Triethylene glycol stabilized MnFe2O4 nanoparticle: Synthesis, magnetic and electrical characterization

Materials Research Bulletin 48 (2013) 1057–1064

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Triethylene glycol stabilized MnFe2O4 nanoparticle: Synthesis, magnetic and electrical characterization M. Gu¨nay a, H. Erdemi b, A. Baykal a,*, H. So¨zeri c, M.S. Toprak d,e a

Department of Chemistry, Fatih University, 34500, B. Cekmece, Istanbul, Turkey Department of Polymer Engineering, Yalova University, 77100, Yalova, Turkey c TUBITAK-UME, National Metrology Institute, PO Box 54, 41470, Gebze-Kocaeli, Turkey d Division of Functional Materials, KTH-Royal Institute of Technology, 16440, Kista-Stockholm, Sweden e Yildirim Beyazit University, Dept of Materials Science and Engineering, Ulus, Ankara, Turkey b

A R T I C L E I N F O

A B S T R A C T

Article history: Received 23 July 2012 Received in revised form 30 October 2012 Accepted 25 November 2012 Available online 12 December 2012

In this study, triethylene glycol (TEG) stabilized superparamagnetic MnFe2O4 nanoparticle (TEG@MnFe2O4) was synthesized via a simple polyol route. Composition of nanoparticle was confirmed as MnFe2O4 and surface conjugation of TEG was confirmed by Fourier transform infrared spectroscopy to be via oxygens on TEG covalently attached to metal centers on the NPs. Vibrating sample magnetometer measurements confirmed the superparamagnetic character of these nanoparticles without reaching to a saturation, and has no coercivity. The TEG@MnFe2O4 nanocomposite showed thermally activated conductivity, significant interfacial polarization effect also in the dielectric permittivity evaluations. Temperature dependency of conductivity is a strong evidence for thermally activated polarization mechanism. dc conductivity is classified into two regions over a limited temperature range and shows maximum conductivity of about 4  107 S cm1 at 120 8C. Temperature and frequency dependence of dielectric permittivity indicated interfacial polarization and temperature-assisted-reorganization effects. Crown Copyright ß 2012 Published by Elsevier Ltd. All rights reserved.

Keywords: A. Magnetic materials A. Nanostructures B. Chemical synthesis C. X-ray diffraction D. Magnetic properties D. Dielectric properties

1. Introduction Manganese ferrites belong to a group of soft ferrite materials characterized by high magnetic permeability and low losses. Manganese ferrite (MnFe2O4) has received great attention in the area of magnetic storage devices, microwave, and electronic devices because of its high magnetic permeability and high electrical resistance [1]. As a potential candidate of contrast agents in MR imaging, the super-paramagnetic manganese ferrite (MnFe2O4) NPs have been found to have a very high magnetization and large relaxivity owing to their large magnetic spin magnitude [2,3]. In general, polymer coating can reduce the aggregation and improve the colloidal stability of the magnetic NPs for biomedical applications [4]. Wan et al. prepared the monodisperse and watersoluble magnetite nanoparticles by polyol process for highperformance magnetic resonance imaging [5]. Lu et al. [6] synthesized superparamagnetic MnFe2O4 nanoparticles via thermal decomposition route for magnetic resonance imaging at tissue, cellular or even molecular levels.

* Corresponding author. Tel.: +90 212 866 33 00/2061; fax: +90 212 866 34 02. E-mail address: [email protected] (A. Baykal).

In this present study, with the help of polyol route, we developed a facile polyol route to prepare superparamagnetic TEG@MnFe2O4 NPs directly via a one-pot approach [7]. We used water-soluble triethylene glycol (TEG) molecules. Here, TEG was used both as solvent and as surfactant to control the particle growth and to prevent the aggregation of particles in high temperature reaction medium. 2. Experimental 2.1. Instrumentations X-ray powder diffraction (XRD) analysis was conducted on a Rigaku Smart Lab Diffractometer operated at 40 kV and 35 mA using Cu Ka radiation. Fourier transform infrared (FT-IR) spectra were recorded in transmission mode with a Perkin Elmer BX FT-IR infrared spectrometer. The powder samples were ground with KBr and compressed into a pellet. FT-IR spectra in the range 4000– 400 cm1 were recorded in order to investigate the nature of the chemical bonds formed. The thermal stability was determined by thermogravimetric analysis (TGA), Perkin Elmer Instruments model, STA 6000. The TGA thermograms were recorded for 5 mg of powder sample at a

0025-5408/$ – see front matter . Crown Copyright ß 2012 Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.materresbull.2012.11.097

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heating rate of 10 8C/min in the temperature range of 30–600 8C under nitrogen atmosphere. The real (e0 ) and imaginary (e00 ) parts of complex dielectric permittivity e*[=e0 (v) + ie00 (v)] were measured with a Novocontrol dielectric-impedance analyzer. The dielectric data (e0 ,e00 ) were collected as a function of temperature and frequency. These films were sandwiched between gold blocking electrodes and the conductivities were measured in the frequency range of 0.1 Hz to 3 MHz at 10 8C intervals. The temperature change was controlled with a Novocontrol cryosystem with a precision of 0.01 8C. High resolution transmission electron microscopy (HR-TEM) analysis was performed using a JEOL JEM 2100 microscope. A drop of diluted sample in alcohol was dripped on a TEM grid. VSM measurements were performed by using a Quantum Design Vibrating sample magnetometer (QD-VSM). 2.2. Chemicals Triethylene glycol (TEG), iron acetate (Fe(acac)3) and manganese acetate (Mn(ac)24H2O) was purchased from Merck and used as received, without any further purification. 2.3. Procedure 0.25 g Mn(Ac)24H2O and 0.71 g Fe(acac)3 were dissolved in TEG (20 mL) was in three-neck round bottomed flask equipped with condenser, magnetic stirrer, heating mantle and stirred under argon and refluxed at 210 8C for 2 h. Then the reflux process was continued at 280 8C for additional 1 h. Finally the dark-brown mixture was cooled down to room temperature and then, washed with ethanol/H2O. Dark-brown solid was separated from the solution by magnet. The obtained black precipitate was washed again with ethanol for three times, which could be easily dispersed in water. Then solid ppt was dried at 100 8C in domestic oven. 3. Results and discussion 3.1. XRD analysis

3.2. FT-IR analysis To characterize the surface nature of the as-prepared particles, FT-IR spectra of TEG and TEG@MnFe2O4 NPs were investigated in detail (Fig. 2). In both spectra, a broad band at around 3400 cm1 was corresponding to O–H stretching vibration of TEG and absorbed water molecules. The peaks at 2900 cm1 and 2800 cm1 are assigned to the asymmetric (nas) and symmetric (ns) stretching vibrations of methylene (– CH2) of TEG molecules, respectively. The triplet band of the C–O– C stretching vibrations with maxima at about 1147 cm1, 1116 cm1, and 1062 cm1 in curve b are typical features for the TEG molecules [9–14]. However, in the IR spectrum of the asprepared NPs (curve a), there is only one peak at 1060 cm1, which is downshifted from 1067 cm1 with triplet band for the free TEG, indicating the strong interactions between NP surfaces and TEG molecules, leading to weakened C–O–C stretching vibrations of the TEG molecules. Taken together, the strong metal–O absorption band around 566 cm1, along with the IR features of TEG molecules for the NPs further supports the stabilizing role played by TEG. Conjugation scheme of TEG onto the surface of MnFe2O4 is presented in Scheme 1. 3.3. TG analysis The presence of TEG on the surface of the MnFe2O4 NPs is supported by TGA measurement. Fig. 3 shows the TGA curve of TEG and TEG@MnFe2O4 NPs. TEG starts to decompose at 150 8C and the decomposition reaches completion around 240 8C in one step. TEG@MnFe2O4 NPs show a three-stage for weight loss in the temperature ranges of 50–250 8C, 260–430 8C and 440–700 8C with a total weight loss of 25%. The first slight amount of weight (5%) loss corresponds to the evaporation of adsorbed water and ethanol. The second (15%) and third steps could be due to the decomposition of TEG residues on the nanoparticles at different configurations/conjugations attached onto MnFe2O4 NPs. It is important to note that the decomposition of TEG started very late,

311

The XRD pattern of TEG@MnFe2O4 NPs is presented in Fig. 1. All of the observed diffraction peaks are indexed by the cubic structure of MnFe2O4 spinel (JCPDS card no. 10-0319, 73-1964) phase. The broadening of the diffraction peaks distinctly indicates

the nanocrystalline nature of the materials. To determine the crystallite size of the sample, the XRD profile was fitted according to Eq. (1) in Wejrzanowski et al. [8] and Pielaszek [9] which allows the estimation of average crystallite size and its standard deviation from XRD. The experimental line profile, shown in Fig. 1 was fitted for 5 peaks (2 2 0), (3 1 1), (4 0 0), (5 1 1) and (4 4 0). The calculated average crystallite size, DXRD, of the product was 10  2 nm.

1,0

(a)

440

0,8

% Transmittance

511

400

220

CPS (a.u.)

0,9

0,7

(b) 0,6

0,5

0,4

0,3

20

30

40

50

60

70

2 Theta (deg.)

4000

3500

3000

2500

2000

1500

1000

-1

Wavenumber, cm Fig. 1. XRD powder diffraction pattern of TEG@MnFe2O4 NPs and its line profile fitting.

Fig. 2. FT-IR spectra of (a) TEG and (b) TEG@MnFe2O4 NPs.

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Scheme 1. Schematic presentation of synthesis of TEG@MnFe2O4 NPs.

around 240 8C, which may indicate the stabilization of TEG due to its conjugation with the NP’s surface. Based on the TGA the content of organic and inorganic phases were determined as 25% and 75% respectively [12–14]. 3.4. TEM analysis Morphology of TEG@MnFe2O4 NPs, along with their size distribution histogram is evaluated with TEM analyses and results are presented in Fig. 4. Some agglomeration is observed from the micrographs, which may be due to the sample preparation. MnFe2O4 NPs are observed to have nearly spherical morphology with an estimated average size of 11  1 nm. Size estimated from TEM micrographs agrees well with the crystallite size estimated from XRD line profile fitting (8  2 nm), which may reveal nearly single crystalline character of Fe3O4 nanoparticles. 3.5. VSM analysis The magnetic properties of TEG@MnFe2O4 NPs were investigated at room temperature by vibrating sample magnetometer. M–H hysteresis curve of the sample was taken up to an external field of 15 kOe and shown in Fig. 5. The figure reveals that the specific saturation magnetization (Ms) of the composite is nearly 48 emu/g, without reaching to a saturation, and has no coercivity. The specific saturation magnetization of the sample, 48 emu/g, should be normalized to the weight of the magnetic core, which is

100

(a)

    mH kB T M ¼ M s coth  kB T mH where m denotes the mean magnetic moment of a single particle, H is applied field and kBT corresponds to the thermal energy of the particles. The Langevin relation considers each particle as a magnetic monodomain and can be used to calculate average particle size from the measured M–H hysteresis curves as described in Refs. [19,20]. Thus, we have determined the average magnetic domain size as 7.75  1.50 nm; this size is slightly smaller than the average sizes estimated from XRD and TEM which may reflect the presence of non-magnetic surface layer on the NPs. 3.6. Dielectric and electrical properties of TEG@MnFe2O4 NPs 3.6.1. ac conductivity The alternating current (ac) conductivity of the TEG@MnFe2O4 NPs was measured from 20 to 120 8C using impedance spectroscopy as a function of frequency and temperature. The frequencydependent ac conductivity (sac(v)) graph of TEG@MnFe2O4 NPs presented in Fig. 6 has been obtained using the following equation:

s 0 ðvÞ ¼ s ac ðvÞ ¼ e00 ðvÞve0

(1)

60 DTG

% Weight Loss

80

about %75 of the total. Then, Ms of the MnFe2O4 NPs becomes 64 emu/g, which is still far from the theoretically predicted value (i.e., 120.8 emu/g) and smaller than some values reported in the literature [15,16]. The low value of the saturation magnetization of these nanoparticles is mainly due to the presence of non-magnetic TEG layer (i.e. magnetically dead layer), surface spin disorder and spin canting effects [17,18]. These are typical features observed in superparamagnetic particles magnetization of which can be described by Langevin function.

40 200 400 600 Temperature ( C)

20

(b)

0 0

100

200

300

400

500

600

700

Temperature ( 0C) Fig. 3. TGA thermogram of (a) TEG@MnFe2O4 NPs and (b) TEG.

800

where s0 (v) is the real part of conductivity, v(=2pf) is the angular frequency, e00 is the imaginary part of complex dielectric permittivity (e00 ) and e0(=8.852  1014 F cm1) is the vacuum permittivity. As seen in Fig. 6, sac curve exhibits frequency dependent linear decrease at low frequency regime which is assigned to the polarization of the blocking electrodes. As temperature increases, one can see that the curves of sac versus frequency for various temperatures comprise conductivity plateau regions, while the plateau region of sac of hot nanocomposite is well developed particularly at low frequencies. The plateau region at low temperatures is slightly well-distinguished at medium frequencies, which shifts toward higher frequencies with increasing

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Fig. 4. TEM micrographs of TEG@MnFe2O4 NPs at different magnifications and particle size histogram obtained thereof.

temperature. That is to say, the conductivity plateau regions broaden toward higher frequencies at higher temperatures. It is also seen from Fig. 6 that the sac are strongly temperature dependent and shows conductivity values of 3.3  108 and 4.3  107 S cm1 at 20 and 120 8C (100 Hz), respectively, while the conductivities were found to be 8.6  106 and 2.7  105 S cm1at 1 MHz at the same temperatures. In previous studies, It has been shown that this phenomenon is a strong indication for ionic conductivity [21,22]. This significant enhancement in conductivity of TEG@MnFe2O4 with temperature may be

assigned to the increasing mobility of charge carriers. At low temperatures, MnFe2O4 nanoparticle surrounded by TEG can form a random network. When temperature is increased regularly, the nanoparticles become more organized and exhibit more capacitive behavior which results in an improvement of electrical conductivity. Additionally, MnFe2O4 nanoparticles may also interact with each other leading to a percolated path that provides electrical current to flow through both semiconducting MnFe2O4 nanoparticles and TEG with the temperature as reported in the literature [23,24]. This fact is called as temperature-assisted organized network formation of the

80

Normalized Curve Langevin Fit TEG @ MnFe2O4

60 40

M, emu/g

20 0 -20 -40 -60 -80

-15000

-10000

-5000

0

5000

10000

15000

H, Oe Fig. 5. Room temperature M–H hysteresis curve of TEG@MnFe2O4 NPs. Solid line, Langevin fit, is used to calculate average magnetic domain size.

Fig. 6. The frequency-dependent ac conductivity of TEG@MnFe2O4 NPs at various temperatures.

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TEG surrounded MnFe2O4 NP’s which results in adjustment of overall conductivities with temperatures. The phase transition temperature of TEG@MnFe2O4 is around 45 8C at which the reorganization of TEG@MnFe2O4 is completed. Consequently, the sac increases steadily with reciprocal temperature in Arrhenius plot up to 40 8C. Subsequently, a transition region takes place between 40 and 50 8C. The sac continues to increase linearly with reciprocal temperature over 50 8C which is basically due to the effect of the thermal energy applied on polymer-like material at about 45 8C. Regarding the frequency dependency of sac, it has been observed that TEG@MnFe2O4 NP’s exhibit at low temperature and low frequency regime while it shows more temperaturedependent behavior as temperature increases at low frequencies. Nevertheless, the sac of nanoparticles is independent of temperature and it is nearly the same for all temperatures at higher frequencies particularly over 10 kHz which may also be an indication for ionic conductivity [25]. The temperature independent behavior of several low mobility polymers and even for crystalline materials, non-crystalline and liquid semiconductors can be described by sharply defined power law:

s ac ðvÞ ¼ Avn

(2)

where v is the angular frequency, n is the frequency exponent and A is a temperature independent constant [26,27]. Frequency exponent (n) values were calculated from the slopes of log sac  log v plot and the variations of n value with temperature are given in the inset of Fig. 7. The sac of TEG@MnFe2O4 NP’s represents a regular increment as a function of frequency in Arrhenius plots as shown in Fig. 7 while it remains nearly constant at high temperatures and high frequencies. The increment becomes more apparent over 45 8C. As mentioned before, complete reorganization of TEG@MnFe2O4 NP’s occurs over a temperature of 45 8C, and then the sac increases again linearly with reciprocal temperature in Arrhenius plot after transition region takes place at a temperature range of 40–50 8C. This is essentially due to the influence of the thermal energy (kBT) exerted on TEG@MnFe2O4 NP’s. On the other hand, n values are strongly temperature-dependent and were found to be in the range of 0.76– 1.16 which is consistent with previous studies [28,29]. The variation of n values with temperature suggests thermally activated polarization mechanism as a result the lower n values from the ac measurements, the stronger electrode polarization if

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Fig. 8. Temperature-dependent dc conductivity of TEG@MnFe2O4 NPs.

we assume that the conduction mechanism is based on ion migration in the applied electric field. Consequently, it indicates that a kind of temperature dependent reorganization of TEG@MnFe2O4 NP’s occurs within the range of phase transition in nanocomposite matrices. 3.6.2. dc conductivity The direct current (dc) conductivity of TEG@MnFe2O4 NP’s versus reciprocal temperature is shown in Fig. 8. The frequency independent conductivity is described with the dc conductivity (sdc) which was obtained from plateau region in graphs of sac vs freq. by linear fittings. In non-plateau regions, the middle region was fitted linearly to reduce the effect of electrode polarization and dispersion. It is clearly seen that the sdc of the sample strongly temperature-dependent and it can be classified into three regions the temperature range of 20–120 8C interval including transition region between 40 and 50 8C. The sdc exhibits Arrhenius behavior before and after transition zone which takes place in the range of 20–40 8C and 50–120 8C, respectively. It has been already indicated that the conductivity of TEG@MnFe2O4 NP’s is thermally activated and it can be interpreted according to the following equation: log s dc ¼ log s 0 

Ea kB T

(3)

where sdc is the dc conductivity, s0 is the pre-exponential term, Ea is the activation energy, kB is the Boltzmann constant (8.617  105 eV K1) and T is the temperature in K. The activation energy, Ea, values were calculated from the slopes of Arrhenius plots before (20–40 8C) and after (50–120 8C) transition region, and Ea values were found to be Ea1 = 0.102 eV, Ea2 = 0.298 eV, respectively. The activation energies and related conductivities are comparable with previously studied similar work [11].

Fig. 7. Log sac  log v plot of TEG@MnFe2O4 NPs. In the inset is given Frequency exponent (n) as a function of temperature.

3.6.3. Frequency and temperature dependence of dielectric permittivity parameters (e0 , e00 ) The complex permittivity parameters of real (e0 ) and imaginary (e00 ) parts of TEG@MnFe2O4 were studied and the real part of permittivity as a function of frequency at various temperatures is shown in Figs. 9 and 10. The e0 of the TEG@MnFe2O4 NP’s reduces significantly with frequency up to 10 kHz at low temperatures while the decline in real part is relatively slow at high temperatures. This result is certainly due to the reorganization of nanocomposite at high temperature range as a whole and phase

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Fig. 9. The real part of permittivity (e0 ) as a function of frequency at various temperatures for TEG@MnFe2O4 NPs.

transition (45 8C) when passed through the medium temperature range [30]. The e0 exhibits a slight curvature over transition temperature around 50 Hz which shifted slightly to higher frequencies with temperature. The variation of the e0 can be found to be less important to some extent at high temperature in applying external electric field with high frequency. For polymers or polymer-like materials, the e0 decreases regularly with increasing frequency which has also been observed for TEG@MnFe2O4 NP’s. The change of e0 at higher frequencies is rather small and remains almost constant which can be described by frequency dependence of the polarization mechanisms which indicates that the variation of dielectric constant with frequency showed the existence of material electrode interface polarization processes taking place at low frequencies [31]. The dielectric constant is associated to how fast the polarizable units in polymers or polymer-like materials orient to keep up with the oscillations of an applied alternating electric field. The orientational polarization declines with increasing frequency since the orientation of dipole moments needs a longer time than electronic and ionic polariza-

Fig. 10. The imaginary part of permittivity (e00 ) as a function of frequency at various temperatures for TEG@MnFe2O4 NPs.

tions that result in a considerable reduction in the real part of dielectric permittivity. Regarding the temperature dependency of the e0 , in general, it increases with temperature due to the molecular orientation and arrangement [30]. It is seen Fig. 9 that the e0 increases significantly with temperature because of the enhancement of interfaces between TEG and MnFe2O4 NP’s. Accordingly, the e0 of TEG@MnFe2O4 NP’s increases remarkably with temperature at low frequency regime while the increment is rather small particularly at higher frequencies. The rubber-like behavior of TEG stabilized MnFe2O4 NP’s enables them to readily interact with each other providing a percolated path that will assist the conduction as reported in previously similar work [25]. In general, the dielectric constant increases with temperature as seen in semiconductors. The thermal energy converts the bound charges to the charge carriers which results and increase in charge carrier concentration as a result it always leads to easy alignment of dipoles in the applied ac electrical field and accordingly increases in dielectric constants. The mobility of the charge carriers enhances by elevating the temperature because of the increase in thermal energy. It is also known that the measured effective permittivity depends both on the microstructure and on the permittivity of the nanoparticles [28]. This micro structural phenomenon can be elucidated by Clausius–Mosotti equation, which explains the relation between dielectric constant and polarizability coefficient of inhomogeneous double structure including the highly conducting grains separated by moderately poor conducting grain boundaries. The conductivity difference between grains and grain boundaries exhibits dissimilar resistances which lead to the accumulation of charge carriers in separated boundaries and an enhancement in dielectric constants of nanoparticles with increasing temperature. The curves of imaginary part of the permittivity (e00 ) as a function of frequency for various temperatures are displayed in Fig. 10. The e00 of TEG@MnFe2O4 NP’s shows a linear decrease with frequency which is more substantial at higher temperatures and lower frequency regime. It has been explained that the grain boundaries are effective at lower frequencies and the grains at higher frequencies. At lower frequencies, therefore, more energy is required for the polarization in the grain boundaries, which results in high-energy loss whereas less energy is required for polarization in the grains and causing small energy loss at high frequencies [32,33]. As a result, interface polarization is leading at lower frequencies while other mechanisms such as electronic and ionic are more effective at higher frequencies [34]. Accordingly, the imaginary part of permittivity turns out to be less sensitive to both frequency and temperature and stays almost constant at high frequency regime (1 kHz to 3 MHz) particularly at lower temperatures as reported previously [35,36]. 3.6.4. Frequency and temperature dependence of modulus (M0 , M00 ) The electrical modulus formalism is a quite useful method to study the influence of polarization on conductivity of nanoparticles. The real (M0 ) and imaginary (M00 ) parts of the complex dielectric modulus versus frequency are depicted in Figs. 11 and 12, respectively. The electrical modulus has been calculated by using the equation given in the previously reported literature [20,37] and the related graphs of M0 and M00 are shown in Figs. 11 and 12. The frequency dependence of M0 can be interpreted by the universal power law with a small deviation in the low frequency region according to the equation of M 0 ðT; vÞ ¼ M 00 ðTÞvn , where n varies from 0.015 to 0.44 with variation of temperature (inset Fig. of F). It is strong evidence that the variation of n with temperature is due to the polarization effect. At low temperature

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electrical functions. In other words, it may be ascribed to influence of the temperature-assisted-reorganization due to the significant structural difference of nanoparticles. The maximum of the peaks designates the conductivity relaxation times and as a consequence related to the conductivity which is strongly dependent on temperature. Subsequently, it should be underlined that the interfacial polarization effect on this nanoparticle is significant to neglect. 4. Conclusion

Fig. 11. The real part of modulus (M0 ) as a function of frequency at various temperatures for TEG@MnFe2O4 NPs.

The triethylene coated superparamagnetic manganese ferrite (TEG@MnFe2O4 NPs) were synthesized through a simple polyolreflux method. A comparison of average particle size (8  2 nm), crystallite size (10  2 nm) and magnetic domain size (7.75  1.50 nm) indicates nearly single crystalline character of these NPs. The average magnetic domain size as 7.75  1.50 nm; this size is slightly smaller than the average sizes estimated from XRD and TEM which may reflect the presence of non-magnetic surface layer on the NPs which was also proved by thermal analysis (TG). The ac conductivity shows both temperature and frequency dependent behavior at low temperature regime (<50 8C) while it is rather temperature dependent particularly at low frequencies as temperature increases. Temperature-dependent dc conductivity provides two activation energies, 0.102 and 0.298 eV, before and after transition regimes, respectively. The variation of n values with temperature is a strong indication for thermally activated polarization mechanism. Analysis of dielectric permittivity functions suggest that conduction mechanism depends both on temperature and the nature of the reorganization of the nanocomposite. Dielectric constants increased with temperature whereas decreased significantly with frequency as seen in semiconductors. The real and imaginary parts of electrical modulus shows a maximum which shifts to higher frequencies. These materials may be promising candidates for applications and the devised fabrication route is rather versatile that can be applied for the fabrication of other ferrite or oxide nanoparticle–nanocomposite systems. Acknowledgments This work is supported by Fatih University under BAP grant no. P50021104-B.

Fig. 12. The imaginary part of modulus (M0 ) as a function of frequency at various temperatures for TEG@MnFe2O4 NPs.

regime (20–40 8C), the frequency dependency of M0 is insignificant while it shows more temperature-dependent behavior. However, frequency dependency is considerable at higher temperatures (50– 120 8C) particularly at lower frequencies and this frequencydependent tendency continues up to 5 Hz for a temperature of 50 8C, whereas it exhibits this dependency up to 100 Hz for 120 8C. At higher frequencies, the change of M0 with frequency and temperature is remarkably low, and eventually reached a constant value being independent of frequency and temperature over 1 kHz as it has been in previously reported studies [25,38]. The imaginary part (M00 ) of electrical modulus is a suitable method in order to suppress the effect of electrode polarization at lower frequencies. The M00 increases gradually with temperature, and then it shows a maximum at low frequencies which is well developed at higher temperatures (>70 8C) and finally it shifts to higher frequencies (from 1 Hz to 25 Hz). This non-Debye behavior is a consequence of distributions of relaxation time and because of non-exponential approach of

References [1] Y.W. Ju, J.H. Park, H.R. Jung, S.J. Cho, W.J. Lee, Compos. Sci. Technol. 68 (2008) 1704–1709. [2] J.H. Lee, Y-M. Huh, Y-W. Jun, J-W. Seo, J-T. Jang, H-T. Song, et al. Nat. Med. 13 (2007) 95–99. [3] U.I. Tromsdorf, N.C. Bigall, M.G. Kaul, O.T. Bruns, M.S. Nikolic, B. Mollwitz, et al. Nano Lett. 7 (2007) 2422–2427. [4] M.R. Phadatare, V.M. Khot, A.B. Salunkhe, N.D. Thorat, S.H. Pawar, J. Magn. Magn. Mater. 324 (2012) 770–772. [5] J. Wan, W. Cai, X. Meng, E. Liu, Chem. Commun. 305 (2007) 5004–5006. [6] J. Lu, S. Ma, J. Sun, C. Xia, C. Liu, Z. Wang, X. Zhao, F. Gao, Q. Gong, B. Song, X. Shuai, H. Ai, Z. Gu, Biomaterials 30 (2009) 2919–2928. [7] C. Feldmann, H.O. Jungk, Angew. Chem. Int. Ed. 40 (2001) 359–362. [8] T. Wejrzanowski, R. Pielaszek, A. Opalin´ska, H. Matysiak, W. Lojkowski, K.J. Kurzydlowski, Appl. Surf. Sci. 253 (2006) 204–209. [9] R. Pielaszek, Appl. Crystallography Proceedings of the XIX Conference, Krakow, Poland, (2003), p. 43. [10] W. Cai, J. Wan, J. Colloid Interface Sci. 305 (2007) 366–370. [11] H. Deligoz, A. Baykal, E.E. Tanrıverdi, Z. Durmus, M.S. Toprak, Mater. Res. Bull. 47 (2012) 537–543. [12] H.M. Xiong, X. Zhao, J.S. Chen, J. Phys. Chem. B 105 (2001) 10169–10174; A. Baykal, H. Deligo¨z, H. Sozeri, Z. Durmus, M.S. Toprak, J. Superconduct. Novel Magnet. (2012), http://dx.doi.org/10.1007/s10948-012-1500-x. [13] M. Gu¨nay, A. Baykal, H. So¨zeri, J. Superconduct. Novel Magnet. (2012), http:// dx.doi.org/10.1007/s10948-012-1627-9. [14] H. Sozeri, Z. Durmus, A. Baykal, Mater. Res. Bull. (2012), http://dx.doi.org/ 10.1016/j.materresbull.2012.05.036.

1064

M. Gu¨nay et al. / Materials Research Bulletin 48 (2013) 1057–1064

[15] J. Xie, S. Peng, N. Brower, N. Pourmand, S.X. Wang, S. Sun, Pure Appl. Chem. 78 (2006) 1003–1014. [16] l. Josephson, J. Lewis, P. Jacobs, P.F. Hahn, D.D. Stark, Magn. Reson. Imaging 6 (1988) 647–653. [17] R.H. Kodama, A.E. Berkowitz, E.J. McNiff Jr., S. Foner, Phys. Rev. Lett. 77 (1996) 394–402. [18] X. Batlle, A. Labarta, J. Phys. D: Appl. Phys. 35 (2002) R15–R42. [19] H. Deligo¨z, A. Baykal, M. Senel, H. So¨zeri, E. Karaog˘lu, M.S. Toprak, Synth. Met. 162 (2012) 590–597. [20] Z. Durmus, H. Erdemi, A. Aslan, M.S. Toprak, H. Sozeri, A. Baykal, Polyhedron 30 (2011) 419–426. ¨ zden, M.S. Toprak, J. Alloys ¨ nal, H. Kavas, Z. Durmus, S¸. O [21] A. Baykal, N. Bıtrak, B. U Compd. 502 (2010) 199–206. [22] B. Unal, Z. Durmus, H. Kavas, A. Baykal, M.S. Toprak, Mater. Chem. Phys. 123 (2010) 184–190. [23] M.Z. Kassaee, H. Masrouri, F. Movahedi, Appl. Catal. A: Gen. 395 (28–33) (2011) 395–402. [24] H. Yang, C. Zhang, X. Shi, H. Hu, X. Du, Y. Fang, Y. Ma, H. Wu, S. Yang, Biomaterials 31 (2010) 3667–3673. ¨ nal, Z. Durmus, A. Baykal, H. So¨zeri, M.S. Toprak, L. Alpsoy, J. Alloys Compd. [25] B. U 505 (2010) 172–178.

[26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38]

R. Bouamrane, D.P. Almond, J. Phys. Condens. Matter 15 (2003) 4089–4094. W. Qu, T.M. Ko, R.H. Vora, T.S. Chung, Polymer 42 (2001) 6393–6401. B. Hallouet, B. Wetzel, R. Pelster, J. Nanomater. 2007 (2007) 11–17. I.M. Afandiyeva, I. Do¨kme, S. Altındal, A. Tataroglu, Microelectron. Eng. 85 (2008) 247–252. S. Muruganand, S.K. Narayandass, D. Mangalaraj, T.M. Vijayan, Polym. Int. 50 (2001) 1089–1094. S. Ukishima, M. Iijima, M. Sato, Y. Takahashi, E. Fukada, Thin Solid Films 308-309 (1997) 475–482. M.J. Iqbal, B. Ismail, J. Alloys Compd. 504 (2010) 440–447. B. Unal, M.S. Toprak, Z. Durmus, H. So¨zeri, A. Baykal, J. Nanopart. Res. 12 (2010) 3039–3048. G.M. Tsangaris, O.C. Psarras, N. Kouloumbi, J. Mater. Sci. 33 (1998) 2027–2032. E. Karaog˘lu, A. Baykal, H. Erdemi, L. Alpsoy, H. Sozeri, J. Alloys Compd. 509 (2011) 9218–9225. J. Mu¨rbe, A. Rechtenbach, T. To¨pfer, Mater. Chem. Phys. 110 (2008) 426–431. Z. Durmus, H. Kavas, A. Baykal, H. Sozeri, L. Alpsoy, S.U¨.C. elik, M.S. Toprak, J. Alloys Compd. 509 (2011) 2555–2561. B. Unal, Z. Durmus, A. Baykal, M.S. Toprak, H. Sozeri, A. Bozkurt, J. Alloys Compd. 509 (2011) 8199–8206.