49

STUDY OF PURE AND ZIRCONIUM (Zr) SUBSTITUTED

BISMUTH FERRITE

3.1 INTRODUCTION

BiFeO3 (BFO) is one of the most candid material amongst all the multiferroic materials

which combines both ferroic properties i.e. antiferromagnetism and ferroelectricity in the same

phase having distorted rhombohedral perovskite structure [121]. In spite of all the favorable

features, applications of BFO are still limited because of the several impediments such as

secondary phase formation, weak magnetic behaviour and high leakage current. Many strategies

like doping in the BFO host lattice either on the Bi-site, Fe-site or at both sites have been adopted

in order to stabilize the perovskite phase, reduce leakage current, and improve ferromagnetic

properties of BFO [122, 123]. Various authors have reported B-site doping by high valance Zr4+,

Nb5+, Mn4+, V5+ and Ti4+ ions, which proved useful in increasing electrical resistivity and

enhancing overall multiferroic properties of BFO. A number of considerable works has been

done for Zr doped and codoped BFO solid state ceramics and thin film. However, synthesis of

pure phase is still a challenge. Therefore, in this chapter, we have prepared pure phase BFO

nanoparticles by optimizing synthesis techniques and studied its structural, magnetic and optical

properties in detail. In addition to pure BFO, Zr4+ doped BFO nanoparticles have also been

prepared and detailed study on the effect of Zr4+ doping on structural, magnetic, dielectric, optical

and impendence properties has been reported in this Chapter.

3.2 EXPERIMENTAL DETAILS

Pure and Zr4+ doped BiFeO3 (BiFe1-xZrxO3; x=0.0, 0.03, 0.07, 0.10 and 0.15)

nanoparticles were prepared by sol-gel method. First the synthesis conditions have been

optimized for the preparation of pure BFO nanoparticles by using citric acid and tartaric acid

route. For the synthesis of BFO nanoparticles from citric acid, Bi(NO3)3·5H2O, Fe(NO3)3·9H2O,

citric acid and ethylene glycol were used to prepare precursor solution. An aqueous solution of

citric acid was prepared in distilled water and then Bi(NO3)3·5H2O and Fe(NO3)3·9H2O were

added with constant stirring at 50–60ºC followed by addition of ethylene glycol to the solution

CHAPTER 3

CHAPTER 3

50

with citric acid/ethylene glycol proportion (70:30). The solution was heated at 100ºC on hot plate

until gel formation occurs after evaporation of water. The dried powder was then sintered at

550ºC for 2 hours to get desired phase of BFO nanoparticles. In second method, synthesis of

BFO nanoparticles was done by tartaric acid route using tartaric acid as a template material and

diluted nitric acid as oxidizing agent. Stoichiometric amount of Bi(NO3)3·5H2O and

Fe(NO3)3·9H2O (Sigma Aldrich, 98 %) were dissolved in deionized water along with small

amount of diluted nitric acid. This was followed by addition of tartaric acid (ratio of tartaric acid

and metal nitrates was 1:2 mole). This solution was heated up to 60ºC for 12 hours with constant

stirring followed by heat treatment in an oven for 4-5 hours in order to obtain fluffy light yellow

color gel. The dried powder was calcined at 550ºC for 2 hours in order to get crystalline phase.

The phase identification of prepared samples was done by XRD technique and it was observed

that tartaric acid based sol-gel route has been found effective in getting pre phase BFO

nanoparticles.

After optimizing the synthesis conditions for pure BFO nanoparticles, rest of the doped

BFO nanoparticles studied in the present thesis have been prepared by tartaric acid based sol-gel

route. For the synthesis of Zr4+ doped BFO nanoparticles, equimolar mixture of Bi(NO3)3.5H2O,

Fe(NO3)3.9H2O and ZrO(NO3)2.xH2O were dissolved in de-ionized water before adding dilute

nitric acid (65-68% HNO3) to the mixture. Tartaric acid (1:2 moles to that of metal nitrates) was

added to the solutions as a chelating agent which helps in controlling shape and size of the

material. Further, the homogenous solution was heated at 50 to 60°C on the hot plate for 12 hours

and was placed in the oven at 100°C for 4-5 hours to obtain fluffy gel. Calcination of dried

powders was performed at 550°C for 2 h in order to obtain crystallized BFO nanoparticles.

The crystal structure of pure and BiFe1-xZrxO3 nanoparticles was examined using X-Ray

diffractometer (Brucker D8 Advanced) with CuKα radiation. Rietveld Refinement was carried

out to find the lattice parameters, atomic positions, bond angles, and various R-factors by using

FULLPROF software. Transmission electron microscopy (TEM) images were taken on (Philips

CM 10) in order to analyze particle size of the samples. Raman analysis was performed by

Renishaw Raman spectrometer (50 mW, 514.5 nm, Ar+ laser). The electrical measurements were

carried on silver coated pellets using LCR meter (PSM 1735) in the temperature range (40-

500oC) at various frequencies. The magnetic properties of the samples were measured by using a

vibrating sample magnetometer (VSM: Lakeshore 7400 series) at room temperature. Optical

CHAPTER 3

51

properties were studied by UV-visible diffuse reflectance spectroscopy. FTIR spectra were

obtained on (Perkin Elmer Spectrum BX-II) in the range 400-4000 cm-1.

3.3 RESULTS AND DISCUSSION

3.3.1 STRUCTURAL STUDIES

3.3.1.1 X-RAY DIFFRACTION STUDIES

Figure 3.1: (a) XRD patterns of BFO nanoparticles prepared by Citric acid route and (b) followed by leaching. (c)

XRD pattern of BFO nanoparticles prepared by tartaric acid route and (d) Rietveld refinement of XRD pattern of

BFO nanoparticles prepared by Tartaric acid route.

Figure 3.1 shows XRD patterns of pure BFO nanoparticles (sintered at 550°C) prepared

by sol-gel method via citric and tartaric acid routes, respectively. It can be seen that majority of

the peaks belong to distorted rhombohedral structure of BFO. The XRD pattern of pure BFO

sample shows the typical peaks of rhombohedral structure (JCPDS card number 71-2494) along

with some impurity peaks of Bi36Fe24O57 and Bi2Fe4O9 for sample prepared by citric acid route

CHAPTER 3

52

(Figure 3.1(a)). Leaching was done in dilute nitric acid to remove impurity phases present in

BFO powder prepared by citric acid route.

Figure 3.1(b) shows XRD pattern of BFO nanoparticles prepared by citric acid route after

leaching in dilute nitric acid. Leaching was effective in reducing impurity phases; however, few

impurity peaks still appeared in the XRD pattern. On the contrary, we obtained pure phase BFO

nanoparticles via tartaric acid route (Figure. 3.1 (c)) even without leaching in dilute nitric acid.

Further characterizations of BFO nanoparticles prepared by tartaric acid route were done and

discussed in this paper. All XRD peaks for sample prepared by tartaric acid route has been

indexed according to rhombohedrally distorted perovskite structure with space group R3c.

Rietveld Refinement of XRD pattern of BFO nanoparticles prepared by tartaric acid route has

been performed to investigate various structural parameters. The difference in the experimental

and simulated data is shown in Figure 3.1 (d). It was observed from fitting parameters that the

structure is very well fitted for rhombohedral system. The lattice parameters and unit cell volume

were calculated as a = b = 5.5785 Å, c = 13.8654 Å and 373.6829 (Å)3, respectively. The

difference in the experimental and simulated data is shown in Figure 3.1 (d). It was observed

from fitting parameters that the structure is very well fitted for rhombohedral system.

Figure 3.2 shows the XRD patterns of Zr4+ doped BFO nanoparticles with varying

concentrations. The phase formation without any trace of impurity has been observed for all

samples. The XRD pattern of pure BFO sample (Figure 3.1) has shown the typical peaks of

rhombohedral structure (JCPDS card number 71-2494). The distinguishable peaks from (006)

and (202) planes around 2θ values 39° were observed for pure BFO sample. However, with

increasing Zr4+ concentration from x=0.03 and 0.15 samples, the peaks merging tendency has

been observed. The reflections from (104) and (110) planes in the 2θ range of 31.5o-32.5o and the

reflections from (006) and (202) planes in the 2θ range of 38.5o-40.5o tend to merge, for which

the dopant substitution on Fe site may stand accountable. For x=0.07 and 0.10 samples, the peaks

from (104) and (110) planes merge into a single peak along with separated peaks from (006) and

(202) planes suggest the phase evolution of triclinic structure. Further, the reflections of (006)

and (202) planes converge into a single peak for x=0.15 sample exhibiting substitution induced

structural phase transformation from rhombohedral to triclinic phase with P1 space group

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53

Figure 3.2: Room temperature XRD patterns for BiFe1-xZrxO3 nanoparticles with x = 0.03-0.15.

This merging of the peaks is attributed to the reduced crystallite size with increased

doping, due to which peak broadening occurs. Such type of structural transformation has also

been reported in Eu doped BFO and Nd doped BFO [124, 125]. The merging of peaks indicates

the distortion in the rhombohedral structure. However, the structural phase transformation is not

evident up to x=0.10 sample. The structural distortion in perovskite structure can be understood

by Goldschmidt’s tolerance factor. The perovskite structure can be described by Goldschmidt’s

tolerance factor which is an indicator of stability and structure distortion in the samples [126].

Phase evolution from distorted rhombohedral structure to triclinic structure can be understood by

tolerance factor which is given by:

[ ]

( )[ ]−++

−+

++−

+=

243

23

12OZrFe

OBi

rxrrx

rrt (3.1)

where +3Bi

r , +3Fe

r , −2O

r , and +4Zr

r , are the atomic radii of Bi3+, Fe3+, O2- and Zr4+ ions

respectively. In case of cubic structure, the value of t lies between 0.90<t<1 and the structure is

distorted perovskite for 0.08<t<0.89. Here, the calculated values of t were 0.900, 0.890, 0.888,

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54

0.887 and 0.885 for x = 0.0, 0.03, 0.07, 0.10 and 0.15 samples, respectively, indicating the

distortion in the structure. The value of t decreases with increasing Zr4+ substitution which further

acts as the driving force for increased octahedral rotation. Alteration in octahedral rotation results

in structural distortion and subsequently the phase evolution from rhombohedral structure to

triclinic structure with increasing amount of Zr4+ in the BFO host lattice.

Figure 3.3: Rietveld Refinement XRD patterns of BiFe1-xZrxO3 samples and enlarged view of merging (104) and

(110) planes and (006) and (202) respectively.

For detailed structural analysis, the XRD patterns of Zr4+ doped BFO samples were

simulated using Rietveld refinement program (using the software FULLPROF) in which the

x=0.03, 0.07 and 0.10 samples were refined for rhombohedral structure with R3c space group

whereas for x=0.15 sample, triclinic structure with space group P1 was utilized. The refinement

was carried out by considering some fixed parameters, co-ordinates of Bi (000), Fe (00z) and

angles α=90°, β=90°, γ=120° for rhombohedral structure whereas co-ordinates of Bi (000), Fe

(xyz) and angles α=90°, β=90°, γ=90° for triclinic structure. The absence of impurity peaks

suggested the incorporation of Zr4+ ions in the BFO host lattice and hence the observed and

calculated patterns matched with each other. Larger ionic radius of Zr4+ ions (0.72 Å) as

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55

compared to Fe3+ ions (0.64 Å) ensures the increase in lattice parameters which increase

substantially with increasing Zr4+ concentrations. The Fe-O-Fe bond angle controls the magnetic

exchange interaction in the BFO lattice and any change in the Fe-O-Fe bond angle contributes to

the magnetic properties of the compositions [126]. Lattice parameters, Bond angles and various

R-factors obtained from Rietveld refinement for x = 0.0 to 0.15 samples are listed in Table 3.1.

Table 3.1: Rietveld Refined structural parameters for BiFe1-xZrxO3 samples

Compo-

-sitions

Crystal

Structure Lattice

Parameters Atoms x y z

R-Factors

(%) and

Fe-O-Fe angle

x = 0.0

Rhombohedral

a = 5.578 (Å)

c = 13.865 (Å)

V = 373.682 (Å3)

Bi

Fe

O

0.0

0.0

0.4491

0.0

0.0

0.0006

0.0

0.2247

0.9564

Rp = 6.78

Rwp = 9.87

RBragg = 6.38

Rf = 6.88

Fe-O-Fe=159.87

x = 0.03

Rhombohedral

a = 5.588 (Å)

c = 13.840 (Å)

V = 374.326 (Å3)

Bi

Fe/Zr

O

0.0

0.0

0.4636

0.0

0.0

0.0378

0.0

0.2234

0.9530

Rp= 5.23

Rwp = 6.88

RBragg = 5.13

Rf = 4.09

Fe-O-Fe=154.04

x = 0.07

Rhombohedral

a = 5.595 (Å)

c = 13.849 (Å)

V = 375.554 (Å3)

Bi

Fe/Zr

O

0.0

0.0

0.4700

0.0

0.0

0.0354

0.0

0.2251

0.9491

Rp = 5.22

Rwp = 6.69

RBragg = 3.68

Rf = 4.00

Fe-O-Fe=153.07

x = 0.10

Rhombohedral

a = 5.604(Å)

c = 13.854 (Å)

V = 376.748 (Å3)

Bi

Fe/Zr

O

0.0

0.0

0.4686

0.0

0.0

0.0483

0.0

0.2230

0.9592

Rp = 4.71

Rwp = 5.83

RBragg = 3.60

Rf = 3.00

Fe-O-Fe=155.71

x = 0.15

Triclinic

a = 3.965(Å)

c = 3.973(Å)

V = 62.88(Å3)

Bi

Fe/Zr

O

0.0

0.4871

-0.0556

0.0

0.4263

0.4182

0.0

0.4986

0.6538

Rp = 4.97

Rwp = 6.93

RBragg = 7.54

Rf = 7.53

Fe-O-Fe=144.20

The crystallite size was calculated from the X-ray line broadening by using Debye-Scherer's

formula: dXRD = Kλ/β(θ)cosθ, where d is the crystallite size, K is a constant, λ is the x-ray

wavelength, β(θ) is the Full width at half maxima (FWHM) of the diffraction line and θ is the

CHAPTER 3

56

angle of diffraction. The crystallite sizes were calculated to be 40, 30, 27, 26.5, 21 nm for x=0.0,

0.03, 0.07, 0.10 and 0.15 samples, respectively.

3.3.1.2 FTIR STUDIES

Figure 3.4: FTIR spectra for BFO nanoparticles sintered at different temperatures.

To study the bond formation and vibrational study in present samples, FTIR spectroscopy

has been utilized. Figure 3.4 shows FTIR spectra of pure BFO nanoparticles (sintered at 250°C,

350°C, 450°C and 550°C) in the wave number range 400–1000 cm−1. It is clear that pure BFO

nanoparticles sintered at 250°C and 350°C did not show any absorption indicating the absence of

bond formation. However, the sample sintered at 450°C showed two weak broad absorption

peaks around 450 cm−1 and 550 cm−1, respectively. On further increasing the sintering

temperature, intensity of broad absorption peaks increases indicating the completion of bond

formation and high degree of crystallization. The broad absorption peaks at 550 cm−1 and 440

cm−1 are due to overlapping of Fe–O and Bi–O group vibration. Absorption peaks at 550 cm−1

and 440 cm−1 are the characteristics of Fe–O stretching and bending vibrations of FeO6 octahedral

group, respectively [127]. BiO6 octahedral structure unit also possesses absorption peaks at 525

cm−1 and 450 cm−1 [128]. Peak corresponding to 667 cm−1 is attributed to water absorption from

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57

environment. As the calcinations temperature increases these peaks becomes prominent which

confirms the Bi–O and Fe–O bonds formation at 550°C.

Figure 3.5 shows the room temperature FTIR spectra for Zr4+ doped BFO samples. The

broad band between 3500-3000 cm-1 is assigned to the stretching of O-H vibration of the water

molecule. The intense peak at 1630 cm-1 is attributed to the bending vibration of H2O while the

band at about 1382 cm-1 corresponds to the trapped nitrates [129]. Two vibration bands at 550

cm-1 and 443 cm-1 are assigned to Fe-O stretching and bending vibrations, respectively, presents

in FeO6 octahedral unit and bands at 525 cm-1 is designated to BiO6 octahedral structure unit

[130, 131]. The appeared broad nature of vibration bands is due to occurrence of absorption

peaks of FeO6 and BiO6 octahedra at nearly the same frequency. The peaks present at 1597, 1350

and 1448 cm-1 are of KBr as shown in the figure. The formation of the perovskite structure in the

present samples can be confirmed by the formation of metal-oxygen bonds.

Figure 3.5: Room temperature FTIR spectra for BiFe1-xZrxO3 samples in the wave number range 500-

4,000cm-1

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58

3.3.1.3 RAMAN SPECTROSCOPY STUDIES

Figure 3.6: Room temperature Raman spectra of BFO nanoparticles

To further explore the structure and to study the effect of dopant on structure, Raman

spectroscopy has been employed. Raman scattering spectra sensitive to atomic displacement give

important information about the structural phase transition and spin-phonon coupling in the

sample. Figure 3.6 shows the measured Raman spectrum of pure BFO nanoparticles at room

temperature. The Raman active modes for rhombohedral structure with R3c space group can be

summarized using irreducible presentation: Γopt = 4A1 + 5A2 + 9E, where A1 modes and E modes

are Raman active, whereasA2 modes remain inactive [132, 133]. The stereochemical activity of

Bi–O lone pair electrons plays main role in changing of A1-1, A1-2, A1-3 modes. The low

frequency E modes (<400 cm-1) are responsible for the ferroelectric nature of the BFO whereas

higher frequency E modes (>400 cm-1) are attributed to change in Fe–O bond [134, 135]. The

measured spectrum for pure BFO sample was fitted and decomposed into individual Lorentzian

components, which is also shown in Figure 3.7.

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59

Figure 3.7: Deconvoluted Raman spectra of BFO nanoparticles

Figure 3.8 represents the deconvoluted Raman spectra of Zr4+ doped BFO nanoparticles.

The modification of the B-site by Zr4+ substitution results in a transition to a triclinic symmetry

with the space group P1. For x=0.03-0.10 samples, 11 (4A1+7E) Raman modes were observed

indicating the distortion in the rhombohedral structure with Zr4+ substitution whereas the

suppression of E-1 mode indicate the structural transformation in the x=0.15 sample consistence

with the XRD results. The observed Raman modes for all the compositions are in good

agreement with the reported data for BFO single crystals [136]. The position of A1-1, A1-2, A1-3

modes, which are governed by Bi-O bonding and is responsible for the ferroelectric ordering,

does not show any shifting with increasing Zr concentration. This suggests that Zr4+ ions get

substituted at Fe-site only and thus does not make any significant change in A1 modes. To study

the effect of Zr4+ doping at Fe3+ site in BFO lattice, higher frequency (>400 cm-1) modes have

been analyzed. The higher frequency deconvoluted Raman modes are shown in Figure 3.9 (b-e).

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60

Figure 3.8: Deconvoluted Raman spectra of BiFe1-xZrxO3 (x = 0.03, 0.07, 0.10 and 0.15) nanoparticles

The higher frequency (>400 cm-1) E-8 and E-9 modes are related with Fe-O1 (O1: axial

ion) and Fe-O2 (O2: equatorial ion) bonds, respectively, which are related to octahedral rotation

critical to weak ferromagnetism [135, 137]. The oxygen rotation influences the

antiferromagnetism in BFO through superexchange interaction, which depends on bond angles.

Thus any distortion in the antiferromagnetic structure should be reflected in the evolution of

Raman spectra and, therefore, different evolution of E-8 and E-9 modes may be attributed to

spin-phonon coupling. It is clear from Figure 3.9 (b-e) that the intensity of E-9 mode reduces for

x=0.03 to 0.07 and then increases for x=0.10 to 0.15 samples. The reduction and later on the

enhancement in intensity of E-9 mode indicate the structural distortion induced spin-phonon

coupling. This spin-phonon coupling predicts that magnetic property of BiFe1-xZrxO3 samples

might increase for x=0.03-0.07 and then decrease for x=0.10-0.15 samples. The comparison of

reported Raman data on thin film by Yang et al [138] and on BFO polycrystalline by Kothari et

al., [133] respectively, and observed Raman modes for BiFe1-xZrxO3 (x = 0.0, 0.03, 0.07, 0.10

and 0.15) nanoparticles is shown in Table 3.2.

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61

Figure 3.9: Measured Raman spectra and decomposed Raman active modes for BiFe1-xZrxO3 samples at room

temperature

Table 3.2: Observed and reported Raman modes of BiFe1-xZrxO3 nanoparticles at room temperature with wave-

number range 100-700 cm-1

Raman modes

(cm-1)

Yang et al.

(cm-1)

Kothari et

al.

(cm-1)

x = 0.0

(cm-1)

x = 0.03

(cm-1)

x = 0.07

(cm-1)

x = 0.10

(cm-1)

x = 0.15

(cm-1)

E -- 98.36 112.7 124.9 123.8 123.7 --

A1-1 139 135.15 136.4 138.0 138.0 138.0 138.2

A1-2 171 167.08 171.0 161.3 161.0 161.0 167.9

A1-3 217 218.11 216.7 204.5 204.2 -- 204.8

E 260 -- 259.9 255.1 254.4 254.6 --

E 274 283 277.3 274.0 271.3 274.9 278.2

E 306 321.47 304.5 315.6 314.0 320.0 324.3

E 344 351.55 344.3 -- -- -- --

E 368 -- 368.0 375.1 372.3 371.6 361.7

A1-4 430 430.95 426.4 430.0 428.9 425.4 424.1

E 468 467.6 467.9 -- -- -- --

E 520 526.22 521.0 528.6 528.2 528.4 540.3

E 611 598 601.2 621.0 620.1 631.5 638.6

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62

3.3.1.4 MORPHOLOGICAL STUDIES

Figure 3.10: TEM images and SAED patterns of BiFe1-xZrxO3 nanoparticles with x=0.0, 0.03 and 0.15

To investigate the particle size and morphology of the nanoparticles, transmission

electron microscopy (TEM) was performed using carbon coated copper grid. Figure 3.10 shows

TEM images and their corresponding SAED patterns for x = 0.0, 0.03 and 0.15 nanoparticles.

These micrographs indicate crystallized nanoparticles wherein particle size decreases with

increasing Zr concentration in the samples. The average particle size of x = 0.03 and 0.15

samples are 35 and 25 nm respectively. Subsequent decrease in particle size was obtained with

increasing doping concentration which may be ascribed to higher valence of dopant ions than

x = 0.0 x = 0.0

x = 0.03 x = 0.03

x = 0.15 x = 0.15

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63

Fe3+ions. The addition of Zr4+ ions at Fe3+ site requires charge compensation, which may be

achieved with the filling of oxygen vacancies and this results in lower grain growth and hence

particle size reduces.

The presence of sharp diffraction spots indicates the formation of crystalline nanoparticles which

is a confirmation of phase purity of the sample in agreement with the XRD studies.

3.3.2 MAGNETIC STUDIES

3.3.2.1 MAGNETIZATION VERSUS MAGNETIC FIELD (M-H) STUDIES

BFO exhibits very interesting magnetic properties, as being antiferromagnetic displays no

net magnetization in bulk form due to spiral spin structure. However, latent magnetization in

BFO can be released by structural transformation/partial structural distortion and reducing the

particle size in nano range. As we have prepared pure and Zr4+ doped BFO nanoparticles, it is

expected to observe enhanced magnetization in these samples. Therefore, to study the magnetic

properties of prepared samples, VSM measurements were carried out. Figure 3.11 shows the

magnetic hysteresis loop of pure BFO nanoparticles measured at room temperature, indicating

ferromagnetic (superparamagnetic) behavior. It is well known that bulk BFO possesses cycloid

spin of order 62 nm and no net magnetization at room temperature. The possible explanation

which can describe the origin of net magnetization in pure BFO nanoparticles could be its size

dependent properties; As the particle size reduces, surface to volume ratio becomes very large

which results in the structure distortion and the modification of the spiral G-type

antiferromagnetic ordering to the collinear G type antiferromagnetic structure in which the canted

components becomes measureable. Moreover, due to large surface to volume ratio in case of

nanoparticles, the antiferromagnetic arrangement of Fe3+ spin often gets disrupted at the surface

which gives rise to canting of surface spin and enhanced magnetization.

Figure 3.12 shows room temperature magnetization-magnetic field (M-H) hysteresis

curves of BiFe1-xZrxO3 (x=0.03, 0.07, 0.10 and 0.15) nanoparticles. The strong ferromagnetic

(superparamagnetic) nature with very thin loops is observed in all the compositions with

saturation magnetization values of 7.42, 7.62, 7.07 and 6.35 emu/g for x=0.03, 0.07, 0.10 and

0.15 samples, respectively. These values are higher than the reported values of Zr4+ doped BFO

and La3+, Zr4+ codoped BFO bulk ceramics [139, 140].

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64

Figure 3.11: Magnetic hysteresis loop of BFO nanoparticles at room temperature, Inset shows the magnified view of

M-H loop.

Figure 3.12: (a) Magnetization hysteresis (M-H) loops for BiFe1-xZrxO3 nanoparticles, (b) its magnified view.

CHAPTER 3

65

These values are higher than the reported values of Zr4+ doped BFO and Pr3+, Zr4+

codoped BFO bulk ceramics [139, 140]. There are several reasons accountable for the

enhancement of saturation magnetization with increasing Zr4+ doping in BFO; (i) the particle size

of Zr4+ doped BFO nanoparticles are below 62 nm (smaller than the period length of cycloid spin

structure of BFO) and modification in cycloid spin structure gives rise to ferromagnetic ordering

instead of antiferromagnetic ordering, (ii) in case of nanoparticles, the surface to volume ratio is

large, giving rise to uncompensated spins at the surface and these uncompensated spins

contribute to high magnetization [141, 142]. Also, the substitution of Zr4+ ions on Fe-site initiates

the lattice distortion at FeO6 octahedra resulting in change in bond lengths and the Fe-O-Fe bond

angle. The Fe-O-Fe bond angle controls the superexchange interaction between the

antiferromagnetically aligned Fe3+ ions through oxygen anions [126].

The change in the Fe-O-Fe bond angle facilitates the dominance of ferromagnetic

interaction over antiferromagnetic interaction and results in the ferromagnetic nature in the

samples. Initially the saturation magnetization increases from x=0.03 to x=0.07 samples which is

ascribed to non magnetic character of Zr4+ ions and hence does not contribute directly to

magnetization but, subsequent replacement of few magnetically active Fe3+ ions with non

magnetic Zr4+ ions, allows ferromagnetic coupling rather than AFM coupling [139].

However, for x=0.10 and x=0.15 samples saturation magnetization slightly decreases and

this behaviour can be explained as: the increasing number of non magnetically active Zr4+ ions

plays dominant role rather than structural distortion induced suppression of cycloid spin

structure. With increasing Zr4+ ions, the formation of additional Fe-O-Zr co-ordination (which

becomes less magnetic sublattices) is responsible for reduced magnetization. In addition, the

evolution of triclinic phase with increasing Zr concentration in x=0.10 sample and phase

transformation to triclinic phase in x=0.15 sample may be accountable for this peculiar magnetic

behaviour.

The spin-phonon coupling as evident from Raman spectra (Figure 3.9 (b-e)) also predicts

this type of variation of magnetization with Zr4+ concentration. Therefore, this peculiar

magnetization behaviour of Zr4+ doped nanoparticles is well supported by Raman spectra

indicating spin phonon coupling.

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66

3.3.2.2 MAGNETIZATION VERSUS TEMPERATURE (M-T) STUDIES)

Figure: 3.13 (a) Zero-field-and field-cooled curves at 500 Oe for BiFe1-xZrxO3 nanoparticles (x = 0.03, 0.10), (b) An

enlarge view of ZFC curves.

Magnetization versus temperature curves for an external field of 500 Oe in zero-field-

cooled (ZFC) and field-cooled (FC) modes are displayed in Figure 3.13 (a) and (b) FC and ZFC

modes remain unperturbed until 293 and 284 K for x = 0.03 and 0.10 samples, respectively. After

which, the bifurcation of FC and ZFC curves gets initiated, i.e. FC mode increases and ZFC

mode decreases linearly with temperature. The large difference between FC and ZFC curves was

observed which was due to the spin glass-type behaviour. This type of behaviour was previously

reported in single crystals of BFO [143]. Spin glass-type behaviour is most dominant in the case

of nanoparticles which may be due to the following reasons: (i) reduction in particle size

increased surface to volume ratio which induces a large strain in the sample as compared to bulk

BFO and (ii) increased in a number of uncompensated Fe3+ spins on the surface results in

distortion in coordination, and magnetic spin-strain interaction increases [144].

These factors have a remarkable effect on magnetization of the sample. Figure 3.13 (b)

shows the zoom view of the ZFC curves of x = 0.03 and 0.10 nanoparticles. The inflection points

are obtained at ∼35 and ∼43 K for x = 0.03 and 0.10 samples, respectively, which are different

from the bulk single crystal. In bulk BFO, similar transition was reported at 50 K as a sharp cusp

[143]. Shifting of inflection point from 50 K reflects the nanosized nature of Zr-doped BFO

samples.

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3.3.3 ELECTRICAL STUDIES

3.3.3.1 DIELECTRIC STUDIES

Figure 3.14: Dielectric constant verses temperature curve for BiFe1-xZrxO3 samples at 1 MHz, 100 kHz and 75 kHz.

Figure 3.14 and 3.15 shows the thermal dependence of dielectric constant (ε′) and

dielectric loss (tan δ) at different frequencies for Zr4+ doped BFO samples. Both ε′ and tanδ

decreases with increasing Zr4+ concentration in BFO lattice. The dielectric anomaly in ε′ and tanδ

has been observed in all the compositions nearly at 300°C which may be related with the

antiferromagnetic Neel temperature (TN). This type of anomaly has been predicted by Landau-

Devonshire theory of phase transition in magnetoelectrically ordered systems as an influence of

vanishing magnetic order on the electric order [35, 145]. This shows the linear magnetoelectric

effect in all the compositions. Both ε′ and tanδ increase with temperature because at low

temperature the electric dipoles cannot align themselves in the direction of applied electric field

and results in low value of ε′. As the temperature increases, the mobility of charge carrier

increases giving rise to hopping conduction mechanism and dipoles acquiring sufficient energy to

follow the applied alternating field. This enhancement in polarization mechanism gives rise to

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higher values of ε′ and tan δ. The value of tanδ reduces with increasing Zr4+ ions concentration

which indicates that the formation of oxygen vacancies in BFO is suppressed by Zr4+ substitution

and thereby improved the resistivity of the materials. The Zr4+ ions act as a donor in BFO and the

substitution of Zr4+ ions requires charge compensation which may be achieved with the filling of

oxygen vacancies and results in improved dielectric properties.

Figure 3.15: Dielectric loss verses temperature curve for BiFe1-xZrxO3 samples at 1 MHz, 100 kHz and 75 kHz.

Figure 3.16 shows the variation of ε′ and tanδ of BiFe1-xZrxO3 samples with frequency

carried out at room temperature. It is clear from the graphs that both ε′ and tanδ decrease with

increasing frequency and remain constant at higher frequencies. High value of the dielectric

constant at the low frequencies and low value of the dielectric constant at high frequencies

indicate dispersion due to the Maxwell-Wagner type of interfacial polarization, which is in

agreement with the Koop’s phenomenological theory [146, 147]. This is because of the inability

of the dipoles to follow the alternating field at high frequencies. This phenomenon is related with

the space charge relaxation and these space charges originated from the oxygen vacancies which

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mainly come from Bi volatility and transitions from Fe3+-Fe2+. At low frequencies these space

charges follow the applied alternating field and as the frequency increases beyond a certain level

these space charges do not follow the field polarity [148]. Since the assembly of the charge

carriers in the dielectric structure requires finite time to line up their axis to the alternating

electric field, the dielectric constant value decreases at the high frequency [149].

Figure 3.16: Frequency dependence of dielectric constant and dielectric loss for BiFe1-xZrxO3 samples at RT.

3.3.3.2 IMPEDANCE STUDIES

To study the conduction mechanism in Zr4+ doped BFO samples, detailed impedance

analysis has been carried out. Figure 3.17 shows the variations of the real part of impedance (Z′)

as a function of frequency at different temperatures (100°C to 220°C). The value of Z′ decreases

with increasing frequency as well as with temperature which indicates the increase in

conductivity of the sample with temperature and frequency Negative Coefficient Temperature

Resistance (NTCR) behavior. At higher frequencies the values of Z′ coincide for all the

temperature which is due to the release of space charges as a result of lowering of barrier

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properties at high temperature and high frequency region [150]. The magnitude of Z′ increases

with increasing Zr4+ concentration which indicates the decrease in conductivity of the samples.

Figure 3.17: Variation of Z' as a function of frequency for BiFe1-xZrxO3 samples at different temperatures.

Figure 3.18 depicts the variation of Z″ with frequency at various temperatures for

different Zr4+ compositions. The magnitude of Z″max decreases with temperature indicating the

increase in resistive nature of the samples. With increase in temperature, the peak position of Z″

shifts towards higher frequency side because of the presence of electrical relaxation in the

sample. This shifting of peaks towards higher frequency shows that the relaxation time decreases

[151]. Asymmetric peak broadening suggests distribution of relaxation time and presence of non

Debye relaxation in the samples [152]. Merging of spectra at higher frequency indicates the

presence of electrical relaxation in the samples; at low temperature immobile species (Fe3+/Fe2+)

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contribute to relaxation process while defects/vacancies are responsible for the relaxation at

higher temperatures [153]

Figure 3.18: Variation of Z" as a function of frequency for BiFe1-xZrxO3 samples at different temperatures

The electric modulus could distinguish between electrode polarization, grain boundary

conduction and grain conduction which further allow us to understand different relaxation

mechanism. Figure 3.19 shows the variation of M′ with frequency at selected temperatures for all

the samples. The value of M′ increases with increasing frequency and possesses very low value at

lower frequencies confirming very small contribution of electrode effect [154]. At high frequency

regions there is a tendency of M′ to saturate at all the temperatures for all the samples. This

behaviour attributes the contribution of short range mobility of the charge carriers to the

conduction. Frequency dependent M′′ graphs are also plotted at selected temperatures in the

frequency range (0.01 Hz-1 MHz) as shown in Figure 3.20.

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Figure 3.19 Variation of M' as a function of frequency for BiFe1-xZrxO3 samples at different temperatures

Figure 3.20 Variation of M" as a function of frequency for BiFe1-xZrxO3 samples at different temperatures

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73

It can be seen that M′′max shifted towards higher frequency with increasing temperature,

which suggests the hopping mechanism of the charge carriers dominates intrinsically and

decreases relaxation time with increasing temperature. Asymmetric broadening of peaks indicates

the spreading of relaxation time with different time constants and hence the relaxation in the

material is non-Debye type.

Figures 3.21 shows Nyquist plots of complex impedance in the wide frequency range at

different temperatures (100-2200C). Two poorly resolved semicircular arcs are observed which

indicate different relaxation processes in the sample. Semi circle at lower frequencies represents

grain boundary relaxation process (Rgb) and at higher frequencies gives the information on grain

relaxation (Rb). It has been clearly understood from these plots that both the grains and grain

boundaries are responsible for the conduction mechanism in the samples. Figure 3.22 shows

Nyquist plots of complex modulus with different temperatures, the emergence of a semicircle in

the electric modulus spectrum is attributed to the relaxation phenomena. The depressed

semicircles in the complex modulus spectra refer to the deviation from the ideal Debye behavior

and indicate multiple relaxations occurring within the samples [155].

Figure 3.21: Variation of Z' with Z'' of BiFe1-xZrxO3 samples at different temperatures

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Figure 3.22: Variation of M' with M'' of BiFe1-xZrxO3 samples at different temperatures

3.3.3.3 CONDUCTIVITY STUDIES

The frequency and temperature dependent conductivity has been measured to study

conduction mechanism in Zr4+ doped BFO samples. The conductivity behaviour can be expressed

by Jonscher’s universal power law [156]:

)5.3()()0()( 01n

TAωσωσσωσ +=+=

Where, )0(σ is frequency independent conductivity and )(1 ωσ is frequency dependent

term for the ac conductivity. Here, A determines the strength of polarizability and the exponent

n can have any value between zero and one which represents the degree of interaction between

mobile ions and lattice around them. Both A and n are temperature and material dependent.

Figure 3.23 illustrates the variation of ac conductivity with frequency at different temperatures.

All the plots show dispersion in σac which shifts towards higher frequency with increasing

temperature. It can be seen from Figure 3.23 that σac increases with temperature which may be

due to increase of hopping of charge carriers with increasing frequency. Curve at low frequency

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75

region tends to flattened with increasing temperature; this suggests the dc conduction behaviour

of the samples [157]. Increase in conductivity with temperature may be attributed to the increase

in oxygen vacancy mobility [158]. The extrapolation of frequency independent plateau gives the

value of dc conductivity. The value of ac conductivity increases with increasing frequency which

may be due to the disordering of cations between neighboring sites and the presence of space

charges. From the nature of the curves it is clear that curves approaches each other and exhibit

low frequency dispersion phenomenon obeying Jonscher’s power law, according to which the

origin of frequency dependent conductivity lies in the relaxation mechanism arising due to

mobile charge carriers [159]. For x = 0.03 sample, step like behaviour was observed which

decreases with increasing Zr4+ concentration, implying increased in stability in A and n

parameters against temperature with increasing Zr4+ content.

Figure 3.23 Variation of σac with frequency for BiFe1-xZrxO3 samples at selected temperatures

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The temperature dependence of ac conductivity at frequency 10 kHz is shown in the

Figure 3.24. The value of the ac conductivity is relatively smaller at lower temperature as

compared to higher temperature; however, it shows different values in the various temperature

regimes. The electrical conductivity of a material is a thermally activated behavior and it obeys

Arrhenius relation given by: σ=σoe(Ea/K

BT), where Ea is activation energy, kB is the Boltzmann

factor and σo is the pre-exponential factor. The variation of activation energy from 0.076-0.066

eV in the temperature range 313-339 K suggests that the space charges play an important role in

the conduction mechanism and the conduction is purely intrinsic (due to the space charge region

which is created at the interfaces). However, as the temperature increases, the movement of

oxygen ions gives rise to the activation energy ranges 0.86-0.82 eV in the temperature range 689-

763 K. The conductivity of the samples decreases with increasing the doping concentration which

is attributed to the reduction in oxygen vacancies.

Figure 3.24: Variation of σac with temperature for BiFe1-xZrxO3 samples at 10 kHz frequency

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During the firing process of the samples electron release is captured by Fe3+ ion which

gets converts to Fe2+ ion and hence hopping of electron occurs between the two different valance

states and thus the conductivity of the grains increase. When Fe3+ ions are replaced by the Zr4+

ions, tendency of the conversion of Fe3+ ions to Fe2+ ions reduces which in turn hinder the

hopping of electron and conductivity of the grains decreases.

3.3.4 OPTICAL STUDIES

3.3.4.1 UV-VIS DIFFUSE ABSORPTION STUDIES

.

Figure 3.25: (a) UV-Vis diffuse absorption spectrum and (b-e) (αhν)2 versus Energy (hν) plots for BiFe1-

xZrxO3 samples.

(a) (b)

(c) (d)

(e) (f)

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78

In recent years optical properties of BFO has attracted much attention due to its potential

in photocatalytic activity and other applications. Therefore, detailed optical properties of pure and

Zr4+ doped BFO nanoparticles have been studied by UV-visible diffuse reflectance spectroscopy.

UV-Visible diffuse absorption spectra for pure and Zr+4 doped BFO samples in the range 1.5-5

eV is displayed in Figure 3.25(a).

However, the 3.2 eV band splits into multiplets followed by a low-energy band at 2.5 eV

for x = 0.03 to 0.15 samples. The splitting of this strongly hybridized O p and Fe d excitations

may be attributed to the distorted FeO6 octahedra and strained nature of Zr-doped BFO

nanoparticles [160]. The two electronic bands present at ~3.2 and 4.3 eV are assigned to minority

channel dipole-allowed charge transfer excitations, however, the 3.2 eV band splits into

multiplets followed by a low energy band at 2.5 eV [161]. The splitting of this strongly

hybridized O p and Fe d excitations may be attributed to the distorted FeO6 octahedra and

strained nature of Zr-doped BFO nanoparticles. Band below at ~2.1 eV assigned to on-site d-d

crystal-field excitations of Fe3+ ion in the all the BFO samples, this transition is spin forbidden

[162].

The energy band gap of the nanoparticles were calculated by using classical Tauc′s

approach i.e. plotting (αhν)2 versus hν and αhν = A(hν-Eg)n where hν is the photon energy, A is

the parameter that related to the effective masses associated with the valence and conduction

bands constant, Eg is energy band gap, α is absorption coefficient. The value of n is depends upon

the nature of electronic transition responsible for absorption and it can have any value ½, 3/2, 2 or

3. For direct band gap value of n taken is ½. The extrapolation of the straight portion (αhν)2

against hν plot to the point α=0 on x-axis provides the direct energy band gap [162]. The band

gap of the samples varies from 2.27-2.17eV (Fig. 3.25(b-f)), which is consistent with the

previously reported band gap of BFO nanoparticles [163-165]. The variation in band gap is also

related with the particle size but in our case the particle size of all the samples is greater than 20

nm which is large enough not to exhibit quantum confinement effect; hence the variation in band

gap may be attributed to the doping ion which alter the electronic structure of pristine BFO.

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3.4 CONCLUSIONS

In summary, pure and Zr4+ doped BFO nanoparticles were successfully prepared by sol-

gel method and synthesis conditions were optimized. Tartaric acid based sol-gel route was

effective in getting pure phase samples calcined at 550ºC. The prepared samples were structurally

characterized by XRD patterns and Raman spectra which exhibited pure phase formation with

rhombohedral structure of space group R3c for x = 0.0-0.10 samples and triclinic phase with

space group P1 for x = 0.15 sample. The distortion in Zr4+ doped BFO nanoparticles is strongly

indicated by prominent Raman modes suggesting the transformation of rhombohedral structure to

triclinic structure for x = 0.15 sample and exhibiting the spin-phonon coupling. The particle size

in ~ 15-50 nm range and structural distortion derived the magnetic behavior of the samples

confirming the spin-phonon coupling. Initially the saturation magnetization increases from

x=0.03 to x=0.07, however, for x=0.10 and x=0.15 samples saturation magnetization slightly

decreases. The reduction in saturation magnetization for x = 0.10 and 0.15 samples is ascribed to

the dominance of increasing non-magnetically active Zr4+ ions over the suppressed spatial spin

modulation. Substitution of Zr4+ ions leads to improved dielectric properties of nanoparticles. It

has been clearly understood from Nyquist plots that both the grains and grain boundaries are

responsible for the conduction mechanism in the samples. The conductivity of the samples

decreases with increasing Zr4+ concentration which is attributed to the reduction in oxygen

vacancies. During the firing process of the samples electron release is captured by Fe3+ ion which

gets converts to Fe2+ ion and hence hopping of electron occurs between the two different valance

states and thus the conductivity of the grains increase. When Fe3+ ions are replaced by the Zr4+

ions, tendency of the conversion of Fe3+ ions to Fe2+ ions reduces which in turn hinder the

hopping of electron and conductivity of the grains decreases. The red shift in the optical band gap

of Zr4+ doped BFO nanoparticles is observed due to distortion in the BFO lattice. The formation

of the perovskite structure can be confirmed by the formation of metal-oxygen bonds in the

samples calcined at 450ºC and 550ºC. The optical band gap in the visible region from 2.27 to

2.17 eV along with improved ferromagnetic nature of the nanoparticles makes them suitable for

potential device applications. .