474-482 - Australian Journal of Basic and Applied Sciences

Australian Journal of Basic and Applied Sciences, 8(3) March 2014, Pages: 474-482
AENSI Journals
Australian Journal of Basic and Applied Sciences
ISSN:1991-8178
Journal home page: www.ajbasweb.com
Complex permeability, Curie temperature and activation energy as a function of
microstructure evolution in a mechanically alloyed Y3Fe5O12single-sample.
1
Rodziah Nazlan, 1Ismayadi Ismail, 1Mansor Hashim, 1Samikannu Kanagesan, 2Norlaily MohdSaidin
1
Materials Synthesis and Characterisation Laboratory, Institute of Advanced Technology (ITMA), Universiti Putra Malaysia, 43400
Serdang, Selangor.
2
Physics Department, Faculty of Science, Universiti Putra Malaysia, 43400 Serdang, Selangor.
ARTICLE INFO
Article history:
Received 12 January 2014
Received in revised form 20
March 2014
Accepted 25 March 2014
Available online 2 April 2014
Key words:
YIG, real permeability, loss factor,
microstructural evolution.
ABSTRACT
We report on an investigation to unravel morphological and magnetic-property
dependence on sintering temperature for mechanically alloyed Y3Fe5O12 single-sample.
The sample was sintered at various sintering temperatures from 600oC to 1400oC. The
dependence of the complex permeability, Curie temperature and activation energy on
the microstructural evolution was studied respectively. The complex permeability
showed an increasing trend parallel to the grain growth of the sample. How this
correspondence depends on the relationship between ordered magnetism and the
microstructure of the sample was duly explained. The curie temperature was
independent of the grain growth and remained almost constant over the grain changes.
Additionally we found that two stages of activation energy sensitively influenced the
sample grain growth, hence influencing the magnetic properties.
© 2014 AENSI Publisher All rights reserved.
To Cite This Article: Rodziah Nazlan, Ismayadi Ismail, Mansor Hashim, Samikannu Kanagesan, Norlaily MohdSaidin., Complex
permeability, Curie temperature and activation energy as a function of microstructure evolution in a mechanically alloyed Y3Fe5O12singlesample. Aust. J. Basic & Appl. Sci., 8(3): 474-482, 2014
INTRODUCTION
In recent years, great interest on yttrium-iron garnet (YIG) has been focused on the dependence of physical
and chemical properties on the grain size. For this reason, there is a renewed interest in the development of new
techniques to produce particles in different sizes, sufficiently monodispersed and with a good degree of
homogeneity. Generally, YIG powder can be prepared either by using a solid state method such as ball milling,
high energy ball milling (HEBM), mechanical alloying (MA) or by using a chemical method such
coprecipitation, sol-gel, microemulsion etc.
However, among the methods mentioned above, the MA method has attracted much interest in recent years
due to its simplicity in the preparation of various interesting solid-state materials. MA milling is used for the
preparation of nanocrystalline powders and takes advantage of perturbation of surface-bonded species by
pressure to enhance thermodynamics and kinetic reactions at room temperature or at least much lower
temperatures than normally required to produce a pure metal (Suryanarayana, C. 2001; Fathi, M.H. and Zahrani,
E.M. 2009). This is due to the energy transferred from the milling media to powder particles which are
continuously submitted to fracture and cold welding processes that will define their final morphology. The
energy transfer to the powder particles in these MA mills takes place by a shearing action or impact of the high
velocity balls with the powder. The size of the nanoparticles depends on several factors namely milling speed;
type and size distribution of balls; ball to powder weight ratio; milling atmosphere and so on (McCormick, P.G.,
et al., 1998). This process has the advantage to produce large quantities of material and is already a commercial
technology. However, contamination from balls should be taken care of.
In the ferrite industry, it is desirable to produce samples with low loss and high permeability. However,Ref.
(Roess, E. 1982) has pointed out that it is impossible to obtain both high permeability and very low loss
material. For this reason, there is a renewed interest in the development of new techniques to produce particles
in different sizes, sufficiently monodispersed and with high homogeneity to achieve the best compromised low
loss-high permeability combination.
The earliest work on correlating grain size and permeability was carried out byRef. (Guillaud, C. 1957) on
manganese zinc ferrite. Ref. (Guillaud, C. 1957) related the inflection at 5 µm to a change from a domain
rotation mechanism to a wall movement mechanism. He considered the limitation at about 20 µm to be due to
the presence of pores included in grains. Permeability appears to increase with increasing grain size by assuming
Corresponding Author: Ismayadi Ismail, Materials Synthesis and Characterisation Laboratory, Institute of Advanced
Technology (ITMA), Universiti Putra Malaysia, 43400 Serdang, Selangor.
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Ismayadi Ismail et al, 2014
Australian Journal of Basic and Applied Sciences, 8(3) March 2014, Pages: 474-482
the other factors, such as porosity, to remain constant. If high magnetic permeability is desired without regard to
high magnetic losses, the presence of grain boundaries will act as an impediment to the domain wall movement.
The fewer the number of grain boundaries present, the larger the grains and the higher the permeability. The
lack of purification during processing, presence of pores, and greater chemical inhomogeneity also prevent the
attainment of high permeability. Ref. (Tseng, T.J. and Lin, J.C. 1989) and (Ravindran, P. and Patil, K.C. 1987)
also confirm that the permeability increased when the grain size increased.
Ref. (Hahn, H.T. 1991) reported a decrease in permeability with porosity in nickel zinc ferrite.
MeanwhileGoldman (1990) claimed that if pores can be suppressed or located at the grain boundaries, the
permeability would increase with grain size.Guillaud (1957) noticed that resistivity in the bulk material was
dominated by grain boundaries via his study on grain boundaries (Guillaud, C. 1957). He also discovered the
importance of Ca2+ which segregated at the grain boundaries as a useful additive in reducing eddy current losses.
The thickness and chemical composition of the grain boundary are two important factors in determining the
magnetic properties of ferrites. Ref. (Akashi, T. 1961) showed that addition of a sensible amount of SiO2 to CaO
could increase resistivity and lower losses, due to an increase in grain boundary resistivity.
Ref. (Ishino, K. and Narumiya, Y. 1987) summarizes the requirements for low-loss ferrites at frequencies
up to 1 MHz. The researchers noticed the following combinations of chemical and microstructural aspects to be
important:
(i) Suppression of electron hopping from Fe3+ to Fe2+ inside the grains,
(ii) Insulating films surrounding the grain boundaries by the addition of Ca2+ and Si2+
(iii) Small and homogenous grain size,
(iv) Reduction of pores for increased density as it creates a demagnetizing field and increases flux density.
Pores should be concentrated at the grain boundaries.
Ferrites with nanometer-sized grains exhibit a much higher coercivity than samples having micron-sized
grain size (Pal, M. and Chakravorty, D. 1999). This is due to the fact that nanometer-sized particles are not
likely to have any grain boundary contribution; however, sintering introduces the effects of grain boundary and
some microstructural defects, thereby resulting in higher values of coercivity. The coercive force is the property
most sensitive to porosity and grain size. The increase in coercive force with porosity is linear at low porosity
and deviates at higher level (Smit, J. and Wijn, H.P.J. 1954). High porosity samples contain smaller particles,
which have larger coercive force. It is correlated with Neel’s theory, which states that the demagnetizing effect
will increase due to non-magnetic inclusions including pores (Treble, T.S. and Craik, T.J. 1969). Morphological
properties are acknowledged as the major factor which influences changes of magnetic properties in ferrites. For
several past decades, studies of the relationship between morphological properties and magnetic properties of
soft ferrites have been focusing only on the product of the final sintering temperature largely neglecting the
parallel evolutions of morphological and magnetic properties. Previously, we have reported the evolution of
microstructure-magnetic properties in NiZn ferrites (Syazwan Mustaffa, M., et al., 2013; Ismail, I., et al., 2013;
Ismail, I., Hashim, M. 2012; Idza, I.R., et al., 2012; Ismail, I., et al., 2012; Ismail, I., M. Hashim, 2011), to the
best of our knowledge no study has been found in the literature attempting to correlate the microstructure with
the magnetic properties of Y3Fe5O12 bulk single-sample produced by mechanically alloyed nanosized powders
and later subjected to sintering temperature from 500oC up to 1400oC. This study is aimed at evaluating the
microstructural evolution and its effect especially on the complex permeability properties of Y3Fe5O12 prepared
by the mechanical alloying method. We also analysed the activation energy involved during the sintering
process and observed its influence on the microstructural changes.
Methodology:
To synthesize YIG powders by using high energy ball milling or mechanochemical reaction, Y2O3 and αFe2O3 were used as precursor reagents. The high energy milling was carried out in a SPEX 8000D machine for 2
hours with the ball-to-powder ratio 10:1. The as-milled powders with particle sizes ranging from 50 to 180 nm
were compacted into a toroid sample which had an outer diameter of 15 mm and an inner diameter of 10 mm
with 2.5 grams in weight by uniaxial pressing at 3 tonne. One green toroid sample was repeatedly annealed in an
ambient temperature from 600⁰C to 1400⁰C and this heating process is denoted as single-sample sintering. The
density of the sample was measured using the Archimedes principle. The surface morphology of the samples
was then observed by using a Nova Nano 230 Field Emission Scanning Electron Microscope (FeSEM). An
Agilent HP4291B Impedance/Material Analyzer (Figure 4.6) was used to measure the complex permeability of
the samples. The measurements were carried out in the frequency range of 10 MHz to 1 GHz at room
temperature. The complex permeability is represented by the complex permeability components which are real
permeability, µ’ and magnetic loss factor, µ’’ as shown in equation (1) below:
∗
=
′
−
′′
(1)
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Curie temperature measurements were carried out in order to determine the transition temperature from
ferromagnetic to paramagnetic state. Before the measurement, atoroidal YIG samples which had been sintered at
9 different temperatures (600-1400oC) were each wound with 20 turns of 150 cm insulated copper wire with
both ends scraped with a sand paper. The samples were heated in the box furnace and connected to an Agilent
HP4294A Precision Impedance/Material Analyzer. For each sample, the value of inductance, Ls,was recorded
from the analyzer at a frequency of 10kHz from room temperature up to 310⁰C after every increment of 20⁰C.
The Curie temperature point was obtained by plotting Ls value at 10 kHz versus temperature graph. Adrastical
drop of the Ls value can be seen and the intersection point on the temperature axis indicates the Curie
temperature.
RESULTS AND DISCUSSION
3.1 Complex Permeability Measurement:
The room temperature real permeability, µ’, of the single-sample of YIG was measured from 10MHz to
1.8GHz (Fig.1). It can be observed that the µ’ value increases to a maximum value and then decreases rapidly to
a very low value. The µ’ value at 10MHz changes linearly with increasing sintering temperature. The observed
increase of µ’ with sintering temperature can be attributed to the increased grain size, which correspondingly
reduces the porosity and grain boundaries inside of the samples. As mentionedearlier, porosity and grain
boundaries, as well as impurities will impede the domain wall movement, thus influencing the permeability of
the sample (Xu, Z., et al., 2009).
Furthermore, an increase in the sintering temperature will result in a decrease of magnetic anisotropy,
reducing the hindrance to the movement of domain walls, thus increasing the value of µ’(Gupta, N., et al.,
2005). The µ’ spectra also show the ferromagnetic resonance in our testing frequency range. These, according to
Snoek’s Law, give the relation µ’fr = constant, where fr represents the ferromagnetic resonance frequency.
Snoek’s Law states that the higher the values of µ’, the lower is the value of the ferromagnetic resonance
frequency. The sloping part of µ’ spectra also indicate that ferromagnetic resonance was occurring at a
frequency inside the testing range.
Table 1 shows the resultsfor the initial permeability, µ i,value for the single-sample sintering. The initial
permeability is qualitatively the magnetization response at very low exciting fields. The µ i value was found to
vary with sintering temperature, where the sample sintered at 14000C has the highest real permeability value of
36.0. This agrees with the expectation that permeability of samples increases with increasing sintering
temperature.
It is also known that magnetization in ferrites is contributed by spin rotation and domain walls
displacement. Considering Fig. 5(a) and the relatively small µ’ and average grain size values for sintering
temperature ≤ 1000oC in Table 1, most of the grains corresponding to sintering temperature 1000oC and below
are single-domain grains. Therefore the permeability for samples sintered at 11000C upward is probably
dominated by both spin rotation and domain wallsforming multi-domain grains, while for samples sintered
below 10000C the permeability is contributed only by spin rotation. Thus the increasing trend of µ’ can be
attributed to the increase in density and grain size and the decrease in porosity with sintering temperature (Table
1). The large grains diminish the number of grain boundaries, therefore reducing the hindrance to the domain
wall motion. Moreover, pores become fewer, thus lessening the impediments to the domain-wall movement.
Generally the permeability of polycrystalline garnet ferrite ceramics is due to two different magnetizing
mechanisms, which are spin rotation and domain wall movement. This relationship can be described as
(Hossein, A.K.M.A., et al., 2006):
µ i = 1 + Xw + Xspin,
(2)
whereXw is the domain wall susceptibility; Xspin is intrinsic rotational susceptibility. Meanwhile the Xw and
Xspin may be written as follow:
Xw= 3πMs2D/4γ
(3)
and
Xspin = 2πMs2/K,
(4)
where Ms is the saturation magnetization, K is the total anisotropy, D is the average grain diameter, and γ is
the domain wall energy. From the formulae given, the domain wall motion is affected by the grain size and
enhanced with the increase of grain size. The initial permeability is therefore a function of grain size and
magnetization as tabulated in Table 1. Larger grains tend to consist of a greater number of domain walls. As the
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Australian Journal of Basic and Applied Sciences, 8(3) March 2014, Pages: 474-482
number of walls increases with the grain sizes, the contribution of wall movement to magnetization increases.
The effect of sintering causes the grains to grow and provides more spaces for the domain movement which is
reflected in the increase of initial permeability as seen in Table 1.
(a) 5.00E+01
4.50E+01
4.00E+01
Real Permeability, µ'
1400degreeC
3.50E+01
1300degreeC
3.00E+01
1200degreeC
2.50E+01
1100degreeC
2.00E+01
1000degreeC
900degreeC
1.50E+01
800degreeC
1.00E+01
700degreeC
5.00E+00
0.00E+00
1.00E+07
600degreeC
1.00E+08
Frequency (Hz)
1.00E+09
Fig. 1: Real permeability measured at room temperature in range of 10 MHz to 1 GHz for single-sample
sintering
Table 1: Sintering temperature, average grain size and initial permeability, µ i at 10MHz for single-sample sintering, measured at room
temperature.
Sintering Temperature (oC)
Average grain size (µm)
Density (g/cm3)
Initial Permeability, µ i at 10MHz
± 0.1
600
0.16
4.09
0.8
700
0.17
4.18
0.9
800
0.18
4.27
1.0
900
0.28
4.39
2.93
1000
0.33
4.51
4.90
1100
0.60
4.59
6.62
1200
1.14
4.63
11.5
1300
1.68
4.70
21.6
1400
2.71
4.74
36.0
Besides that, the increase in the sintering temperature leads to a decrease in the magnetic anisotropy by
decreasing the internal stress and crystal anisotropy, hence the hindrance to the domain wall motion reduces,
thereby increasing the value of the magnetic permeability (Verma, A., et al., 2005).
The results for the magnetic loss factor, µ’’, at different sintering temperatures are shown in Fig. 2. The loss
factor arises due to a lag between the magnetization or flux induction and external applied field (Gupta, N., et
al., 2005). The same trend with magnetic real permeability is observed in the case of magnetic loss factor with
respect to the frequency. The sample shows the increase of loss factor with increasing frequency, attaining the
maximum value at particular frequency and decreases with a further increase in frequency. As can be observed,
the loss factor value is increased as the sintering temperature increased. Ref. (Otsuki, E., et al., 1991) reported
that the larger grain size increases the eddy current loss because the domain wall movement is easier in the
larger grain. Thus eddy current is induced. Besides that, the fraction of grains that are occupied with domain
walls increases in the larger grain thus increasing the eddy current and hysteresis losses.The main types of loss
encountered in ferrites are the hysteresis loss, eddy current loss and residual loss. The grain size increased with
increasing sintering temperature, therefore the domain wall will easily moves in the larger grains. Hysteresis
loss can be minimized if one reduces the hindrance to domain wall motions by reducing the pinning centre to the
domain walls movement such as volume fraction of pores, impurities and dislocations, and internal strain inside
the samples (Kotnala, R.K., et al., 2010). Besides that, the hysteresis loss becomes less important at high-
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Australian Journal of Basic and Applied Sciences, 8(3) March 2014, Pages: 474-482
enough frequencies due to spin rotationat high frequency. The eddy current losses is more important at higher
frequencies as a circulating current is induced in the sample due to changing magnetic field, leading to energy
losses. However, in polycrystalline YIG, the eddy current losses can be neglected due to high resistivity
(~1014Ωcm). Residual loss plays an important role in the high frequency range, therefore to reduce the losses,
the complex permeability has to be made to peak at the high frequency as possible and this can be achieved
using fine-grained samples. Losses at high frequency are always associated with rotational or spin resonance.
This resonance involves energy absorption during its occurrence, followed by dissipation of energy to the lattice
in the form of heat (Smit, J. and H.P.J. Wijn, 1954).
2.00E+01
1.50E+01
1400degreeC
Loss Factor
1300degreeC
1200degreeC
1.00E+01
1100degreeC
1000degreeC
5.00E+00
900degreeC
800degreeC
700degreeC
0.00E+00
1.00E+07
-5.00E+00
1.00E+08
1.00E+09
600degreeC
Frequency,f (Hz)
Fig. 2: Magnetic loss factor, µ’’ from 10 MHz to 1 GHz measured at room temperature for single-sample
sintering
Curie Temperature:
The Curie temperature of the single-sample which was sintered at various temperatures is shown in
Fig.3.The single-sample sintering yielded nearly same value of Curie temperature, i.e ~267± 50C, which is
slightly lower than the Curie temperature reported by Ref. (Gilleo, M.A. and S. Geller, 1958) which is 272 0C.
The difference is believed to be due to experimental error.Curie temperature is defined by the vanishing
temperature for the spontaneous magnetization of the material.
The Curie temperature depends on the superexchange interaction (Rao, A.D., et al., 1999), where it depends
strongly on the distribution of Fe3+ ions in octahedral and tetrahedral sites which changed the magnetization. It
is also depends on the distance between those ions in both of the sites. Since significant magnetization occur
only in well-formed crystalline phase, the single-sample YIG did not show any value of Curie temperature for
samples sintered at 6000C to 8000C due to incomplete crystal structure (Table 2).
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4.5000E-04
4.0000E-04
Inductance, L
3.5000E-04
1400C
3.0000E-04
1300C
2.5000E-04
1200C
1100C
2.0000E-04
1000C
1.5000E-04
900C
1.0000E-04
800C
700C
5.0000E-05
600C
0.0000E+00
0
-5.0000E-05
50
100
150
200
250
300
Temperature (oC)
Fig. 3: The Curie temperature results for the single sample after sintering, at various temperatures.
Table 2: Curie temperature value of single-sample sintering as a function of various sintering temperatures.
Sintering Temperature (oC)
Curie Temperature for single-sample sintering
± 5 (oC)
600
700
800
900
267
1000
267
1100
267
1200
266
1300
263
1400
263
3.3 Activation Energies of the Microstructure Evolution:
Coble’s theory (Coble, R.L., 1961; Shinde, T.J., et al., 2008) mentioned that from the behavior of particle
growth, the activation energy of grain growth can be predicted using the Arrhenius equation below:
=
(5)
where k is the specific reaction rate constant, Q is the activation energy, T is the absolute temperature, and
R is the ideal gas constant. The value of k however can directly be related to grain size according to Ref. [30].
Integration of Eq. (5) would result in the equation below:
Log D = (-Q/2.303R)1/T + A
(6)
where A is the intercept on the log D vs 1/T plot and D is the grain size. By using Eq. (6), one can obtain a
best fitted straight-line plot of grain size where a plot of log D versus the 1/T is as shown in Fig. 5. This method
was adopted from Ref. (Li, X. and G. Wang, 2009). We obtained two slopes of –Q/2.303R of the lines and the
values of the activation energy of grain growth (Q) can be calculated from the Arrhenius plot in Fig. 4 having 2
stages of activation energy which are 17.33kJ/mol and 93.09kJ/mol. These two stages of activation energy show
a significant shift of the grain growth process and can be associated with grain size. As we recall, the starting
powder of the samples is in nanometer size having been processed via mechanical alloying. During the low
sintering temperatures, the activation energy was lowered due to the effect of nanosized particles. The small
particles stimulated the growth reaction by having a large surface area. Higher activation energy was needed in
order to continue the growth of grains. The shift of activation energy can also be correlated with the
microstructural evolution that affects the magnetic properties of the samples. The low activation energy
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Australian Journal of Basic and Applied Sciences, 8(3) March 2014, Pages: 474-482
(sintering temperature from 600 to 1000oC) which was due to the smaller size of starting powders which were
initially mechanically alloyed, reducing their dimensions into the nanosized region. This caused the surface area
of the starting powders to be increased and lower down the activation energy. Since the sample has been
sintered repeatedly, causing it to have a thermal history and higher activation energy needed for the second stage
of the grain growth.
3.6
3.4
y = -4861.5x + 6.3333
3.2
Log D
3
2.8
2.6
y = -905.28x + 3.1898
2.4
2.2
2
0.0005
0.0007
0.0009
0.0011
0.0013
1/T
Fig. 4: Plots of log D versus the reciprocal of absolute temperature (1/T)
Density remains almost the same but the grain size were increasing with the sintering temperature as can be
seen in Fig. 5 and Table 1. The grain growth effect has influence the formation of domain wall and initial
permeability was observed to rise significantly within this temperature range. Looking at the microstructure of
the sample sintered at 1000oC, it could be deduced that the shape and size of grains contributed dominantly to
the coercivity through shape anisotropy. No in-grain domain wall was formed for this sintering temperature. A
columnar shape of grains formed during this stage and it is not favorable for the magnetic moment rotation
compared to the equiaxed grain shape in the samples sintered at 1100oC and above. Fig. 5 shows micrographs of
single-sample sintered at 1000oC, 1100oC and 1200oC respectively. The picture clearly shows the shape grainchanges in the samples which contributed to the rise of complex permeability of the sample.
a
b
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c
Fig. 5: SEM micrographs of YIG single-sample sintered at: (a) 1000oC, (b) 1100oC and (c) 1200oC.
Conclusions:
Real permeability corresponded to the increase of sintering temperature. Increased grain size (due to
sintering process) resulted in increase of µ’. Domain walls dominated the grain, giving rise to the sample
sintered above 1100oC. This effect was also observed with the loss factor. The Curie temperature for YIG
single-sample was 267oC; itremained the same even with changes in grain size. This shows that Curie
temperature is independentof microstructural effects and relies onlyon exchange forces. Interestingly, two stages
of activation energy were found using the Arrhenius equation. The lower activation energy was due to the
higher reactivity of nanometer-sized starting powder. Reduced surface reactivity from repeated sintering of the
single-sample caused higher activation energy needed for the grain growth.
ACKNOWLEDGEMENT
The authors are thankful to Univerisiti Putra Malaysia (UPM), Malaysia for providing the Research
University Grant Scheme (RUGS) with the vot number 9357800.
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