Fe oxide. The Fe core is single crystal, as confirmed by the high-resolution lattice imaging. The difference in the shades is related to the nonuniform thickness of the sample. The oxide shell is polycrystalline, as evident by the various orientations of the lattice planes. The shell in the figure is approximately 2–3 nm thick. Fig. 2 illustrates the effects of cooling sample Ir 0.49 in a magnetic field. In Fig. 2a, the room temperature magnetization loop is given. The coercivity was found to be 440 Oe; therefore, the sample does not exhibit superparamagnetism. The
Fig. 1. (a) XRD patterns for three Fe/Fe-oxide samples. The a-Fe peaks are indexed. (b) TEM micrograph of an Fe-oxide-coated Fe nanoparticle.
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ARTICLE IN PRESS C. Baker et al. / Journal of Magnetism and Magnetic Materials 280 (2004) 412–418 300 Ir 0.20 Ir 0.41 Ir 0.49
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sample was then cooled in a field of 2 T, from 300 to 5 K. Fig. 2b shows the magnetization loop after the FC. It is evident that the coercivity has increased considerably and a loop shift or exchange bias of 277 Oe is present. The increased coercivity is the result of increased anisotropy energy barriers with decreasing temperatures. The loop shift occurs as the sample is cooled below a blocking temperature, Tb; characteristic of the oxide layers. The effect occurs when the sample is cooled from above Tb of the ferrimagnetic shell, where the moments are oriented at random, and below TC of the FM core, where the moments are aligned with the field. As the sample is cooled below the Tb ; the ferrimagnetic moments begin to align with the field and consequently with the FM moments and undergo an exchange interaction with the moments of the FM core. If the field is reversed, the FM moments tend to rotate, but their reversal is prevented due to their coupling with the ferrimagnetic moments that encounter a large anisotropy in the oxide layers. In this case, there is a large energy barrier for the moments to overcome if they are to align with the direction opposite to that of the cooling field direction. The result is a loop shift [8]. In Fig. 3, the exchange bias as a function of temperature is plotted for three samples. Exchange bias is observed as the temperature is lowered below B80 K, which may also correspond to a Tb for the ferrimagnetic shell. It is evident that the exchange bias increases as the temperature is decreased for each sample. It is also evident that the exchange bias increases with oxide content in the samples. Sample Ir 0.49 has the greatest exchange bias at 5 K, which corresponds to the highest ratio of oxide to Fe peak intensity. This suggests that the thicker oxide layer has heightened anisotropy energy. This is similar to the finite size volume effect first presented by Parker et al. [9] from experiments performed on g-Fe2O3 nanoparticles. A similar dependence, i.e., increasing exchange bias with oxide thickness, was also found by Del Bianco et al. [10] for oxide passivated Fe nanoparticles of 6–15 nm. They also observed values comparable to our results of the exchange bias. In contrast, Peng et al. [11] studied samples
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that were both intentionally passivated in situ and samples that were not intentionally passivated but did acquire an oxide passivation layer upon removal from the growth chamber. They found the exchange bias to be as high as 1200 Oe after FC to 5 K for a 9 nm diameter sample of in situ passivated particles. This value is B4 times greater than the highest value found in our samples. However, a sample in which the only surface oxidation was due to the exposure to the atmosphere was found to have a value for exchange bias similar to our samples. This suggests that the intentionally oxidized samples in their work may have had considerably more oxide coating than our samples, and that the exchange bias increases dramatically for large amounts of oxidation in samples smaller than 10 nm. Fig. 4 shows a Scanning Electron Micrograph (SEM) of a Fe/PMMA sample. Although the majority of the particles are situated below the surface of the film, and thus are not visible in the image, some appear on the surface. It can also be seen that some agglomeration of the particles has occurred. The coercivity, Hc vs. temperature for sample Ir 0.41, both packed into a disc and dispersed in a PMMA film, is given in Fig. 5a. For all
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temperatures, Hc is lower for the sample packed into a disc than when it is dispersed in the nanocomposite film. Suppressed coercivity due to inter-particle interaction has been observed in a Monte Carlo modeling study [12]. The lowered Hc value for the packed sample may be the result of a lowered anisotropy barrier due to dipolar interactions. When the average inter-particle distances are increased, as is the case when they are dispersed in the film, the anisotropy barrier effectively increases. In Fig 5b, Heb for Fe dispersed in PMMA films for three concentrations is given. It is seen that at 5 K the exchange bias decreases with the increase in concentration of the Fe in the films.
This is also the result of the inter-particle dipolar interactions effectively increasing in strength and becoming dominant over the intra-particle exchange interaction. Thermoremanent magnetic relaxations for sample Ir 0.41, both pressed into a disc and dispersed in a PMMA film, are plotted in Fig. 6. The samples were cooled in a field of 5000 Oe and the field was then removed. The relaxation is measured at 6 and 20 K. Initially, the moments in the sample are aligned with the field, but as the field is removed a number of moments attain thermal energy to overcome the anisotropy barrier energy and arrange themselves at random. The sample that was compressed into a disc has a much faster relaxation than the sample dispersed into the polymer. This is due to the lowered anisotropy barrier for the particles pressed into the pellet. The magnetization (m) vs. time (t) for the powder pressed into a pellet can be modeled with a logarithmic function as t m ¼ m0 1 S log 1 þ ; ð1Þ t where S is the magnetic viscosity, m0 is the initial value of the moment (at t ¼ 0), and t is a time constant obtained from the fit (typically about 40 s). The magnetic relaxation being represented by a logarithmic function is common for a system of particles with a distribution of barrier energies. Here the magnetic viscosity S; which is the rate of relaxation of the moments, is related to the distribution of the energy barriers in the sample [13].
Fig. 4. SEM micrograph of Fe/Fe-oxide nanoparticles dispersed in a PMMA film.
350
2000
0.5 mg Fe
Ir 0.41 300
Ir 0.41/PMMA
1.0 mg Fe
1600
1.5 mg Fe
1200
H eb (Oe)
Hc (Oe)
250
800
200 150 100
400
0
(a)
50 0 0
100
200 300 Temperature, T (K)
0
400
(b)
20
40
60
80
100
Temperature, T (K)
Fig. 5. (a) Coercivity Hc for sample Ir 0.41 both pressed into a disc and dispersed in a PMMA film. (b) Exchange bias Heb as a function of temperature for sample Ir 0.41 dispersed in varying concentrations in PMMA nanocomposites.
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0.153 6K
0.00478
Ir 0.41
Ir 0.41/PMMA
6K
M (emu)
M (emu)
0.1528 0.1526 0.1524
0.00472
0.00466 0.1522 0.0046
0.152
0
500 1000 Time (sec)
1500
0
2000
1000 2000 Time (sec)
3000
0.1435 0.00428
Ir 0.41
20 K
Ir 0.41/PMMA
20 K
M (emu)
M (emu)
0.143
0.1425
0.00422
0.00416
0.142
0.1415
0
1000 2000 Time (sec)
3000
0.0041
0
1000 2000 Time (sec)
3000
Fig. 6. Theromoremanent relaxation in sample Ir 0.41 pressed into a pellet (left) and dispersed into a polymer, (right) at 6 K, (top) and at 20 K (bottom).
S is determined from the fit of the data to Eq. (1) to be equal to 0.0008 at 6 K and 0.0025 at 20 K. The temperature-dependent increase in relaxation is expected since the available energy (kT) for excitation of the moments over the anisotropy barrier increases with temperature. The relaxation rate depends on the ratios of the energies
R¼
kT ; KV
ð2Þ
where K is the anisotropy constant for the material, V is the particle volume, and k is Boltzman’s constant. However, it can also be noted that there is virtually no relaxation in the Fe/PMMA sample at both temperatures. This is due to the heightened anisotropy constant K for the case where the particles are separated too far for dipolar interactions to be effective. It is only at 40 K where relaxation begins to occur in the Fe/ PMMA samples. In this case, the magnetic viscosity has a value of 0.0034.
4. Conclusion Fe-oxide/Fe core shell nanoparticles were synthesized in varying sizes. The exchange bias of the particles was found to depend on the amount of Fe-oxide coating the samples. This effect occurs as a result of a heightened anisotropy barrier that results from the ferrimagnetic oxide shell and the exchange interaction with the Fe core. The nanoparticles were dispersed into PMMA films in varying concentrations. It is observed that interparticle effects such as dipolar interaction are reduced as the particles are separated from each other. The result is an increase in coercivity, and a much slower relaxation rate in thermoremanent magnetization experiments.
Acknowledgements The authors would like to acknowledge Professor K.M. Unruh for the use of his XRD apparatus and help with the diffraction analysis. The authors
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would also like to acknowledge the NSF INT0138151 US-Pakistan Cooperative Research Grant.
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