Physica B. Doping effect of Y 3þ ions on the microstructural and electromagnetic properties of Mn Zn ferrites

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1 Physica B 407 (2012) Contents lists available at SciVerse ScienceDirect Physica B journal homepage: Doping effect of Y 3þ ions on the microstructural and electromagnetic properties of Mn Zn ferrites Qingkai Xing a, Zhijian Peng a,n, Chengbiao Wang a, Zhiqiang Fu a, Xiuli Fu b,n,n a School of Engineering and Technology, China University of Geosciences, Beijing , PR China b School of Science, Beijing University of Posts and Telecommunications, Beijing , PR China article info Article history: Received 23 August 2011 Received in revised form 20 October 2011 Accepted 3 November 2011 Available online 6 November 2011 Keywords: Mn Zn ferrite Y 3þ ion Doping Electromagnetic properties abstract Mn Zn ferrites doped with different contents of Y 3 þ ions were prepared by conventional two-step synthesis method. The microstructure and electromagnetic properties of the as-prepared Mn Zn ferrites were investigated. It was found that all the samples consisted of ferrite phases of typical spinel cubic structure, and when Y 3 þ ion content was upto, yttriumirongarnet (Y 3 Fe 5 O 12 ) phase with garnet structure was detected. With increasing doping content of Y 3 þ ions, the lattice constant and grain size increased, and after an increase to its maximum value, the sample apparent and relative densities dropped down. Through the analysis of magnetic properties, it was revealed that the saturation magnetization, and both the real and imaginary parts of permeability of the as-prepared samples raised with increasing doping content of Y 3 þ ions but decreased with more Y 3 þ ions, while their coercivity showed an opposite change trend; and the Curie temperature increased monotonously. The measurement of dielectric properties indicated that the dielectric constant of the doped Mn Zn ferrites presented a rise with increasing Y 3 þ ion content, and dropped down gradually when more Y 3 þ ions were doped, while the dielectric loss tangent would decrease with Y 3 þ content upto, but after that, it increased. & 2011 Elsevier B.V. All rights reserved. 1. Introduction Manganese zinc (Mn Zn) ferrites, which have been studied for several decades because of their high initial permeability, low loss, high saturation magnetization and relatively high Curie temperature, have widely served as fundamental materials in electronic and information industries, being used as recording heads, choke coils and communication pulse transformers and so on [1,2]. It is well known that Mn ferrites are of inverse spinel structure because 80% of Mn ions occupy tetrahedral site (A-site), which is surrounded by four O 2 ions, and the left 20% of Mn ions occupy octahedral site (B-site), which is surrounded by six O 2 ions. The cation distribution for Mn ferrites can be represented by [Mn þ Fe þ ] A [Mn þ Fe þ Fe þ ] B O 4 [3,4]. In Mn Zn ferrites, Zn 2 þ ions have strong preference for A-sites [5], leading a decrease of Fe 3 þ in A-sites. This kind of spinel structure allows the introduction of different metallic ions into the lattice, thus modifying the magnetic and electric properties considerably [6]. According to Ref. [7], additive materials can affect the microstructure and properties of Mn Zn ferrites by three main mechanisms. Some n Corresponding author. Tel.: þ ; fax: þ nn Corresponding author. addresses: pengzhijian@cugb.edu.cn (Z. Peng), xiulifu@bupt.edu.cn (X. Fu). additives such as Co 3 þ,gd 3 þ and Eu 3 þ [8,9] can enter into the spinel lattice, replacing the metallic cations in regular A- or B-sites, thus altering the microstructure and properties of the ferrites. Some can affect the grain growth by different mechanisms. For example, CuO and Bi 2 O 3 can act as grain growth accelerator because these additives can create a thin layer of liquid phase during sintering [10]; tungsten and vanadium ions with high valence would increase cation vacancy in lattice, promoting the pore mobility and grain growth [11,12]; and when introducing TiO 2, the grain growth of Mn Zn ferrites can be restrained and the microstructure tends to be more uniform [13]. Certain additives such as CaO can reduce the eddy-current loss of ferrites and finally reduce the total power loss, because this kind of dopants can create an electrical insulating film around the ferrite grains, increasing the electrical resistivity of the material [14]. In the work of Ishaque et al. [19], nonmagnetic Y 3 þ ions were introduced into Ni ferrites. Their results revealed that the structural and transport properties of Ni ferrites could be improved by doping of Y 3 þ ions. However, to the best of our knowledge, no information was reported in literature on the microstructure and electromagnetic properties of Y 3 þ substituted Mn Zn ferrites and the doping effect of Y 3 þ ions on Mn Zn ferrites need to be established. So in this work, aiming at obtaining Mn Zn ferrites with comparatively high performance in magnetic and electric properties, such as saturation magnetization, coercivity, dielectric constant, and so /$ - see front matter & 2011 Elsevier B.V. All rights reserved. doi: /j.physb

2 Q. Xing et al. / Physica B 407 (2012) on, Y 2 O 3 was introduced into Mn Zn ferrites, and the doping effect of Y 3 þ ions on the ferrites was investigated. 2. Experimental procedures 2.1. Sample preparation Using MnO 2, ZnO, Fe 2 O 3 and Y 2 O 3 of analytical reagent grade as raw materials, 0.5 wt% PVA as binder, and 0.5 wt% Davon C as dispersant, Y 3 þ -doped Mn Zn ferrites, in which the nominal composition of the original ferrite was Mn 0.5 Zn 0.5 Fe 2 O 4 and the content of extra added Y 3 þ ions ranged from 0 to in a step of, were fabricated by a conventional two-step synthesis method [15]. For all the prepared ferrite samples, the powder mixture was milled for 24 h in a ball mill with de-ionized water as milling medium and TZP ZrO 2 balls as grinding medium. For the ball-milling, the weight ratio of raw powder to grinding media was 1:2. Then, the milled slurries were dried at 120 1C in an oven. After that, the powder mixture was pre-sintered at 800 1C in air for 2 h. Then, the pre-sintered powder mixture was divided into 5 groups evenly, re-milled for 24 h after PVA, Davon C and different amounts of Y 2 O 3 were added, re-dried at 120 1C in oven, and grinded and sieved into fine powders. Then toroid samples (20 mm in out diameter, 10 mm inner diameter and 3 mm thickness), small pellets (6 mm diameter and 3 mm thickness) and bigger pellets (18 mm diameter and 3 mm thickness) were prepared with the resultant fine powders, respectively. Finally, all the green samples were sintered at optimum temperature C for 4 h under controlled N 2 /O 2 atmosphere and cooled under equilibrium conditions [16] Materials characterization Using the small pellets, the phase compositions of the samples were identified by X-ray diffraction (XRD, D/max-RB, CuKa radiation, and l¼ Å) with a continuous scanning mode at a speed of 61/min, and the lattice constants were calculated from the XRD results by software MDI Jade. The microstructure investigations were carried out on the fracture surfaces of the samples by thermal field emission scanning electron microscope (SEM, LEO-1530) equipped with an energy dispersion X-ray spectroscopy (EDX), and the grain sizes of the samples were calculated from the micrographs by software Lince PC. In order to obtain the average apparent density, the apparent densities of all the sintered samples were measured by Archimedes method according to international standard (ISO18754). The theoretical densities of the samples were calculated on the hypothesis that all the doped Y 3þ ions were reacted into yttriumirongarnet phase on the basis of the obtained XRD results and the grains of all phases in the samples were packed in the most closed way, and the relative density is the percentage of apparent density to theoretical density. Also using the small pellets, the hysteresis loops and M T curves of the samples were measured by employing a vibrating sample magnetometer (VSM, LakeShore 7307) with a maximum magnetic field of 10 KOe, and the saturation magnetization (M s ), coercivity (H c ) and Curie temperature (T c )were calculated from the measured hysteresis loops and M T curves, respectively. With the toroid samples, the characteristics of complex permeability (m¼m 0 jm 00 ) of the obtained ferrite samples were recorded by an impedance analyzer (Agilent E4991A) within the frequency from 1 to 100 MHz. Using the bigger pellets, the measurement of complex permittivity (e¼e 0 je 00 )ofthesampleswas performed by an impedance analyzer (Agilent E4991A) within the frequency from 1 to 40 MHz. All these measurements except T c were carried out at room temperature. 3. Results and discussion 3.1. Structural properties The XRD patterns of the as-prepared Mn Zn ferrites doped with different contents of Y 3þ ions are shown in Fig. 1. It can be seen from this figure that the predominant phase in all the specimens is Mn Zn ferrite phase of spinel cubic structure. However, when the doping content of Y 3þ ions is high (more than in this case), except for the typical diffraction peaks of spinel phase of Mn Zn ferrites, other small peaks were also clearly detected, indicating the presence of secondary phase, which were identified by PDF card as yttriumirongarnet (Y 3 Fe 5 O 12 ) phase of garnet structure. In the Mn Zn ferrite samples with small contents of Y 3þ ions, the yttriumirongarnet phase might not be found in the XRD detection limit due to its low content. The lattice constants of ferrite phases in the as-prepared samples calculated from the XRD data are listed in Table 1. The results reveal that the lattice constant of the samples increases with increasing doping content of Y 3þ ions, implying that the doped Y 3þ ions can also enter into the lattice of the as-prepared Mn Zn ferrites. So it can be inferred that after the addition of Y 3þ ions, the main phase of the samples was Y-doped Mn Zn ferrite. The increase trend in lattice constants of the as-prepared ferrite samples can be explained on the basis of the ionic radii. The radius of Y 3þ ion (0.95 Å) is larger than that of Fe 3þ ion (0.67 Å). Y 3þ ions with larger radius can behave as other rare-earth ions like Tb, La, Ce and Th, entering into the crystal lattice of the as-prepared Mn Zn ferrites and locating in the B-sublattice with adequate space [19,21]. The replacement of Fe 3þ ions in B-sites by Y 3þ ions will cause the expansion of unit cell, resulting in larger lattice constants. The SEM micrographs taken from the fracture surfaces of the as-prepared samples doped with different contents of Y 3 þ ions are illustrated in Fig. 2. The calculated grain sizes of the corresponding Mn Zn ferrites from these graphs are presented in Table 1. It can be seen that the grain size of the samples increases with increasing doping contents of Y 3 þ ions. The melting point of Y 2 O 3 is C, which is much higher than the sintering temperature of the ferrites in this study. So the main cause accounting for the increase of bulk grain size is not liquid-phase sintering. It can be explained as follows. When the divalent ions were replaced by Y 3 þ ions in the grain boundary region, the metallic ion vacancies in vicinity increased so as to balance the electric charges. As a result, the speed of the grain boundary movement Relative Intensity (a.u.) 20 2θ (degree) Mn-Zn ferrites YIG Fig. 1. XRD patterns of Mn Zn ferrites doped with different contents of Y 3 þ ions.

3 390 Q. Xing et al. / Physica B 407 (2012) Table 1 Basic structural parameters of Mn Zn ferrites doped with different contents of Y 3 þ ions. Content of Y 3 þ (mol%) Lattice (Å) Grain size (mm) Average apparent density (g/cm 3 ) Relative density (%) Fig. 2. SEM images of Mn Zn ferrites doped with different contents of Y 3 þ ions: (a), (b), (c), (d) and (e). increases, thereby promoting the grain growth [6,11]. At the same time, after the introduction of Y 3 þ ions, yttriumirongarnet (YIG) with garnet structure appears and accumulates in the grain boundary region as solid solution, leading to the increase of oxygen vacancies. These oxygen vacancies can serve as material transportation path, promoting the grain growth of ferrites further [17]. In addition, the measured SEM EDX results revealed that the doped Y 3 þ ions distributed in both the ferrite grains and grain boundaries, which can indirectly support the conclusion from XRD results that the ferrite phase of the samples was Y-doped. The apparent densities and relative densities of the as-prepared samples are listed in Table 1. It can be seen from this table that the relative density of the as-prepared samples keeps at a relatively high level, indicating that all the ferrite samples in this study are dense. It can also be seen from the table that the apparent density and relative density of the as-prepared samples increase with increasing doping content of Y 3 þ ions, reaching a maximum value when the concentration of Y 3 þ ions is, but then decrease with further increased doping content of Y 3 þ ion. This result reveals that an appropriate doping content of Y 3 þ ions can improve the densification of the samples. As is mentioned before, the electric charges and oxygen vacancies can promote the grain growth of ferrites after the introduction of Y 3 þ ions, and this is conducive for the expelling of closed pores in the as-prepared samples. So the density and relative density of the as-prepared samples increase when the doped Y 3 þ ion content is less than. However, owing to the increasing

4 Q. Xing et al. / Physica B 407 (2012) Table 2 Basic magnetic parameters of Mn Zn ferrites doped with different contents of Y 3þ ions mol % Content of Y 3þ (mol%) T c M s (emu/g) H c (Oe) cation and oxygen vacancies when more Y 3 þ ions were added, the density and relative density of the samples decreased, although the dopant might still promote the grain growth of ferrites. Permeability μ'' μ' 3.2. Magnetic performance 0 Table 2 lists the basic magnetic parameters of the as-prepared Mn Zn ferrites doped with different contents of Y 3þ ions. The results reveal that the Curie temperature T c of the samples increases with increasing doping content of Y 3þ ions. The saturation magnetization M s of the samples increases with Y 3þ ion content upto, presenting a maximum value of emu/g, after which it decreases with further increasing Y 3þ ion content, and the coercivity H c shows an opposite trend to M s. It is known [4,18] that the magnetic behavior of Mn Zn ferrite is governed by the iron iron interaction, which can be easily modified if the native cations are replaced by foreigners. As mentioned above, Y 3þ ions with larger radius can enter into the crystal lattice of Mn Zn ferrites, and behave as other rare-earth ions like Tb, La, Ce and Th locating in the B-sublattice with adequate space do [19], causing the alteration of magnetic properties. Curie temperature depends on the A B exchange effect (the strongest one in all the three kinds of exchange effects, A A, B B and A B), in which the interaction between Fe ions plays a leading role [20]. As mentioned in Ref [4], B-sites contain more Fe ions than A-sites. When Y 3 þ ions enter into the B-sites, some Fe 3 þ are replaced and pushed into A-sites, enhancing the quantity of iron ions in A-sites. So the A B exchange effect is enhanced, and the measured T c increases. According to the molecular magnetization, the net magnetization M is the result of the difference in the sublattice moments of the two sites (M¼9M b M a 9, where M b is the magnetic moment in B-sites and M a is the magnetic moment in A-sites), which depends upon the cation occupancy [3,21,22]. Due to the substitution of Fe 3 þ ions by nonmagnetic Y 3 þ ions, the magnetization of B-sites decreases, resulting theoretically in the decrease in saturation magnetization for the obtained Y 3 þ -substituted ferrite samples. However, as shown in Table 2, M s presents an opposite trend at first, and then drops down. This can be attributed to the enlarged grain size. A larger grain size makes the reversible displacement or domain wall movement in the direction of the applied magnetic field much easier. So the resultant M s increases at first, and then drops down, due to the enlarged grain size but the decreased magnetization in B-sites. Coercivity is the amount of reverse magnetic field, which must be applied to a magnetic material to drive the magnetic flux return to zero. For soft ferrites, it is caused by the resistance of domain wall displacement. The coercivity of ferrites depends inversely on grain sizes. Larger grains tend to consist of a greater number of domain walls. The magnetization or demagnetization caused by domain wall movement requires less energy than that required by domain rotation [21]. In contrast with the contribution of domain rotation, the contribution of domain wall movement, which needs less energy when the materials are magnetized or demagnetized, increases as the number of walls increases with increasing grain sizes. Therefore, the samples 10k having relatively larger grains are expected to have lower coercivity. However, impurities distribute in the grain boundary area, breaking and acting against the displacement of domain walls. So due to the foreign phase of yttriumirongarnet and possibly Y 2 O 3, the samples doped with more Y 2 O 3 are expected to have larger coercivity as is shown in Table 2. The initial complex permeability dispersion spectra (m¼m 0 jm 00 ) of the as-prepared samples doped with different contents of Y 3þ ions are presented in Fig. 3. It is noticed that both m 0 and m 00 increase with increasing doping content of Y 3þ ions, reaching a maximum for the samples with Y 3þ ion content being, and decrease with further increasing doping content of Y 3þ ions. The complex permeability is correlated to two different magnetizing mechanisms: the spin rotational magnetizing inside the domains and the domain wall motion [23,24]. With increasing doping content of Y 3 þ ions, the enlarged crystalline grains make the initial permeability increase. At the same time, with further increasing doping content of Y 3 þ ions, more Y 3 Fe 5 O 12 generates and segregates at grain boundaries pining at the domain walls, and the domain energy is enhanced. Thus the decrease of permeability can also be expected Dielectric performance 100k 1M 10M 100M Fig. 3. Frequency dependence of permeability dispersion spectra of Mn Zn ferrites doped with different contents of Y 3 þ ions. The variations of dielectric constant (e 0 ) and dielectric loss tangent (tan d) with frequency for the as-prepared ferrites doped with different contents of Y 3þ ions are shown in Figs. 4 and 5, respectively. Similar dielectric properties can also be found in Refs [3], [8] and [25]. Itcanbeseenfromthesetwo figures that compared with the nondoped Mn Zn ferrite, the e 0 of the Y 3þ -doped Mn Zn ferrite presents a rise with increasing Y 3þ ion content, and drops down gradually when more Y 2 O 3 is added. In low and intermediate frequency, the tan d of the ferrites decreases with Y 3þ ion content upto and increases when Y 3þ ion content is, and in high frequency, no obvious distinction of tan d can be found. Scientists have established a strong correlation between the conduction mechanism and dielectric constant of ferrites. In their studies, they explained the dielectric behavior of ferrites on the basis of the number of available Fe 2þ ions in the B-sites. The electronic exchange such as Fe 2þ 2Fe 3þ and Mn 2þ 2Mn 3þ results in a local displacement of electrons, which determines the polarization, and thus the dielectric constant of ferrites [26]. After the introductionof Y 3þ ions, an insulating intergranular secondary phase (Y 3 Fe 5 O 12, which contains Fe 3þ ions only) forms, segregating at the grain

5 392 Q. Xing et al. / Physica B 407 (2012) Dielectric constant (ε') k boundaries and impeding the oxidation of Fe 2þ ions inside the grains during the slow cooling of sample temperatures [4], increasing the content of Fe 2þ and decreasing the content of Mn 3þ,thus leading to an increase of e 0 and decrease of tan d. However,the replacement of Fe 3þ ions by Y 3þ ions in the doped samples will cause a decrease of Fe 3þ in B-sites when Y 3þ ions are introduced. Owing to this replacement, the content of iron ions decrease and the electron exchange interaction Fe 3þ 2 Fe 2þ is limited. So after hoping to a maximum, e 0 drops down gradually and decreases to a minimum, and tan d starts to increase when more Y 2 O 3 is added. 4. Conclusions 1M 10M Fig. 4. Frequency dependence of dielectric constant of Mn Zn ferrites doped with different contents of Y 3þ ions. Dielectric loss tangent (tanδ) k 100k 10M Fig. 5. Frequency dependence of dielectric loss tangent of Mn Zn ferrites doped with different contents of Y 3þ ions. Mn Zn ferrites doped with different contents of Y 3þ ions were prepared by conventional two-step synthesis method. The conclusions can be summarized as follows: (1) All the samples doped with Y 3 þ ions mainly contain ferrites of a typical spinel cubic structure. With increasing doping content of Y 3 þ ions, more Y 3 þ ions enter into the B-sites of the ferrite, causing the increase of lattice constant. The doping of Y 3 þ ions can promote the grain growth of the ferrites, and an appropriate doping content of Y 3 þ ions can improve the densification of the samples. (2) The Curie temperature of Mn Zn ferrites increases with increasing doping content of Y 3 þ ions. The saturation magnetization rises after the doping of Y 3 þ ions, but drops down with more Y 3 þ ions. The coercivity of the ferrites shows an opposite trend to the saturation magnetization. Both the real and imaginary parts of permeability of the ferrites increase with increasing doping content of Y 3 þ ions. So the introduction of an appropriate doping content of Y 3 þ ions can improve the magnetic properties of Mn Zn ferrites. (3) The introduction of Y 3 þ ions is against obtaining Mn Zn ferrites of low dielectric constant, but an appropriate doping content of Y 3 þ ions can reduce the dielectric loss. Compared with the non-doped Mn Zn ferrite, the dielectric constant of the Y 3 þ -doped Mn Zn ferrites presents a rise with increasing doping content of Y 3 þ ions, and drops down gradually when more Y 2 O 3 is added. And the dielectric loss of the samples decreases with Y 3 þ ion content upto, and increases when Y 3 þ ion content is. Acknowledgments The authors would like to thank the financial support for this work from the National Natural Science Foundation of China (Grant no and ), the Transfer and Industrialization Project of Sci-Tech Achievement (Cooperation Project between University and Factory) from Beijing Municipal Commission of Education and the Cultivating Foundation for Young Scientists in China University of Geosciences at Beijing from the Fundamental Research Funds for the Central Universities (Grant no. 2011PY191). References [1] R. Arulmurugan, B. Jeyadevan, G. Vaidyanathan, S. Sendhilnathan, J. Magn. Magn. Mater. 288 (2005) 470. [2] H. Fujioka, T. Ikeda, K. Ono, S. Ito, M. Oshima, J. Cryst. Growth 241 (2002) 309. [3] M.A. Ahmed, N. Okasha, M.M. El-Sayed, Ceram. Int. 33 (2007) 49. [4] Q.M. Wei, Jian-biao Li, Yong-jun Chen, Yong-sheng Han, Mater. Charact. 47 (2001) 247. [5] S. Modak, M. Ammar, F. Mazaleyrat, S. Das, P.K. Chakrabarti, J. Alloys Compd 473 (2009) 15. [6] H. Su, H.W. Zhang, X.L. 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