Journal of Alloys and Compounds 622 (2015) 548–556

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Journal of Alloys and Compounds

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The effects of mechanical activation energy on the solid-state synthesisprocess of BiFeO3

http://dx.doi.org/10.1016/j.jallcom.2014.10.1350925-8388/� 2014 Elsevier B.V. All rights reserved.

⇑ Corresponding author at: School of Metallurgy and Materials Engineering,University of Tehran, P.O. Box 14395-553, Tehran, Iran. Tel.: +98 21 82084084; fax:+98 21 88006076.

E-mail address: aataie@ut.ac.ir (A. Ataie).

M. Ahmadzadeh, A. Ataie ⇑, E. MostafaviSchool of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, Tehran, Iran

a r t i c l e i n f o

Article history:Received 28 September 2014Received in revised form 26 October 2014Accepted 27 October 2014Available online 4 November 2014

Keywords:Bismuth ferriteMechanical activationMilling energyNano-structured

a b s t r a c t

The effects of milling energy induced during intermediate mechanical activation of precursors on the syn-thesis of nano-structured BiFeO3 powders have been systematically investigated. X-ray diffractometer,laser particles size analyzer, field emission scanning electron microscope, vibrating sample magnetome-ter and electrical evaluation techniques were used to study phase composition, particles size distribution,morphology, magnetic properties and ferroelectric properties of the products, respectively. Applying atotal energy of 171.18 kJ/g during milling led to formation of an amorphous structure which resultedin decreasing the formation temperature of bismuth ferrite phase by about 100 �C, although smallamounts of secondary phases were detected. This sample shows the mean particles size of 170 nm andthe mean crystallite size of 40 nm, when calcined at 750 �C. Saturation magnetization (MS) increasedfrom 0.054 to 0.071 A m2/kg and coercive field (HC) decreased from 32.63 to 6.37 kA/m by increasingthe milling energy from 13.48 to 171.18 kJ/g. In addition, electrical hysteresis loops demonstrated adecrease in the current leakage by increasing the milling energy and lowering the calcinationtemperature.

� 2014 Elsevier B.V. All rights reserved.

1. Introduction

In recent years, multiferroic materials that exhibit simultaneousmagnetic and ferroelectric orders, have become one of the mostattractive topics of materials science due to their fascinating phys-ical properties and potential applications in multifunctionaldevices [1]. BiFeO3 (BFO) is one of the most important reportedmultiferroics which simultaneously possesses both antiferromag-netic (TN � 370 �C) and ferroelectric (TC � 830 �C) order at andabove room temperature [1–3]. BFO with a rhombohedrally dis-torted perovskite structure is suitable for applications in radio,television, microwave and satellite communication, bubblememory devices, audio–video, sensors, optical filters, smartdevices, high-density ferroelectric random access memory and dig-ital recording fields [4,5].

Currently, the synthesis of single-phase BFO ceramics is achallenging issue. Moreover, nanostructured BFO exhibits uniqueproperties, compared to the bulk sample, due to its low dimension-ality and quantum confinement effect [6]. Therefore, various tech-niques including solid state and wet chemical methods have been

developed in order to synthesize nanosized BFO ceramics. Themost common technique to fabricate perovskite-type structurematerials such as BFO is conventional solid state route whichincludes the calcination of a mixture of bismuth and iron oxides[7,8]. However, materials prepared by this method, which requirehigh temperatures and long annealing times, inevitably containBi2Fe4O9 and Bi25FeO40 as impurity phases, which considerablyweaken the properties of the material [9]. Furthermore, conven-tional route has poor reproducibility and results in coarse powders[10]. Currently, in order to overcome the deficiencies of conven-tional processing and avoid complicated high-cost wet chemicalmethods, mechano-thermal synthesis is being employed. Itincludes mechanical activation (MA) of starting materials and postheat treatment. MA carried out by high-energy ball milling is usedto improve the reactivity of the precursors so that the desiredphase is formed at a reduced calcination temperature due tothe enhanced diffusion rates and the microstructural refinement[11–14]. Since increasing the calcination temperature causes thedecomposition of BiFeO3 phase to impurity phases [15,16], lower-ing the BFO formation temperature allows one to carry out synthe-sis in the region where this compound is more stable [9]. Mauryaet al. [12] reported that mechanical activation of precursorslowered the BiFeO3 phase formation temperature by about100 �C and improved the properties of powders, compared toconventionally synthesized sample. Egorysheva et al. [9] have also

M. Ahmadzadeh et al. / Journal of Alloys and Compounds 622 (2015) 548–556 549

synthesized BiFeO3 by mechanical activation assisted method,under different milling times and synthesis temperatures. Theyconcluded that mild milling enhances the purity and magneticproperties of products. However, it is necessary to investigate theeffects of mechanical activation energy in a formulated manner;to our best knowledge, it has not been studied earlier. Althoughmechanical milling is a stochastic process and the number of vari-ables involved in the process is large [17], several reports [18–20]have developed milling modeling studies in order to calculate theapplied energy to the powders during the milling process. Report-ing the milling energy instead of all milling parameters facilitatesthe study and comparison of milled samples.

This work is mainly aimed to systematically investigate theeffects of mechanical activation energy on the phase composition,formation temperature, magnetic and ferroelectric properties ofbismuth ferrite produced via mechano-thermal route.

2. Experimental

Commercial pure Bi2O3 (99%) and Fe2O3 (99%) used as starting materials. Thesepowders were accurately weighed according to the stoichiometric proportion. Thepowder mixture was milled using a high-energy planetary ball mill with hardenedsteel vial and balls under air atmosphere. Milling was carried out for 20 h at variousrotation speeds and ball to powder weight ratios (BPR). More details of milling con-ditions employed for four various milling energy are listed in Table 1. The milledsamples were calcined at different temperatures from 600 to 800 �C for 1 h witha heating rate of 10 �C/min.

The phase analysis of as-milled and calcined samples was performed usingX-ray diffraction (XRD) on a PhilipsPW-1730 X-ray diffractometer with Cu Ka radi-ation (k = 1.5406 Å). The mean crystallite size of the BFO perovskite was calculatedusing Scherrer’s formula [21]:

S ¼ Kkb cos h

ð1Þ

where K is constant, k is the wave length of Cu Ka radiation, and b is the full width athalf maxima of XRD peaks. The morphology and chemical composition of powderswere investigated using a field emission scanning electron microscope, HitachiS4160 FESEM equipped with energy dispersive spectroscopy (EDS) point chemicalanalysis. Magnetic properties of synthesized BFO perovskite were determined by avibrating sample magnetometer (VSM) under the maximum applied magnetic fieldof about 800 kA/m at room temperature. Pressed pellets of powders with the diam-eter of 10 mm and the thickness of 0.4–0.5 mm were used for ferroelectric measure-ments. The surfaces of pellets were painted by silver base electrode.

3. Theory/calculation

The milling process of the precursors has been affected by mill-ing parameters such as ball-to-powder weight ratio, rotationspeed, size of balls and milling time. All the results can be betterunderstood if the milling parameters being reported as a singleterm of ‘‘milling energy’’. Therefore, in this study, the total energytransferred during ball milling was calculated for four various lev-els of energy. The energy released by each ball per hit during themilling process is given by [18,19,22]:

DEb ¼12

mbðV2b � V2

s Þ ð2Þ

where mb is the mass of a ball, Vb and Vs are the absolute velocitybefore and after the hit.

Table 1Milling conditions of the samples.

Sampleidentity

Rotation speed(rpm)

BPR Millingtime (h)

Diameter ofball (mm)

A 200 5:1 20 15B 300 10:1 20 15C 300 20:1 20 15D 400 30:1 20 15

DEb can be calculated following the process given by Burgioet al. [18]. They have presented collision model by which Vb andVs can be obtained having milling parameters in a high-energyplanetary ball mill. DEb is calculated as:

DEb ¼ �mb½x3vðRv � rbÞ=xp þxpxvRp�ðRv � rbÞ ð3Þ

where xp and xv are the absolute angular velocity of mill plate andof one vial. Rp, Rv and rb are the radius of plate, vial and ball, respec-tively. The energy released by one ball in a system containing Nb

balls can be given as DE�b ¼ ubDEb, where ub < 1 is the empiricalfactor for different degrees of filling of the vial.

Finally, the total energy transferred per unit weight of powderfor a given milling time, t, can be written as:

DEt ¼DE�bNbf bt

Wpð4Þ

where Wp is the weight of powders in the vial and fb is thefrequency with which the balls are launched and can be expressedby:

f b ¼ Kðxp �xvÞ

2pð5Þ

where K is a proportionality constant and approximately equals to1. Since in the present study the milling time was constant (20 h)for all samples, the total energy per unit time and per unit weightof powders (P*) was calculated. P* can be written as:

P� ¼ DE�bNbf b

Wpð6Þ

The above parameters were calculated for four different levelsof milling energy evaluated in this study and are presented inTable 2.

4. Results and discussion

XRD patterns of mechanically activated starting materials infour various levels of energy as a function of milling energy areshown in Fig. 1. The results reveal that the characteristic peaks ofstarting materials (Bi2O3 and Fe2O3) tend to broaden as the millingenergy increases and their intensities decrease. These observationsare known to be due to the increasing of the strain and defects andthe crystallite refinement, without forming any new phase.Increasing the milling energy in the sample C, resulted in an amor-phous structure. Therefore, applying the energy of 171.18 kJ/g isnecessary to form an amorphous structure for the present oxide-oxide system. It can be observed that bismuth oxide amorphizationhas occurred at the first steps of mechanical milling, while thecharacteristic peaks of iron oxide exist even in sample C. This canbe related to the fact that the iron oxide is a harder reactant incomparison with the bismuth oxide [9]. Similar disordered struc-tures can also be found in the previous literature for the similarhigh-energy ball-milled oxide-oxide systems [9,23,24]. Mechanicalactivation of starting oxides resulted in the following steps: initialdistortion of the crystal structure of reactants; formation,accumulation, and interaction of linear and point defects; anddisintegration of the materials. By increasing the milling energy,

Table 2Milling energy parameters of the samples.

Sample identity ub DE�b (J) fb (1/s) P⁄ (J/g s) DEt (kJ/g)

A 0.99 5.02 � 10�2 10 0.19 13.48B 0.98 11.15 � 10�2 15 1.24 89.86C 0.94 10.62 � 10�2 15 2.38 171.18D 0.85 17.18 � 10�2 20 7.69 553.86

Fig. 1. XRD patterns of as-milled samples under various milling energy.

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the reaction surface and the reactivity of reactants will increase. Itis observed that mechanical activation leads to crystal structuredisappearing which contributes to formation of desired phase inlower temperatures. However, the XRD result of sample D withthe highest energy showed a crystalized bismuth phase. Thepresence of contamination in the powders is suspected when anunexpected phase appears to be present in the milled powder.Thus, it seems that the powder contamination played a crucial rolein this phenomenon. The magnitude of contamination, whichdepends on milling energy, reaches its maximum in the sampleD. In this experiment, since the main source of contamination is

Fig. 2. (a) Room temperature magnetic hysteresis loop and (b) EDS p

the milling media which has been made of Fe alloys, the impurityis primarily iron.

The magnetic hysteresis loop of the sample D in Fig. 2a demon-strates a saturation magnetization (MS) of 4.66 A m2/kg. Neverthe-less, bismuth as a detected phase in the XRD pattern exhibitsdiamagnetic behavior [25]. Moreover, as possible present phases,Fe2O3 is a canted antiferromagnet which shows a weak ferromag-netism at room temperature with MS � 0.4 A m2/kg [26] and Bi2O3

shows diamagnetic properties at ambient temperature [27,28].Therefore, the relatively high saturation magnetization of the sam-ple D could not be due to the existence of the above phases. It

oint analysis of sample D. The insets show the obtained values.

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seems the presence of Fe contamination can be inferred from themagnetic properties with the high MS. However, the characteristicpeaks of Fe did not appear in the XRD patterns of sample D due tothe low percentage of Fe in the powder and the lower molecularmass of Fe compared to Bi (MFe = 55.84, MBi = 208.98 g/mol).Additionally, the EDS analysis of this sample in Fig. 2(b) provesthe existence of iron and oxygen besides bismuth.

According to the Gibbs free energies of this compounds(DG0

f ðBi2O3Þ ¼ �493:24, DG0f ðFe2O3Þ ¼ �741:53 and DG0

f ðFe3O4Þ ¼�1014:48 kJ=mol), some of bismuth oxide has been probablyreduced by Fe-based impurities. The reduced bismuth crystalizedowing to local high temperature resulted from applying highmechanical energy. As a result, the peaks of bismuth appearedin the XRD pattern. It is worth mentioning that the abovevalues are the Gibbs free energies at room temperature; but thelocal temperature would be slightly higher although the powderwas milled at room temperature. Nevertheless, the values at realtemperature should not be far from these ones. It should be men-tioned that due to the undesirable contamination in the sample D,the data obtained for this sample was excluded from the rest ofthis report.

Fig. 3 presents the FESEM micrographs of bismuth and ironoxides used as starting materials as well as their mixture prior tomilling. Bismuth and iron oxides particles show rod-like andalmost spherical shapes, respectively. As expected, the mixturemicrograph reveals both oxides particles together in which theyhave not been uniformly mixed.

The particles size and uniformity of milled powders can influ-ence the properties of final products. Therefore, the typical FESEMmicrographs and the particle size distributions of the samplesmilled under different milling energy are shown in Figs. 4 and 5,respectively. It is obvious that the morphology and size distribu-tion of milled particles depend on milling energy. Increasing themilling energy changed the morphology from plate-like particleswith a mean thickness of 45 nm to uniform-size spherical nano-particles with a mean size of 30 nm, which are arranged within

Fig. 3. FESEM micrographs of (a) mixture of bismuth and iron oxides before milling,(b) iron oxide and (c) bismuth oxide.

hard agglomerates. The particle shape of the as-milled powdersmainly depends on the thermodynamics and kinetics of reactions.Due to the heavy work hardening of as-milled powder, irregularityin the shape of mechanically milled powder is expected [29,30]. Itis worth mentioning that after milling for a certain length of time,steady state stage is attained. At this stage, the composition ofevery powder particle is the same as the proportion of the ele-ments in the starting powder mix. As a general rule, the timestaken to achieve the steady-state conditions are short for a high-energy ball milling [17]. LPSA results (Fig. 5) show that by increas-ing the milling energy, the particle size distribution of as-milledsamples deviates from normal distribution and the mean diameterof particles increases from 2 to 4.7 lm. It worth noting that, in theLPSA results, the particle size can be meant by agglomerate size.Cold welding and fracturing which occurred repeatedly duringthe milling period determine the particles size. Due to the verylow milling energy of sample A, the large particles were fracturedand then plastically deformed. The plate-like shape of particles inthe FESEM image of this sample is an indicative of this fact. Byincreasing the milling energy in the samples B and C, the work-hardened particles tended to fracture more which led to a large

Fig. 4. FESEM micrographs of as-milled samples A (a), B (b) and C (c).

Fig. 5. Particle size distributions of as-milled samples A (a), B (b) and C (c). Theinsets show the obtained values.

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area of new fresh surfaces. These new fresh surfaces encourage thecold welding process among the particles. This repeated cold weld-ing and fracturing of particles caused the wide size distribution inthe sample C.

Heat treatment is necessary in order to complete the solid-statereaction between the oxide precursors and form BFO phase, asreported earlier [9,12,23,24]. The XRD patterns of samples A, Band C after calcination at various temperatures of 600, 700, 750and 800 �C are shown in Fig. 6. The amount of BiFeO3 phase is

affected by milling energy, particularly at lower calcination tem-peratures. The XRD results of sample D after calcination at differenttemperatures (not shown here), indicated a large amount of impu-rity phases. The BiFeO3 phase has started to form at 600 �C and700 �C in the samples C and A, respectively. The main impurityphases formed during the calcination are mullite (Bi2Fe4O9) andsillenite (Bi25FeO40). In addition, by increasing the calcination tem-perature in each sample, the BiFeO3 content increased. Rising themechanical activation energy results in increasing the crystaldefects such as dislocations, vacancies, stacking faults, and grainboundaries [17]. The higher defect density enhances the diffusivityof reactants and increases the nucleation sites. Furthermore, themicrostructural refinement induced by mechanical activationdecreases the diffusion distances. Consequently, increasing themilling energy leads to the formation of BFO at lower tempera-tures. These XRD results show that the formation temperature ofBFO phase was lowered by about 100 �C in the sample C.

The mean crystallite sizes of calcined samples as a function ofmilling energy and calcination temperature are shown in Fig. 7.Increasing the milling energy from 13.48 to 171.18 kJ/g in the sam-ples calcined at 750 �C reduces the mean crystallite size from 48 to40 nm. This was expected due to the larger numbers of defects thatacted as nucleation sites in the sample C. Increasing the calcinationtemperature of sample C from 700 to 800 �C resulted in an increasein the mean crystallite size from 33 to 45 nm which is related tohigh diffusivity at higher temperatures.

The effects of milling energy on the morphology of the productscan be seen in Fig. 8 which shows the FESEM images of samples A,B and C after calcination at 750 �C. By increasing the millingenergy, the morphology of particles changed from faceted particleswith the mean size of 330 nm (sample A) to semi-spherical parti-cles with the mean size of 170 nm (sample C). The shape of parti-cles in sample A-750 can be associated with the grain growth habit,where the general particles morphology acquires the unit cellshape. Santos et al. [23] have reported that the faceted grains wereassociated to Bi-rich phase (Bi25FeO40 phase, cubic symmetry)which is the main secondary phase of sample A-750, consideringthe XRD results. It is seen that particles growth has been limitedowing to the higher defects density and more nucleation sitesinduced by intensive milling in the sample C-750.

The room temperature M–H loops of samples A, B and C aftercalcination at 750 �C are shown in Fig. 9. The thin and linearfield-dependent loops suggest an antiferromagnetically orderedstate which was reported earlier for BiFeO3 [24,31]. The localshort-range magnetic ordering of BiFeO3 is the G-type antiferro-magnet in which each Fe3+ with spin-up is surrounded by six ofthe nearest Fe neighbors with spin-down [32]. The results revealedthat the coercive field (HC) decreased from 32.63 to 6.37 kA/m byincreasing the milling energy from 13.48 (sample A) to 171.18 kJ/g (sample C). The thicker loop of sample A-750 is probably relatedto the existence of more content of secondary phases, Bi25FeO40

specifically, which has been confirmed by XRD results. The resultsalso showed that by increasing the milling energy, the maximummagnetization slightly increased from 0.054 in sample A-750 to0.071 A m2/kg in the sample C-750. This behavior has beenreported earlier by Maurya et al. [12]. They suggested that theincrease in defects induced by mechanical activation causes anincrease in oxygen vacancies, which may give rise to the formationof Fe2+ ions. It may result in enhanced magnetization in sampleC-750. In addition, an increase in the spin canting, which resultedfrom relatively higher degree of disorder in this sample, can also bea contributor to the magnetism.

Fig. 10 presents the polarization (P) versus electric field (E)loops for samples A, B and C after calcination at 750 �C. The resultsshowed that ferroelectric properties of BFO samples depend on theapplied energy to the powders during mechanical activation of

Fig. 6. XRD patterns of milled samples A (a), B (b) and C (c) calcined at various temperatures from 600 to 800 �C for 1 h.

M. Ahmadzadeh et al. / Journal of Alloys and Compounds 622 (2015) 548–556 553

precursors. Fig. 10(a) shows a loose-shape loop for sample A-750,which indicates a high level of leakage current, even at low appliedelectric fields (�4 kV/cm). The leakage is far lower in the samples

B-750 and C-750, which have been mechanically activated underhigher milling energy. It is known that the large leakage currentin BFO is attributed to existence of Fe ions in a mixed valence state

Fig. 7. Mean crystallite size of synthesized samples as a function of calcinationtemperature and milling energy.

Fig. 8. FESEM micrographs of samples A (a), B (b) and C (c) all calcined at 750 �C for1 h.

Fig. 9. Room temperature magnetic hysteresis loops of samples A, B and C aftercalcination at 750 �C for 1 h. The inset shows the enlarged M–H curves.

Fig. 10. Room temperature polarization–electric field hysteresis loops of samples A(a), B (b) and C (c) all calcined at 750 �C for 1 h, measured at the frequency of 10 Hz.

554 M. Ahmadzadeh et al. / Journal of Alloys and Compounds 622 (2015) 548–556

Fig. 11. Room temperature polarization–electric field hysteresis loops of sample Ccalcined at (a) 700, (b) 750 and (c) 800 �C for 1 h, measured at the frequency of 1 Hz.

M. Ahmadzadeh et al. / Journal of Alloys and Compounds 622 (2015) 548–556 555

(Fe2+/Fe3+) which causes oxygen vacancies in order to compensatefor the change in charge [33]. In addition to oxygen vacancies, leak-age current is affected by defects density, porosity and phase pur-ity. The effect of calcination temperature on the ferroelectricproperties can be seen in Fig. 11, which shows the room tempera-ture P–E hysteresis loops of sample C after calcination at various

temperatures. As indicated, the leakage current slightly decreasedby increasing the calcination temperature. In this study, thedecrease in the leakage current possibly arose from reducing theconducting secondary phases as a result of increasing the millingenergy and calcination temperature. It is seen that all samples donot show saturated hysteresis loops as reported previously in somestudies for solid-state synthesized BFO [12,34,35]. The relativelylow ferroelectricity of the samples is most likely due to the second-ary phases existing in the powders. However, in the presence ofconducting phases, to suppress the leakage and extract the intrin-sic behavior of BiFeO3, low temperature measurements will beneeded [12].

5. Conclusion

Nano-structured BiFeO3 powders were synthesized via mechan-o-thermal route. The applied milling energy was calculatedthrough collision model and its influence on the synthesis processwas investigated. Although a single-phase BiFeO3 was notobtained, applying the total milling energy up to an optimum valueof 171.18 kJ/g during mechanical activation of precursors resultedin forming of the BiFeO3 phase at relatively lower temperature.However, applying the high energy of 553.86 kJ/g caused powderscontamination which resulted in forming undesirable phases. Forthe powders calcined at 750 �C, the sample with the applied opti-mum energy showed an antiferromagnetically ordered state andless leakage current. Moreover, increasing the milling energy ledto decreasing the mean crystallite and particles size of the samplesto 40 and 170 nm, respectively, for the samples calcined at 750 �C.

Acknowledgements

The financial supports of this work by the University ofTehran and Iran Nanotechnology Initiative Council are gratefullyacknowledged.

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