Fabrication and Characterization of Chemically Modified Multiferroic Bismuth Ferrite Thin Films

FABRICATION AND CHARACTERIZATION OF CHEMICALLY MODIFIED MULTIFERROIC BISMUTH FERRITE THIN FILMS

YAN FENG (M. Sc)

A THESIS SUBMITED FOR THE DEGREE OF DOCTOR OF PHILOSOPHY DEPARTMENT OF MECHANICAL ENGINEERING NATIONAL UNIVERSITY OF SINGAPORE 2011

Acknowledgements

I would like to express my sincere gratitude to my supervisors, Prof. Lu Li, and Prof.

Lai Man On from Department of Mechanical Engineering, National University of

Singapore (NUS), and Prof. Zhu Tiejun from Department of Materials Science and

Engineering, Zhejiang University (ZJU) for giving me the opportunity to work on this exciting project. I would especially like to thank all of them for their guidance and support throughout my PhD study at NUS. I benefited from their guidance in every aspect during my Ph.D research, such as the discussions we held and intellectual suggestions they made regarding my work.

I am grateful for the insights and advice Prof. Zeng Kaiyang shared with me. His suggestions and guidance were a great help in performing piezoresponse force microscopy (PFM) and Kelvin probe force microscopy (KPFM) on the microscopic ferroelectric properties as related to the experiments in Chapter 6.

I would like to thank Prof. Rüdiger-A. Eichel from Institut für Physikalische Chemie,

Albert-Ludwigs-Universität Freiburg (German) for his kind assistance and advice on the understanding of the magnetic phase transition and magnetoelectric coupling effect via electron paramagnetic response (EPR). It is a great honor for me to have opportunity to work in his group for four months. Also, I would like to thank Dr. Emre

Erdem and Dr. Peter Jakes in this group for their kind advice and suggestions.

I would like to thank Prof. I. M.Reaney from Department of Engineering Materials,

University of Sheffield for his kind assistance the understanding of the microstructure of the investigated thin films via HRTEM.

I would also like to thank Dr. Zhang Zhen as a collaborator. All the ideas we shared and discussed have proven very useful. This project would not have gone so well without his contribution at the forefront of this research.

I specially thank Dr. Wang Shijie for his help and advice at the beginning of my studies. I would like to express my gratitude for the help of my colleagues and collaborators: Dr. Xiahui, Mr. Wang Hailong, Mr. Ye Shukai, Mr. Xiao Pengfei, Miss.

Zhu Jing, Mr. Song Bohang and other members in Prof. Lu’s research group.

In addition, I would like to give my special thanks to the staff of Materials Science

Laboratory, Department of Mechanical Engineering, and National University of

Singapore.

Finally, I want to thank my family for their understanding, encouragement, and endless love throughout my life.

Table of Contents

Acknowledgements ...... I

Table of Contents ...... III

Abstract ...... VII

List of Tables ...... VIII

List of Figures ...... IX

List of Symbols and Abbreviations ...... XVII

List of Publications ...... XVIII

Chapter 1 Introduction ...... 1

1.1 Overview & Motivations...... 1

1.2 Outline ...... 3

Chapter 2 Literature Review ...... 4

2.1 Magnetoelectric effect and multiferroic materials ...... 4

2.1.1 Magnetoelectric effect[14] ...... 4

2.1.2 Multiferroics ...... 6

2.1.3 Single phase multiferroics [3] ...... 8

2.1.4 Multiferroic composites ...... 10

2.2 Multiferroic BiFeO3 ...... 10

2.2.1 Structure ...... 10

2.2.2 Ferroelectricity...... 12

2.2.3 Dielectric Properties ...... 13

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2.2.4 Magnetism ...... 14

2.2.5 Magnetoelectric Coupling ...... 15

2.2.6 Domains and domain walls[46] ...... 16

2.3 Device application ...... 18

2.3.1 Data storage [1] ...... 18

2.3.2 Optoelectronic devices ...... 19

2.4 Thin film deposition and characterization ...... 21

2.4.1 Pulsed Laser Deposition (PLD) ...... 21

2.4.2 X-Ray diffraction (XRD)...... 22

2.4.3 Atomic force microscopy (AFM) ...... 22

2.4.4 Macroscopic Ferroelectric measurement ...... 23

2.4.5 Dielectric measurement ...... 24

2.4.6 Piezoresponse force microscopy (PFM) ...... 25

2.4.7 Switching Spectroscopy PFM (SS-PFM) ...... 26

2.4.8 Kelvin probe force microscopy (KPFM) ...... 27

2.4.9 Vibration sample magnetometer (VSM) ...... 28

2.4.10 Magnetic force microscopy (MFM) ...... 28

Chapter 3 Experimental Procedures...... 30

3.1 Fabrication of Targets ...... 30

3.1.1 Pure BiFeO3 target ...... 30

3.1.2 Chemically Modified BiFeO3 targets ...... 30

3.2 Thin film growth ...... 31

3.2.1. Substrate and target cleaning ...... 31

3.2.2 Film deposition ...... 31

3.3 Thin film characterization ...... 32

Chapter 4 Pure BFO thin films ...... 34

4.1 Introduction ...... 34

4.2 Epitaxial BiFeO3 thin films ...... 34

4.2.1 Structures of (001), (011) and (111) oriented BiFeO3 thin films ...... 34

4.2.2 Surface morphology ...... 37

4.2.3 Macroscopic electrical properties of (001), (011) and (111) oriented BiFeO3

thin films ...... 37

4.2.4 Ferroelectric domain structure using PFM ...... 40

4.2.5 Magnetoelectric coupling of BiFeO3 thin films via PFM and MFM ...... 46

4.3 Polycrystalline BiFeO3 thin films ...... 48

4.3.1 BFO on Pt/TiO2/SiO2/Si substrate ...... 48

4.3.2 Size effect on the piezoelectric response of BFO on Pt/TiO2/SiO2/Si

substrate ...... 56

4.4 Effects of bottom electrode on switching behavior at nanoscale ...... 61

4.4.1 Structure and multiferroic properties of BFO on LaNiO3/Si and SrRuO3/Si

substrates ...... 61

4.5 Effect of bottom electrode on the surface potential of polycrystalline BFO ...... 70

Chapter 5 Chemically modified BFO thin films ...... 74

5.1 Introduction ...... 74

5.2 La doped BFO ...... 76

5.3 Ru doped BFO ...... 84

5.4 La, Ru codoped BiFeO3 ...... 93

5.5 Pb(Zr0.52Ti0.48)O3 (PZT) codoped BFO ...... 101

5.6 Conclusions ...... 110

Chapter 6 Leakage mechanisms of BFO ...... 113

6.1 Introduction ...... 113

6.2 Deposition temperature dependent leakage mechanism ...... 115

6.3 Oxygen pressure dependent leakage mechanism ...... 119

6.4 Leakage mechanism in La and Ru codoped BFO ...... 123

6.5 Conclusions ...... 130

Chapter 7 EPR study of BFO ...... 132

7.1 Introduction ...... 132

7.2 EPR study of pure BFO ...... 134

Chapter 8 Conclusions and Future work ...... 143

8.1 Conclusions ...... 143

8.2 Future Work ...... 145

Reference ...... 146

Abstract

Magnetoelectric multiferroics exhibit coexisting magnetic and ferroelectric phases, with coupling between magnetic and electric ordering. In this work, pulsed laser deposition (PLD) technology has been used to deposit multiferroic pure and chemically modified BiFeO3 (BFO) thin film on different substrates. The novelty of our work is to present a systematically of chemically modified BFO thin films combining with the macroscopic and microscopic ferroelectric properties and local domain switching behavior. The microstructures, electrical and magnetic properties of the as-grown films are systematically investigated.

A, B-site and A, B codoped effects have been determined, suggesting that the dopants could greatly impact the multiferroic properties of BFO films. In addition, the leakage mechanism of the multiferroic films have been studied at different growth conditions and different measurement conditions, such as temperature and electric field.

Furthermore, defect chemistry and low temperature magnetic phase transition have been evaluated via electron paramagnetic resonance (EPR).

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List of Tables

Table 3.1 Deposition parameter for thin films…………………………………...32

eff Table 4.1 The orientation dependence of d33, Ps, and effective coefficient, Q33 .43

Table 7.2 Spin counting for the samples annealed in different atmospheres…...136

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List of Figures

Figure 2.1 Phase control in ferroic and multiferroics………………………………6

Figure 2.2 The relationship between multiferroic and magnetoelectric…………....8

Figure 2.3 Schematic of the crystal structure of BFO and the ferroelectric polarization (arrow) and antiferromagnetic plane (shaded planes)…...12

Figure 2.4 Polarization of BiFeO3: (a) bulk single crystal and (b) epitaxial thin film…………………………………………………………………….13 Figure 2.5 Schematics of the 64 nm antiferromagnetic circular cycloid. The canted antiferromagnetic spins (blue and green arrows) give rise to a net magnetic moment (purple arrows) that is specially averaged out to zero due to the cycloidal rotation. The spins are contained within the plane defined by the polarization vector…………………………………….14 Figure 2.6 Schematics of the planes of spin rotations and cycloids k~1 vector for the two polarization domains separated by a domain wall (in light gray). Rotating the polarization by 71o results in a change of the magnetic easy plane, meaning that sublattice magnetization can be switched by an applied voltage………………………………………………………...16 Figure 2.7 Schematic of the three types of ferroelectric domain walls. The ferroelectric polarization (bold arrows) and antiferromagnetic plane (shaded planes)………………………………………………………..17 Figure 2.8 MERAMs based on exchange-bias coupling between a multiferroic that is ferroelectric and antiferromagnetic (FE-AFM, green layer), and a thin ferromagnetic electrode (blue layer). A tunneling barrier layer between the two top ferromagnetic layers provides the two resistive states……19 Figure 2.9 The variation of photocurrent with sample rotation under illumination with a linearly polarized light. The experimental sketch is shown in the inset. The PV effect becomes maximum when the polarized-light electric field is parallel to the in-plane component of the ferroelectric polarization and minimum when the field is perpendicular to the in- plane component...………………………………………. ………….20 Figure 2.10 Schematic of the pulsed laser deposition system…… ………………..21

Figure 2.11 Schematic of the experimental set-up for ferroelectric hysteresis loop measurements………………………………………………………….23 Figure 2.12 Schematic structure of the electrical measurement……………………24

Figure 2.13 Schematic of PFM operation…………………………………….……25

Figure 2.14 Switching spectroscopy PFM diagram………………………………..27

Figure 4.1 X-ray diffraction spectra for epitaxial BFO films deposited on SRO buffered STO substrates with different orientations…………………35 Figure 4.2 Pole figure plots of the corresponding to BFO (110) diffraction for (a) (001) (b) (011) (c) (111)-oriented thin films, indicating a cube-on-cube growth behavior……………………………………………………….35 Figure 4.3 Room temperature Raman spectra of the (100), (110), and (111)- oriented rhombohedral BFO thin films and of a (111) STO substrate..36 Figure 4.4 Atomic force microscopy (AFM) images of BFO thin films deposited at (a) (001), (b) (011), (c) (111)-oriented substrate……………………...37 Figure 4.5 Ferroelectric hysteresis (P-E) loops of (001), (011), (111)-oriented BFO thin films measured at room temperature……………………………..38 Figure 4.6 Leakage current density for (001), (011), (111)-oriented BFO thin films measured at room temperature……………………………………… 39 Figure 4.7 Leakage current and time relationship for (001), (011), (111)-oriented BFO thin films measured at room temperature……………………….40 Figure 4.8 (a), (d), and (g) AFM images (1×1 µm 2), (b), (e) and (h) out-of-plane piezoresponse (OP-PFM) amplitude images, and (c), (f) and (i) OP- PFM phase images of (001), (011), (111)-oriented BFO thin films. Yellow (bright) and purple (dark) on the piezoresponse image correspond to negative (upward) and positive (downward) domains, respectively……………………………………………………………41 Figure 4.9 Calculated average piezoresponse of PFM images obtained at each voltage step vs ac voltage for (001), (011), (111)-oriented BFO thin films. 95% confidence error bars are smaller than 3% and too negligible to be visible on the graph…………………………………… ……… 42 Figure 4.10 Out of plane PFM images demonstrating switching of (001)-oriented films: (a) amplitude and (b) phase images before DC bias and (c) amplitude and (d) phase images after -10 and +10 V DC voltages….. 44 Figure 4.11 (a) AFM images (1×1 µm 2), (b) out-of-plane piezoresponse (OP-PFM) amplitude images, and (c) OP-PFM phase images of (001)-oriented BFO thin films with the tip position was donated as A; (d) and (f) PFM phase voltage hysteresis loop and (e) and (g) amplitude voltage butterfly loop of (001)-oriented BFO thin films, in both “ON” (f, g) and “OFF” (d, e) states…………………………………………….………46

Figure 4.12 (a) Topography (20 × 20 m2) (b) Amplitude and (c) magnetic domain imaging of (001)-oriented BFO thin film using magnetic force microscopy (MFM)……………………………………………………47 Figure 4.13 Color Multimode scanning probe images of (001) BFO thin film. (a), (d) Topography and magnetic domain imaging of BFO thin film using

magnetic force microscopy (MFM). (b), (e) Topography and piezoresponse imaging of BFO thin film using piezoresponse force microscopy (PFM). The image shows the ferroelectric domain switching performed by the poling DC bias. (c), (f) Topography and domain imaging of BFO thin film using magnetic force microscopy after writing ferroelectric domains (e). This image gives information of both ferromagnetic and ferroelectric domain structures…………...………..48 Figure 4.14 X-ray diffraction (XRD) spectra of the BFO deposited on Pt substrate at 550 oC 50 mTorr……………………………………………………….49 Figure 4.15 (a) AFM images (1×1 µm 2), (b) out-of-plane piezoresponse (OP-PFM) amplitude images, and (c) OP-PFM phase images of BFO thin films deposited on Pt/TiO2/SiO2/Si substrate………………………………50 Figure 4.16 (a) Ferroelectric hysteresis loops and (b) leakage current density for BFO thin films deposited on Pt/TiO2/SiO2/Si substrate measured at room temperature…………………………………………………………….50 Figure 4.17 Dielectric properties of pure BFO thin films measured as a function of frequency at room temperature………………………………………51

Figure 4.18 Variation in dielectric constant (r) and loss tangent (tan of pure BFO thin films deposited on Pt/TiO2/SiO2/Si substrate measured as a function of temperature………………………………………………………..52 Figure 4.19 Magnetic hysteresis loops for polycrystalline BFO thin films deposited on Pt/TiO2/SiO2/Si measured at room temperature. …………………53 Figure 4.20 Calculated average piezoresponse of PFM images obtained at each voltage step vs. ac voltage for the BFO thin films deposited on the Pt/TiO2/SiO2/Si substrate. ……………………………………………..54 Figure 4.21 (a) AFM images (1×1 µm 2), (b) out-of-plane piezoresponse (OP-PFM) amplitude images, and (c) OP-PFM phase images of polycrystalline BFO thin films with the tip position was donated as “A”; (d) and (f) PFM phase voltage hysteresis loop and (e) and (g) amplitude voltage butterfly loop of polycrystalline BFO thin films deposited on Pt/TiO2/SiO2/Si, in both “ON” (f, g) and “OFF” (d, e) states………….55 Figure 4.22 (a) Topography (20 × 20 m2) (b) Amplitude and (c) magnetic domain imaging of polycrystalline BFO thin film deposited on Pt/TiO2/SiO2/Si using MFM……………………………………………………………..56

Figure 4.23 XRD spectra for the BFO films of varying thickness on Pt/TiO2/SiO2/Si substrate. ………………………………………………………………57 Figure 4.24 Out of plane lattice parameter as a function of thickness of BFO thin films on Pt/TiO2/SiO2/Si substrate. ……………………………………58 Figure 4.25 (a), (d), and (g) AFM images (2×2 µm 2), (b), (e) and (h) out-of-plane piezoresponse (OP-PFM) amplitude images, and (c), (f) and (i) OP- PFM phase images for 10, 80, 160 nm BFO thin films on

Pt/TiO2/SiO2/Si substrate, respectively. Yellow (bright) and purple (dark) on the piezoresponse image correspond to negative (upward) and positive (downward) domains, respectively. …………………………..59

Figure 4.26 d33 as a function of the thickness of BFO film deposition on Pt/TiO2/SiO2/Si substrate. ………………………….………………….60

Figure 4.27 XRD spectra for polycrystalline BFO thin films deposited on (a) LaNiO3 (LNO) and (b) SrRuO3 (SRO)-coated Si substrate……………………62 Figure 4.28 (a) and (d) AFM images (1×1 µm 2), (b) and (e) out-of-plane piezoresponse (OP-PFM) amplitude images, and (c) and (f) OP-PFM phase images for BFO thin films on LNO and SRO-coated Si substrate, respectively. Yellow (bright) and purple (dark) on the piezoresponse image correspond to negative (upward) and positive (downward) domains, respectively…………………………………………………..65 Figure 4.29 (a) ferroelectric hysteresis loops and (b) Leakage current density for BFO thin films deposited on LNO and SRO-coated Si substrate measured at room temperature…………………………………….…..67

Figure 4.30 (a) dielectric constant (r) and (b) tangent loss (tan dependence on frequency of pure BFO thin films deposited on LNO and SRO coated Si substrate measured at room temperature………………………..……..68 Figure 4.31 Calculated average piezoresponse of PFM images obtained at each voltage step vs ac voltage for the BFO thin films deposited on the LNO and SRO coated Si substrate. …………………….……………………69

Figure 4.32 (a) Topography (5×5 m2), (b) magnetic domain imaging of polycrystalline BFO thin film deposited on LNO/Si and SRO/Si substrate using magnetic force microscopy (MFM), respectively……..70 Figure 4.33 (a) and (d) AFM images, (b) and (e) out-of-plane PFM phase image and (c) and (f) KPFM surface potential distribution of the area scanned with the applied biases from -10 ~ +10V with a 2V step to the cantilever tips for BFO deposited on Pt/TiO2/SiO2/Si and LNO/Si substrate, respectively. …………………………………………………..……….71 Figure 4.34 Surface potential line profiles obtained from the KPFM images……..73

Figure. 5.1 (a) X-ray diffraction spectra for BFO and BLFO thin film, (b) A typical FE-SEM cross-sectional image of the films…………………………..77 Figure. 5.2 (a), (d) AFM images, (b), (e) out-of-plane PFM amplitude and (c), (f) phase images for BFO and BLFO thin film(1 × 1 µm2), respectively. The Height data is the topography in the image, and the phase data (piezoresponse) is mapped on as color. Bright and dark indicate the upward and downward domain orientation in Figure 5.2c and 5.2f….79 Figure.5.3 Leakage current density of the Pt/BFO/Pt and Pt/BLFO/Pt capacitors at room temperature…………………………………………………80

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Figure 5.4 Ferroelectric hysteresis loops of the Pt/BFO/Pt and Pt/BLFO/Pt capacitors at room temperature……………………………………….81 Figure 5.5 Electric field dependences of remanent polarization and coercive field of the Pt/BLFO/Pt capacitors at room temperature…………………..81 Figure 5.6 Dielectric properties of the Pt/BFO/Pt and Pt/BLFO/Pt capacitors at room temperature……………………………………………………..82 Figure 5.7 Magnetic properties of BFO and BLFO thin film measured at room temperature………………………………………………………….83

Figure 5.8 (a) X-ray diffraction spectra of the BFO and BFRO thin films with inset of magnified patterns showing a diffraction at 2θ = 31.76°; and (b) and (c) cross-sectional FESEM images of the BFO and BFRO, respectively………………………………………………………….85 Figure 5.9 (a) and (d) AFM images (2×2 µm2), (b) and (e) out-of-plane piezoresponse amplitude images, and (c) and (f) piezoresponse phase images of BFO and BFRO thin films. Yellow (bright) and purple (dark) on the piezoresponse image correspond to positive (upward) and negative (downward) domains, respectively. ………………………86 Figure 5.10 Leakage current density for BFO and BFRO thin films measured at room temperature…………………………………………………….87 Figure 5.11 Ferroelectric hysteresis loops for BFO and BFRO thin films measured at room temperature…………………………………………………88 Figure 5.12 Dielectric properties for BFO and BFRO thin films measured at room temperature……………………………………………………………89 Figure 5.13 Leakage current and time relationship for BFO and BFRO thin films measured at room temperature……………………………………..90 Figure 5.14 Magnetic hysteresis loops for BFO and BFRO thin films measured at room temperature……………………………………………………91 Figure 5.15 (a) XPS spectra of the Fe2p lines of the BFO and BFRO films, the insets are the peak fitting simulations. (b) and (c) XPS spectra of Ru 3d and 3p showing variable oxidation state for the BFRO film. ………93 Figure 5.16 XRD spectra of the BFO, BLFO, BFRO and BLFRO thin films with inset of magnified patterns showing a diffraction at 2θ = 32° (a) and 46° (b) ……………………………………………………………………94 Figure 5.17 Raman spectra with inset of magnified patterns showing composition dependent shift in the wave number between 200 to 240 cm-1 (a) and 550 to 700 cm-1(b) of the BFO, BLFO, BFRO and BLFRO thin films………………………………………………………………...96 Figure 5.18 (a) and (e) AFM images (2×2 µm2), (b) and (f) out-of-plane piezoresponse amplitude images, (c) and (g) piezoresponse phase images, and (d) and (h) magnetic domain images (20×20 µm 2) via

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MFM of BFO and BFRO thin films. Purple (dark) and yellow (bright) on the piezoresponse image correspond to positive (upward) and negative (downward) domains, respectively. ……………………….98 Figure 5.19 (a) XPS spectra of the Fe2p lines of the BFO and BLFRO films, the insets are the peak fitting simulations and (b) XPS spectra of the La3d5 lines and fitting simulation of BLFO and BLFRO thin film………..99 Figure 5.20 (a) Magnetic hysteresis loop (M-H), (b) Polarization-electric field hysteresis loops of BLFRO thin film as a functions of electric field recorded in the 100~350 K temperature range………………………100 Figure 5.21 (a) XRD spectra of the pure, PZT2% and PZT5%BFO thin films with inset of magnified patterns showing a diffraction at 2θ = 32.34°; (b) a typical cross-sectional FESEM images of the PZT 2mol% modified BFO film; (c) Raman spectra of the pure, PZT 2% and PZT 5%BFO thin films with inset of magnified patterns showing composition dependent shift in the wave number between 200 to 240 cm-1; (d) -1 indicates the peak intensity ratio I620/I154 and 220 cm line peak position versus PZT mol percent. The solid lines are guides for eyes….……103 Figure 5.22 (a), (b), and (c) AFM images (2×2 µm2), (d), (e) and (f) out-of-plane piezoresponse (OP-PFM) amplitude images, and (g), (h) and (i) OP- PFM phase images of pure, PZT2% and PZT5%BFO thin films. Yellow (bright) and purple (dark) on the piezoresponse image correspond to negative (downward) and positive (upward) domains, respectively………………………………………………………..105 Figure 5.23 Magnetic hysteresis loop (M-H) measured at room temperature for the pure, PZT2% and PZT5%BFO films. ………………………………106 Figure 5.24 (a) Polarization-electric field hysteresis loops, (b) Leakage current density, (c) dielectric properties of pure, and (d) leakage current vs. time relationship of the pure and PZT2% and PZT5%BFO thin films measured at room temperature……………………………………..108 Figure 6.1 Typical J – E characteristics of Pt/BFO/LNO capacitors for different deposition temperature at a fixed 50 mTorr oxygen pressure………116 Figure 6.2 Log (J) vs log (E) plots at positive bias for the SCLC mechanism….117

Figure 6.3 Log (J) vsE0.5 plots at positive bias for the Schottky mechanism…117

Figure 6.4 Log (J/E) vsE0.5 plots at positive bias for the PF mechanism……...118

Figure 6.5 Log (J/E2) vs1/E plots at positive bias for the FN mechanism………119

Figure 6.6 Leakage current density-electric field characteristics of Pt/BFO/LNO capacitor films deposited at different oxygen pressure at 500 oC……120

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Figure 6.7 Log (J) vs. log (E) for the BFO films deposited at different oxygen pressure……………………………………………………………..121 Figure 6.8 Log (J) vsE0.5 plots at positive bias for the Schottky mechanism for the films deposited as a function of oxygen pressure. …………………..121 Figure 6.9 Log (J/E) vsE0.5 plots at positive bias for the PF mechanism………..122

Figure 6.10 Log (J/E2) vs1/E plots at positive bias for the FN mechanism at different oxygen pressure. ……………………………………...……123 Figure 6.11 (a) X-ray diffraction spectra for the BLFRO thin film on a (001) Nb- STO substrate. (b) Pole figure image of the BLFRO thin film on (011) plane. (c) AFM image and (d) out-of-plane piezoresponse phase image of BLFRO thin film. …………………………………………….…..124 Figure 6.12 (a) Polarization-electric field hysteresis loops at room temperature, (b) Leakage current density as a function of temperature, (c) dielectric properties as a function of frequency and (d) Temperature dependence of the dielectric constant measured at different frequency of BLFRO thin films……………………………………………..………………125 Figure 6.13 (a) J-E curves of the BLFRO film at low field, fitting according to the modified Langmuir-Child law; the inset presents the fitting parameters  and  as function of temperature. (b) Log (J/E) vs E1/2 curves of BLFRO film at high fields at various temperatures; the inset shows the trap ionization energy as a function of E1/2. (c) Log(J) vs Log(E) plots at positive bias and temperature of 100, 140 and 170K; (d)ln(J/E2) vs (1/E) at positive bias and temperature of 200~350K, the inset shows the apparent potential barrier for BLFRO film………………………….128 Figure 6.14 V1/2 dependence of the apparent potential barrier for BFLRO film deposited on Nb-STO substrate……………………………………..130

Figure 7.1 (a) X-ray diffraction spectra of BiFeO3 polycrystalline powder. (b) Normalized X-band EPR spectra of BiFeO3 polycrystalline powders measured under different microwave energy at room temperature..135 Figure 7.2 X-band EPR spectra for BFO powders treated under different annealing atmospheres recorded at RT……………………………………….136 Figure 7.3 Temperature dependence of the X-band EPR spectra of polycrystalline BFO powder………………………………………………………..137

Figure 7.4 Temperature dependence of the resonance field, BC , g-factor, and the

intensity of EPR, IEPR indicate that magnetic phase transitions near 140, 200, and 260 K for BFO powder. The solid lines are guides for eyes……………………………………………………………….139

Figure 7.5 (a) Temperature dependence of the peak-to-peak EPR linewidth BPP

and (b) ln(BTPP ) as a function of 1000/T. The solid lines are guided for eyes…………………………………………………………….141 Figure 7.6 Variation of A/B ratio with temperature in BFO. The inset illustrates the method adopted to calculate the A/B ratio………………………….142

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List of Symbols and Abbreviations

ζ: electrical conductivity

A: electrode area

C: Capacitance

: Dielectric permittivity

–12 2 –1 –2 : Dielectric permittivity of free space, ε0 = 8.854 × 10 C N m

r: Dielectric constant (no unit) tanδ : Dielectric loss (no unit)

T : Temperature

P : Polarization (unit: μCcm–2) d : Out-of-plane lattice parameter (unit: Å) dfilm : Out-of-plane lattice parameter of thin film (unit: Å) dbulk : Out-of-plane lattice parameter of bulk material (unit: Å)

FE-SEM : Abbreviation of Field-Emission Scanning Electron Microscope”

LNO : Abbreviation of LaNiO3

BFO : Abbreviation of BiFeO3

PLD: Pulsed Laser Deposition

XRD: X-Ray Diffraction

AFM: Atomic force microscopy

PFM: Piezoresponse force microscopy

SS-PFM: Switching Spectroscopy PFM

KPFM: Kelvin probe force microscopy

VSM: Vibration sample magnetometer

MFM: Magnetic force microscopy

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List of Publications

1. F. Yan, M. O Lai, L. Lu and T. J. Zhu. Role of Pb(Zr0.52Ti0.48)O3 substitution in

multiferroic properties of polycrystalline BiFeO3 thin films. J. Appl. Phys., 110,

114116, (2011).

2. F. Yan, T. J. Zhu, M. O Lai, and L. Lu. Influence of La and Ru dopants on

multiferroic properties of polycrystalline BiFeO3 thin films, Appl. Phys. Express.

4, 111502 (2011)

3. F. Yan, T. J. Zhu, M. O Lai, and L. Lu. Effect of bottom electrodes on nanoscale

switching characteristics and piezoelectric response in polycrystalline BiFeO3 thin

films. J. Appl. Phys., 110, 084102 (2011)

4. F. Yan, M. O Lai, L. Lu and T. J. Zhu. Variation of leakage mechanism and

potential barrier in La and Ru co-doped BiFeO3 thin films. J. Phys. D: Appl.

Phys., 44, 435302 (2011)

5. F. Yan, S. Miao, I. Sterianou, I. M. Reaney, M. O. Lai, and L. Lu. Multiferroic

properties and temperature-dependent leakage mechanism of the Sc-substituted

bismuth ferrite lead titanate thin films Scripta Mater., 64, 458-461 (2011).

6. F. Yan, M. O Lai, L. Lu and T. J. Zhu. Enhanced multiferroic properties and

valence effect of Ru-doped BiFeO3 thin films. J. Phys. Chem. C, 114, 6994-6998

(2010).

7. F. Yan, M. O Lai, L. Lu and T. J. Zhu. Enhanced multiferroic properties and

domain structure of La doped BiFeO3 thin films. Scripta Mater., 63, 780-783

(2010).

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8. F. Yan, I. Sterianou, S. Miao, I. M. Reaney, M. O. Lai, and L. Lu. Magnetic,

ferroelectric, and dielectric properties of Bi(Sc0.5Fe0.5)O3-PbTiO3 thin films. J.

Appl. Phys., 105, 074101 (2009).

9. F. Yan, T. J. Zhu, M. O. Lai, L. Lu. Influence of oxygen pressure on the

ferroelectric properties of BiFeO3 thin films on LaNiO3/Si substrates via laser

ablation. Appl. Phys. A. 101, 651 (2010).

10. F. Yan, I. Sterianou, S. Miao, I. M. Reaney, M. O. Lai, and L. Lu. Multiferroic

properties of Bi(Fe0.5Sc0.5)O3-PbTiO3 thin films. Phys. Scr. T139. 014003, (2010).

11. H. Xia, F. Yan, M. O. Lai and L. Lu. Electrochemical properties of BiFeO3 thin

films prepared by pulsed laser deposition. Funct. Mater. Lett. 2, 163 (2009).

12. A. Kumar, F. Yan, K. Y. Zeng, L. Lu. Eelectric, magnetic and mechanical

coupling effects on ferroelectric properties and surface potential of BiFeO3 thin

film. Funct. Mater. Lett. 4, 91 (2011).

Before Ph.D. Period:

13. F. Yan, T. J. Zhu, S. N. Zhang, X. B. Zhao, Microstructure and thermoelectric

properties of cubic AgPb18 Sb1-xTe20 (x = 0.1, 0.3, 0.5) compounds, Phys. Scr.

T129 116, (2007).

14. F. Yan, T. J. Zhu, XB Zhao, and SR Dong. Microstructures and thermoelectric

properties of GeSbTe based layered compounds. Appl. Phys A. 88, 425 (2007).

15. F. Yan, T. J. Zhu, X. B. Zhao, S. R. Dong, A study of the crystallization kinetics of

Ge-Te amorphous system, J. Univ. Sci. Techn. Beijing, S1, 64, (2007).

XIX

16. T. J. Zhu, F. Yan, X. B. Zhao, Preparation and thermoelectric properties of bulk

nanocomposite with amorphous/nanocrystal in-situ hybrid structure, J. Phys. D,

40, 6094 (2007).

17. T. J. Zhu, F. Yan, S. N. Zhang, X. B. Zhao, Microstructure and electrical

properties of quenched AgPb18Sb1-xTe20 thermoelectric materials, J. Phys. D,

40, 3537 (2007).

18. T. J. Zhu, Y. Q. Cao, F. Yan, XB Zhao, Conference "Nanostructuring and

Thermoelectric Properties of Semiconductor Tellurides", International Conference

on Thermoelectrics 2007, Jeju, Korea.

19. F. Yan, T. J. Zhu, X. B. Zhao, Progresses in study on phase-change-materials,

Funct. Mater. 37, 329, (2006).

Chapter 1 Introduction

This chapter is intended to provide a short overview of the applications and development of bismuth ferrite. In addition, the current research issues in bismuth ferrite are briefly discussed. The motivations of this thesis are finally shown at the end of the chapter.

1.1 Overview & Motivations.

Multiferroics represent an appealing kind of multifunctional materials that coexist in several ferroic orders such as ferroelectricity and ferromagnetism [1]. The simultaneous several order parameters bring about novel physical phenomena and exhibit many possibilities for new device functions [2]. Of particular interest is the existence of a cross-coupling between the magnetic and the electric orders in term of magnetoelectric coupling. This coupling enables the control of ferroelectric polarization by a magnetic field and, conversely, the manipulation of magnetization by an electric field [3].

Ferroelectric random access memories (FeRAMs) have recently achieved faster access speeds (5 ns), higher densities (256 Mb) and embodiments in several traditional materials (Pb(Zr,Ti)O3, (Ba,Sr)TiO3), but they are limited by the need for a destructive read and reset operation and suffer from atmospheric contamination. The modified composition of bismuth ferrite (BiFeO3 or BFO) enables data storage capacity to increase to five times greater than the materials currently used in FeRAM production

[4]. By comparison, magnetic random access memories (MRAMs) have been lagging far behind [5]. As ferroelectric polarization and magnetization are used to encode binary information in FeRAMs and MRAMs respectively, the coexistence of magnetization and polarization in a multiferroic material allows the realization of four-

state logic in a single device [6].The magnetoelectric coupling in multiferroics provides an opportunity to be used as a write scheme based on the application of a voltage rather than on large currents to reduce the writing energy of MRAMs.

Several issues should however be solved before realization in devices. The problems are high leakage current density, chemical fluctuation and large coercive field, and inhomogeneous magnetic spin structure [7]. Many attempts have been made to enhance the ferroelectric and ferromagnetic properties of BFO via ion substitution to introduce a “chemical pressure” into the crystal to vary the electronic and crystalline structure. La, Sm, Tb, Gd and Nd partially substitute Bi bring about an improvement of ferroelectricity and enhanced homogenization of spin arrangement [7-11]. Doping with

Mn, Sc, Cr, Ti and Nb ions into Fe-site has been attempted to eliminate oxygen vacancies in order to decrease the leakage current and change the overall magnetic spin structure [12, 13].

The mechanisms behind the enhanced ferroelectric and magnetic of chemically modified BFO films have been intensively explored from the prospects of both experimental and theoretical simulations, and abundant evidences have indicated that these issues are close to the defects chemistry in the BiFeO3 films [2]. However, a systematic study of doping effects on the ferroelectric and magnetic properties of BFO is still lacking. Therefore, the aim of this research is to systematically investigate the effects of elemental substitutions on the ferroelectric and magnetic properties of BFO films via pulsed laser deposition. In this study, the macroscopic ferroelectric, dielectric, magnetic properties and microscopic ferroelectric, magnetic properties were investigated. For each substituted system, the macroscopic ferroelectric switching behaviors and magnetic spin were also discussed. Meanwhile, the nanoscale domain structure, domain wall mobility and nanoscale switching behaviors were studied via

piezoresponse force microscopy. Furthermore, the surface potential distribution and evolution were also looked into for understanding the data storage process. Finally, the defect chemistry and magnetic phase transition in the pure and chemically modified

BFO at low temperature were characterized using electron paramagnetic resonance

(EPR).

1.2 Outline

The thesis is organized as follows:

Chapter 1 gives an introduction to the background and the research motivations and the scope of this study.

Chapter 2 provides a literature review of the structure and applications of bismuth ferrite. The basic physics and unresolved aspects of bismuth ferrite and device applications are summarized.

Chapter 3 describes the experimental procedures and principles applied in this work.

Chapter 4 contains the deposition and characterization of the pure BFO thin films.

Chapter 5 presents the deposition and characterization of the chemically modified BFO thin films

Chapter 6 focuses on leakage current mechanism of the pure and chemically modified

BFO thin films.

Chapter 7 systematically explores the magnetic phase transition in BFO via EPR.

Chapter 8 concludes the main findings and suggests some works that could be carried out to further explore the materials.

Chapter 2 Literature Review

An introduction to magnetoelectric effect and multiferroic materials is mentioned in

Section 2.1while the most important single phase multiferroic BiFeO3 (BFO) is presented systematically in Section 2.2. Moreover, the current research and status of the BFO is shown in detail. Furthermore, several approaches to enhance the multiferroic properties of BFO are discussed in Section 2.2 and the applications of

BFO are discussed in Section 2.3.

2.1 Magnetoelectric effect and multiferroic materials

2.1.1 Magnetoelectric effect[14]

The magnetoelectric (ME) effect is the phenomenon of tuning of magnetic properties with the application of an external electric field and vice versa [14, 15]. The effects can be linear or/and non-linear with respect to the external fields. The magnetoelectric effect in a single crystal is traditionally described [16] in Landau theory [17] by writing the free energy F of the system in terms of an applied magnetic field, H, whose ith component is denoted as Hi, and an applied electric field E whose ith component is denoted as Ei [14]. In a non-ferroic material, both the temperature-dependent electrical polarization Pi(T) and the magnetization Mi(T) are zero in the absence of applied fields and there is no hysteresis. It may be represented as an infinite homogeneous and stress- free medium by writing F under the Einstein summation convention in S. I. units as:

11 ijk FEHEEHHEHEHH(,) 00 ij i j    ij i j   ij i j  i j k 2 2 2  ijk HEEi j k ... 2 (2.1)

Chapter 2 Literature Review

The first term on the right hand side describes the contribution resulting from the electrical response to an electric field, where the permittivity of free space is denoted

as ε0, and the relative permittivity ij()T is a second-rank tensor that is typically

independent of Ei in non-ferroic materials. The second term is the magnetic equivalent

of the first term, where ij ()T is the relative permeability, and 0 is the permeability

of free space. The third term describes the linear magnetoelectric coupling via ij ()T ;

the third-rank tensors ijk ()T and  ijk ()T represent the higher-order magnetoelectric coefficients [14].

The effect can be expressed in the following form:

ijk Piij H j  H j H k ... 2 (2.2)

 ijk 0MEEEi ji j  j k ... 2 (2.3)

In ferroic materials, the above analysis is less rigorous because ij ()T and ij ()T display field hysteresis. Moreover, ferroics are better parameterized in terms of resultant rather than applied fields. This is because it is then possible to account for the potentially significant depolarizing/demagnetizing factors in finite media, and also because the coupling constants would then be functions of temperature alone, as in standard Landau theory. In practice, the resultant electric and magnetic fields may sometimes be approximated by polarization and magnetization respectively.

Chapter 2 Literature Review

2.1.2 Multiferroics

Multiferroic materials which show simultaneously two or three ferroic properties: ferroelectric, ferromagnetic, and ferroelastic ordering, exhibit unusual physical properties, and in turn promise new applications in devices, as a result of the coupling between their dual order parameters[3]. Figure 2.1 shows the phase control in different types of couplings present in the materials. The electric field E, magnetic field H, and stress ζ control the electric polarization P, magnetization M, and strain ε, respectively.

Figure 2.1 Phase control in ferroic and multiferroics. [18]

Ferroelectricity is a spontaneous electric polarization of a material that can be reversed by the application of an external electric field [19]. Ferroelectric materials possess at least two equilibrium orientations of the spontaneous polarization vector in the absence of an external electric field, and the spontaneous polarization can be switched between those orientations by an electric field. Ferroelectric materials undergo a phase transition from a high-temperature phase (paraelectric) state above the Curie temperature (TC) to a low temperature phase (ferroelectric) that contains a spontaneous polarization (PS). The most widely studied ferroelectric materials are oxides with a perovskite structureof the form ABO3.

Chapter 2 Literature Review

Ferromagnetism is a phenomenon that involves spontaneous magnetization below a critical temperature [20]. The spontaneous magnetization can be switched and saturated (MS) along the direction of external magnetic field (H). There is a remanent magnetization (Mr) once the field is removed. On the other hand, the definition of multiferroics can be expanded as to include non-primary order parameters such as antiferromagnetism. In materials that exhibit antiferromagnetism, the magnetic moments of atoms or molecules, usually related to the spins of the electrons, are aligned in a regular pattern with neighboring spins (on different sublattices) pointing in opposite directions. Generally, antiferromagnetic order may exist at sufficiently low temperatures, vanishing at and above a certain temperature, the Néel temperature [21].

Above the Néel temperature, the material is typically paramagnetic.

Ferroelasticity is a phenomenon in which a material may exhibit a spontaneous strain.

When a stress is applied to a ferroelastic material, a phase change will occur in the material from one phase to an equally stable phase either of different crystal structure

(e.g. cubic to tetragonal) or of different orientation (a 'twin' phase). This stress-induced phase change results in a spontaneous strain in the material[22].

In general, a variety of mechanisms can cause lowering of symmetry resulting in multiferroicity, which is charge ordered, geometrically frustrated multiferroics, magnetically driven ferroelectricity, and lone pair multiferroics. Much efforts have been devoted in the studies of these mechanisms, see for example, references [22-26].

In this study, the main investigation is focused on multiferroic materials exhibiting simultaneously ferroelectricity and ferromagnetism. The magnetoelectric multiferroic coupling effects are controlled of spontaneous polarization by applying a magnetic field and vice versa.

Chapter 2 Literature Review

2.1.3 Single phase multiferroics [3]

A single-phase multiferroic material is one that possesses two- or all-three-of the so- called “ferroic” properties, as can be seen in Figure 2.2.

Figure 2.2 Relationship between multiferroic and magnetoelectric.

The first ferromagnetic ferroelectric material discovered was nickel iodine boracite,

Ni3B7O13I [27]. The search for other ferromagnetic ferroelectrics began in Russia in the 1950s, with the replacement of some of the d0 B cations in ferroelectric perovskite oxides by magnetic dn cations [28]. The first synthetic ferromagnetic ferroelectric material, (1-x)Pb(Fe2/3W1/3)O3-xPb(Mg1/2W1/2)O3, was produced in the early 1960s using this approach [29]. Other examples include B-site-ordered Pb2(CoW)O6 and

Pb2(FeTa)O6 [30] which are ferroelectric and antiferromagnetic; and Pb2(FeTa)O6 [31] and Pb2(FeTa)O6, [32] which are both ferroelectric and antiferromagnetic with weak ferromagnetism below around 10 K. As a result of dilution of magnetic ions, all these materials have rather low Curie or Neel temperatures.

Several families of multiferroic fluorides are well known, including the group BaMnF4,

BaNiF4, BaCoF4, whose antiferroelectric transition near 251 K was first analyzed by

Ryan and Scott in 1974 [33] and studied in detail by Fox et al [34]. However, it is important to note that the perovskite fluorides such as KCoF3, KMnF3, RbCoF3,

Chapter 2 Literature Review

TlCoF3, etc., are generally cubic and hence are of no interest with regard to magnetoelectricity. Thus, magnetoelectricity in multiferroic fluorides is usually of more concern to complex structures such as that of BaMF4 or the orthorhombic tungsten bronze K3Fe5F15 which are not so amenable to theoretical modeling.

There exists yet another class of compounds which are often cited as violating the “d0- ness” rule: hexagonal manganites RMnO3 (R = Y or small rare earths). Sometimes there are called hexagonal perovskites, although in fact it is a misnomer: despite apparently similar forumula ABO3, these systems have different crystal and electronic structure. Thus, in contrast to conventional perovskites and even to quasi-one-

3+ dimensional hexagonal perovskites like CsNiCl3, YMnO3 Mn ions are located not in

O6 octahedra, but are in a 5-fold coordination – in the centre of the O5 trigonal biprysm

[35]. There are still other new multiferroic materials reported recently, such as

TbMn2O5 demonstrated by a highly reproducible electric polarization reversal and permanent polarization imprint that are both actuated by an applied magnetic field [36].

TbMn2O5 exhibits a profound interplay between electrical polarization and the applied magnetic field. The long-sought control of electrical properties by magnetic fields has recently been achieved in a rather unexpected class of materials known as “frustrated magnets”, for example the RMnO3. Like many of the rare-earth manganites, RMn2O5, where R denotes rare earths from Pr to Lu, Bi and Y, shows four sequential magnetic transitions: incommensurate sinusoidal ordering of Mn spins at T1 = 42 ~ 45 K, commensurate antiferromagnetic ordering at T2 = 38 ~ 41 K, reentrance transition into the incommensurate sinusoidal state at T3 = 20 ~ 25 K, and finally, an ordering of rare- earth spins below T4 < 10 K [37, 38].

Chapter 2 Literature Review

2.1.4 Multiferroic composites

The choice of single-phase materials exhibiting the coexistence of strong ferromagnetic and ferroelectricity at room temperature is quite limited. In the past decades, various ceramic composites consisting of piezoelectric and magnetic oxide ceramics have been investigated experimentally, including BaTiO3, PZT, Pb(Mg,

Nb)O3(PMN), PbTiO3(PTO) ferroelectric phase with magnetic phase, such as Ni ferrites, Co ferrite, Li ferrite and Cu ferrite [39]. Among them, particulate ceramic composites are most easily prepared via conventional sintering technique. The leakage problem due to high concentration of ferrite phase with low resistivity in the particulate composite ceramics can be eliminated in the laminate composite ceramics.

Magnetoelectric multiferroic composites consist of a ferroelectric phase and ferromagnetic phase, and the coupling between the two orderings is through stress mediation. The magnetoelectric effect is extrinsic in this case since magnetoelectric effect is not exhibited by any of the constituent phases on their own. Various constituent materials have been studied as multiferroic composite materials, such as

BaTiO3-CoFe2O4, BaTiO3-NiFe2O4, PZT-NiFe2O4, PZT-CoFe2O4, BiFeO3-CoFe2O4, and BiFeO3-NiFe2O4 [3, 40-43].

2.2 Multiferroic BiFeO3

2.2.1 Structure

Single crystal BFO owns a rhombohedrally distorted perovskite structure belonging to the space group of R3c with a lattice parameter of a = 5.62 Å and α = 59.35° [44].

Meanwhile, the structure of BFO could also be characterized by two distorted perovskite unit cells (ar = 3.96 Å, αr = 89.3~89.4°) connected along their body diagonal, denoted by the pseudocubic <111>, to form a rhombohedral unit cell [45]. The Fe-O-

Chapter 2 Literature Review

Fe angle is important since it dominates both the magnetic exchange and orbital overlap between Fe and O, and as such it controls the magnetic ordering temperature and the conductivity [46]. One of the Bi-Fe distances also has a local maximum near

Néel temperature, and the atomic vibrations of Bi3+ and O2- ions show a significant anisotropy [47].

The ferroelectric state arises from a large displacement of the Bi ions relative to the

FeO6 octahedra. The ferroelectric polarization lies along the pseudocubic <111> leading to the formation of eight possible polarization variants, corresponding to four structural variants1 [48]. The antiferromagnetic ordering of BFO is G-type, in which the Fe magnetic moments are aligned ferromagnetically within (111) and antiferromagnetically between adjacent (111) [45]. In addition, BFO exhibits a spin cycloid structure in the bulk [49] while the preferred orientation of the antiferromagnetically aligned spins is in the (111), perpendicular to the ferroelectric polarization direction with six equivalent easy axes within that plane (Figure 2.3) [50].

BFO thin films have shown the existence of a large ferroelectric polarization as well as a small net magnetization of the Dzyaloshinskii–Moriya type resulting from a canting of the antiferromagnetic sublattices [51].

Chapter 2 Literature Review

Figure 2.3 Schematic of crystal structure of BFO and ferroelectric polarization (arrow) and antiferromagnetic plane (shaded planes) [45].

2.2.2 Ferroelectricity

The values of ferroelectric polarization in bulk BFO are along the diagonals of the perovskite unit cell of <111>pseudocubic and strongly sample-dependent of Pr such as 6,

8.9, 11.7, 40, and 60 Ccm-2 [52-55]. It took more than 30 years before they were proved right by measurements on high-quality thin films, single crystals and ceramics[46]. Thin films of BFO have demonstrated excellent ferroelectric

-2 properties[56] (Pr = 60 Ccm along <100>pseudocubic) and the films often present different crystallographic structure compared with that of single crystal.

Figure 2.4 shows the polarization of bulk single crystal and epitaxial thin film. A major problem in the BFO films is their low resistivity and high leakage, caused by defects and nonstoichiometric compositions in the BFO materials. Considerable attempts have been made to enhance the ferroelectric and properties of BFO via ion substitution/codopping A/B-sites [2]. “Chemical pressure” has been introduced into the perovskite crystal (ABO3) to vary the electronic and crystalline structure. Ca, La, Sm,

Tb, Gd and Nd ions partially substituting A-site bring improvement in ferroelectricity and enhance homogenization of spin arrangement [57-60]. Doping of Mn, Sc, Cr, Ti

Chapter 2 Literature Review and Nb ions into B-site could eliminate oxygen vacancies, decrease the leakage current, and change the overall magnetic spin structure [61-63]. In the mean time, other ABO3 perovskite materials have also been introduced to form solid solutions with BFO. For example, ferroelectric properties have been enhanced in PbTiO3, PbZrO3, and BaTiO3 modified BiFeO3 ceramics and thin films [64-66]. The insertion of other ABO3 perovskite into BFO not only stabilizes the BFO perovskite phase but also forms a morphotropic phase boundary (MPB) with BFO due to their different crystal symmetries.

Figure 2.4 Polarization of BiFeO3: (a) bulk single crystal [55] and (b) epitaxial thin film [56].

2.2.3 Dielectric Properties

The voltage dependence of dielectric constant of (001)-oriented BFO thin film is about

100, which is smaller than those of typical perovskite ferroelectrics such as Pb(Zr,

Ti)O3 (PZT), (Ba, Sr)TiO3, BaTiO3 etc [67]. The low dielectric constant is associated with the low piezoelectric coefficient, which is the reason for the relatively small sensitivity of BFO to epitaxial strain [68]. The refractive dielectric constant could be estimated to be n=2.5while the optical frequency dielectric constant of K=n2=6.25 [69].

BFO is in fact strongly birefringent with ∆n=ca. 0.34, meaning that the dielectric constant at optical frequencies is very anisotropic [46].

Chapter 2 Literature Review

The dielectric constant could be modified by introducing dopants such as Ba, La, Mn,

Co, Nb, etc [70-73]. At low or high temperatures, colossal dielectric constants have been observed due to finite conductivity leading to Maxwell-Wagner (M-W) behavior

[74, 75]. The temperature at which M-W effects set in depends on the sample conductivity and the dopants.

2.2.4 Magnetism

The local short range magnetic ordering of BFO is G-type antiferromanget: each Fe3+ spin is surround by six antiparallel spins on the nearest Fe neighbors. The spins are in fact not perfectly antiparallel as there is a weak canting moment caused by the local magnetoelectric coupling to the polarization [46]. The antiferromagnetic spin ordering is not homogeneous but is manifested as an incommensurated cycloid structure with a wavelength of ~ 64 nm along <110>, as can be seen in Figure 2.5 [76]. The spin rotation plane can also be determined because the magnetic scattering amplitude depends on the component of magnetic moments perpendicular to the scattering vector

[49]. The magnetic Néel temperature is about 643 K and the cycloid could be distorted at low temperatures [77].

Figure 2.5 Schematics of the 64 nm antiferromagnetic circular cycloid. The canted antiferromagnetic spins (blue and green arrows) give rise to a net magnetic moment (purple arrows) that is specially averaged out to zero due to the cycloidal rotation. The spins are contained within the plane defined by the polarization vector [46].

Chapter 2 Literature Review

Meanwhile, spin reorientation transition has been identified in BFO at 140 and 200K, by different research groups independently [78, 79]. The magnon anomalies observed at 140K and 200K suggest phonon–magnon coupling in each case and unscreened strain and the consequent possibility of mean-field behavior [80].

Sc, Mn, Nb, Zr, Co and Ti ions substitute of Fe ions would increase the magnetization by changing the Fe valence state due to charge compensation and varying the Fe-O-Fe bond angle to increase the spin [81-85]. For Bi-site substitution, Yb, La, Gd, Nd, Pb,

Ca, Sr ions may also impact the magnetization of the BFO stem from the spatial homogenization of spin arrangement. The disturbed spin cycloid structure when La ions positions in Bi-site, the formation of partial Fe2+ due to the Bi ions volatilization or oxygen vacancies, varied canting angle of Fe-O-Fe bond due the distortion of FeO6 octahedron caused by introducing La ions. The increased tensile strain changes the balance between the antiferromagnetic and ferromagnetic interactions [86-89].

As for the impact for dopants on the magnetic phase transition, there are very few report studies. In this work, the doping effect on the spin-reorientation transition via the EPR as a function of temperature will be investigated. This would be discussed in

Chapter 7.

2.2.5 Magnetoelectric Coupling

The magnetoelectric coupling in BFO manifests the relationship between ferroelectric polarization and magnetic symmetry such as ferroelectric control of magnetism or magnetic control of ferroelectricity. The magnetic moments rotate with the plane defined by the polarization and the cycloid propagation vector [46], as can be seen in

Figure 2.6. The magnetic easy planes could be rotated by applying a voltage and switching the polarization by 71o in single crystal [49]. Meanwhile, the thin films

Chapter 2 Literature Review reported by Zhao et al [90] have demonstrated that the films had no spin cycloid and had the homogeneous G-type antiferromagnetism. The ferroelectric structure was determined using piezoelectric force microscopywhile X-ray photoemission electron microscopy as well as its temperature dependence were used to detect the antiferromagnetic configuration and antiferromagnetic domain switching induced by ferroelectric polarization switching [90].

Figure 2.6 Schematics of the planes of spin rotations and cycloids k~1 vector for the two polarization domains separated by a domain wall (in light gray)[49]. Rotating the polarization by 71o results in a change of the magnetic easy plane, meaning that sublattice magnetization can be switched by an applied voltage[46].

2.2.6 Domains and domain walls[46]

Research on domains and domain walls has intensified recently because firstly, the behavior of domains is directly responsible for switching characteristics, and secondly, domain size scales with sample size, so thin films can have very small domains and therefore, a high volume density of domain walls. As for BFO, the ferroelectric polarization can point along any of the four diagonals of the perovskite unit cell, with two antiparallel polarities for each direction: hence there are eight different polar domains in BFO. There are three types of ferroelectric domain wall which are usually

Chapter 2 Literature Review labeled according to the angle formed between the polarization vectors on either side of the wall. This is schematically depicted in Figure 2.7.

Figure 2.7 Schematic of the three types of ferroelectric domain walls, ferroelectric polarization (bold arrows) and antiferromagnetic plane (shaded planes) [90].

Different domain shapes have been observed by different research groups: for thick film, the domains of epitaxial BFO are striped domains [91, 92] on different substrates, such as SrRuO3 buffered SrTiO3 and TbScO3. For ultrathin films, the domains in BFO are no longer striped but instead form irregularly shaped mosaic structures or fractal domains [93]. The reason for the thickness-induced transition to fractal morphology remains unknown, but irregular walls are elastically costly, thus they cannot be the equilibrium configuration in a perfect crystal.

Domain walls have their own local symmetry and hence also their own properties. The domain walls of BiFeO3 are much more conductive compared with the domain inside

[94]. The conductivity of the domain walls is related to the type of domains. 180o walls are the most conductive, followed by 109 o and finally the 71 o walls which do not have any measureable transport enhancement [46]. There could be two reasons for the enhanced conductivity of the walls. First, the polarization normal to the domain wall is observed not to be constant across it; this generates an electrostatic depolarization field

Chapter 2 Literature Review that may attract charge carriers. Second, the electronic bandgap is considerably reduced for the 180o and 109 o domain walls.

A plausible explanation for the band gap decreases has to do with the local distortion of the Fe-O-Fe bond angle. The bond angle which controls the orbital overlap. In the middle of the domains, the octahedra are quite buckled and hence the gap is big. If the unit cell expands, the bucking angle can become straighter, thereby increasing the orbital overlap and reducing the band gap. The local suppression of polarization at the domain walls leads to precisely such a volume expansion via the cancelling of the spontaneous strain and hence, the local straightening of the bond angles at the walls reduces the gap in much the same way as temperature or pressure would. The absolute

o value of the polarization is smallest in the middle of the 180 walls (Pwall =0),

o o intermediate in the 109 walls (Pwall =0.57P0), and maximum in the 71 walls ( Pwall

=0.82 P0). Therefore, the volume change (and the associated change in conductivity) will be biggest for the 180o walls, intermediate in the 109 o walls and smallest in the 71 o walls.

2.3 Device application

2.3.1 Data storage [1]

The magnetoelectric coupling in multiferroic provides an opportunity for magnetoelectric random access memories (MERAMs) to combine the magnetoelectric coupling with the interfacial exchange coupling between a multiferroic and a ferromagnet to switch the magnetization of the ferromagnetic layer by using a voltage.

As shown in Figure 2.8. BFO could act not only as the magnetoelectric active layer, but also as the tunneling barrier. In MERAMs, the magnetoelectric coupling enables an electric field to control the exchange coupling at the interfaces of the multiferroic with

Chapter 2 Literature Review the ferromagnet. The exchange coupling across the interface then controls the magnetization of the ferromagnetic layer, so that ultimately this magnetization can be switched by the electric polarization of the multiferroic [1].

Figure 2.8 MERAMs based on exchange-bias coupling between a multiferroic that is ferroelectric and antiferromagnetic (FE-AFM, green layer), and a thin ferromagnetic electrode (blue layer). A tunneling barrier layer between the two top ferromagnetic layers provides the two resistive states [1, 46].

2.3.2 Optoelectronic devices

The band gap of the BFO is about 2.8 eV yielded from the screened-exchange density functional theory [95]. The ability to tune the oxidation state of the Fe ion in BFO may enable researcher to engineer the gap and conductivity to enhance the photoferroelectric properties of BFO [96]. The experimental results indicate that BFO not only possesses a band gap in the visible range but also that it displays a significant photoresponse which improves with increasing oxygen vacancy concentration [96].

The bulk electric polarization in BFO with a small optical gap edge of ~ 2.2 eV is highly nonlinear and unidirectional [97]. The photovoltaic effect (PV) in ferroelectric is distinctly different from the typical PV effect in semiconductor p-n junctions. The observation of switchable diode and photovoltaic effect in BFO reveals unusual and

Chapter 2 Literature Review intriguing charge conduction behavior in leaky ferroeelctrics, and warrants advanced studies of BFO based multifunctional devices [97]. The change in photocurrent at zero bias with q shown in Figure 2.9 with blue circles, closely follows a sinusoidal form with a periodicity of 180°. The maximum of photocurrent was observed when the polarized-light electric field was along the in-plane ferroelectric polarization, and the current was minimal when the light electric field was perpendicular to the in-plane ferroelectric polarization. After the initial rotation experiment, the polarizer was rotated by 90°, and lighting conditions were readjusted for an optimum photocurrent.

The change in photocurrent (green circles) with q after this polarizer rotation is similar to that before rotating the polarizer, except for a 90° phase shift of q. This demonstrates that the reproducible angular dependence is not due to any artifact of the optical setup [97].

Figure 2.9 Variation of photocurrent with sample rotation under illumination with a linearly polarized light. The experimental sketch is shown in the inset. The PV effect becomes maximum when the polarized-light electric field is parallel to the in-plane component of the ferroelectric polarization, and minimum when the field is perpendicular to the in-plane component [97]

Chapter 2 Literature Review

2.4 Thin film deposition and characterization

2.4.1 Pulsed Laser Deposition (PLD)

PLD, as shown schematically in Figure 2.10 is a physical vapor deposition process. A pulsed laser is focused onto a target to be grown at a vacuum chamber with a rotating target holder and a substrate heater. The ablation plume provides the materials of target flux for thin film growth. Multiple targets can be performed during the deposition process, enabling the deposition of multilayered films.

Generally, a background gas is introduced as a component of the flux and the reactant gas required for phase formation. High quality film growth could be obtained through optimized deposition parameters such as substrate temperature, laser energy density and frequency, target-to-substrate distance and background gas pressure. One of the most important and enabling characteristics in PLD is the ability to realize stoichiometric transfer of ablated materials from multication targets for many materials.

This advantage arises from the non-equilibrium nature of the ablation process itself due to absorption of high laser energy density by a small volume of materials[98].

Figure 2.10 Schematic of pulsed laser deposition system

Chapter 2 Literature Review

2.4.2 X-Ray diffraction (XRD)

X-ray diffraction is a method of determining the arrangement of atoms in a crystal.

Here a beam of X-ray strikes a crystal and diffracts into many specific directions.

Crystals are regular arrays of atoms, and X-rays can be considered waves of electromagnetic radiation. The relationship between them can be expressed by Bragg's law: nd 2 sin (3.1) where n is an integer, λ is the wavelength of X-ray, d is the lattice interatomic spacing, and θ is the diffraction angle.

X-ray structural analysis could be used to characterize the microstructure and epitaxial feature of the thin films. θ-2θ scans are used to determine the crystalline orientation of the thin films while the  scans can be used to determine the epitaxy of the films. The

XRD system used in this study was X-ray diffraction (XRD) with Cu Kα radiation using a Shimadzu XRD-7000 diffractiometer attached with a pole figure attachment.

2.4.3 Atomic force microscopy (AFM)

Atomic force microscopy (AFM) is a high-resolution type of scanning probe microscopy, with demonstrated resolution in the order of fractions of a nanometer. The

AFM consists of a cantilever with a sharp tip (probe) at its end that is used to scan the specimen surface. Compared with Scanning Electron Microscope, AFM provides extraordinary topographic contrast direct, height measurements and unobstructed views of surface features (no coating is necessary). In this study, the surface morphologies and grain sizes of films were obtained by AFM (Asylum Research MFP-

3DTM).

Chapter 2 Literature Review

2.4.4 Macroscopic Ferroelectric measurement

Ferroelectric properties of the films were characterized using a commercial Radiant

Precision workstation (Radiant Technologies, USA). Figure 2.11 shows the standard

Radiant hysteresis test. Standard delay is 1 second for the recovery of the ions displacement.

Figure 2.11 Schematic of the experimental set-up for ferroelectric hysteresis loop measurements.

Pt top electrodes of diameter of 100 μm were sputtered on the surface of the films through a shadow metal mask to form metal-insulator-metal (MIM) capacitors for the electrical property tests. The capacitor structure is shown in Figure 2.12. The polarization could be determined as follows [99]:

For an ideal ferroelectric insulator

QPA 2 r (3.2)

where Pr is the remanent polarization and A is the electrode area for a parallel-plate capacitor. For a somewhat conductive sample

Q2 Pr A Ea t (3.3)

Chapter 2 Literature Review

where ζ is the electrical conductivity, Ea is the applied field, and t is the measuring time. Thus Q is a pulsed measuring system depends on the pulse width.

Figure 2.12 Schematic structure of the electrical measurement.

2.4.5 Dielectric measurement

Dielectric constant and loss tangent of the films were measured without DC bias using a precision impedance analyzer (Wayne Kerr Electronics 6500B Series). A small AC bias with 100 mV was applied to the capacitors at various frequencies. The capacitance was recorded and the dielectric constant evaluated from,

Cd   (3.4) 4 0 A

where  is dielectric constant,  0 is the dielectric constant in free space, and C is the capacitance.

The relative dielectric permittivity of a material in an alternating electric field can be represented by a complex form:   '''j  , where the real part  ' represents the r r r r energy stored and the imaginary part  ''represents the energy loss of the field in the r

material. The dielectric constant  r and dielectric loss tanδ are expressed as:

Chapter 2 Literature Review

 '2  ' '2 (3.5) r rr tan ''' /  (3.6) rr

2.4.6 Piezoresponse force microscopy (PFM)

PFM measures the mechanical response when an electrical voltage is applied to the sample surface with a conductive tip of an AFM. In response to the electrical stimulus, the sample then locally expands or contracts as shown in Figure 2.13. When the tip is in contact with the surface and the local piezoelectric response is detected as the first harmonic component of the tip deflection, the phase j of the electromechanical response of the surface yields information on the polarization direction below the tip.

For c -domains (polarization vector oriented normal to the surface and pointing downward), the application of a positive tip bias results in an expansion of the sample, and surface oscillations are in phase with the tip voltage, j = 0. For c+ domains, the response is opposite and j = 180°.

Figure 2.13 Schematic of PFM operation [100].

The relationship between strain and applied electric field (often referred to as the

“inverse piezo effect”) in piezoelectric materials is described by a rank-3 tensor. The most important component of this tensor for typical “vertical” PFM is the d33

Chapter 2 Literature Review component [101] since it is coupled directly into the vertical motion of the cantilever.

The voltage applied to the tip is

Vtip V dc V ac cos( t ) (3.7) resulting in piezoelectric strain in the material that causes a cantilever displacement of

z zdc  A( , V ac , V dc )cos(  t   ) (3.8) due to piezoelectric effect [102]. When the voltage is driven at a frequency well below that of the contact resonance of the cantilever, equation (3.8) becomes

z d33 Vac  d 33 V ac cos( t  ) (3.9)

where d33 is the displacement along the polarization direction. From Figure 2.13, the magnitude of the oscillating response is a measure of the magnitude of d33 while the phase is sensitive to the polarization direction of the films and the direction of sample polarization determines the sign of the response.

2.4.7 Switching Spectroscopy PFM (SS-PFM)

Switching spectroscopy piezoresponse force microscopy (SS-PFM) is used to address quantitatively the local switching characteristics of ferroelectric materials. The theoretical model for the hysteresis loop shape has been developed and the influence of nonideal tip-surface contact and electrostatic force contribution on experimental measurements has been determined [103]. SS-PFM refers to locally generated hysteresis loops in ferroelectric materials. From these hysteresis loops, information on local ferroelectric behaviors such as imprint, local work of switching, and nucleation biases can be obtained.

Chapter 2 Literature Review

As can be seen in Figure 2.14, a sine wave is carried by a square wave that steps in magnitude with time in SS-PFM [103]. Between each ever-increasing voltage step, the offset is stepped back to zero with the AC bias still being applied to determine the bias-induced change in polarization distribution (e.g. the size of the switched domain).

It is then possible to see the hysteresis curve of the switching of the polarization of the surface [100].

Figure 2.14 Switching spectroscopy PFM diagram [103].

2.4.8 Kelvin probe force microscopy (KPFM)

KPFM is a scanning probe method where the potential offset between the probe tip and a surface can be measured using the same principle as a macroscopic Kelvin probe.

The cantilever in the AFM is a reference electrode that forms a capacitor with the surface over which it is scanned laterally at a constant separation. The cantilever is not piezoelectrically driven at its mechanical resonance frequency as in normal AFM although an alternating current (AC) voltage is applied at this frequency. With KPFM, the work function of the surfaces can be observed at atomic or molecular scales [104].

The work function relates to many surface phenomena including catalytic activity,

Chapter 2 Literature Review reconstruction of surfaces, doping and band-bending of semiconductors, charge trapping in dielectrics and corrosion [105]. The map of the work function produced by

KPFM gives information about the composition and electronic state of the local structures on the surface of a solid.

2.4.9 Vibration sample magnetometer (VSM)

A VSM has been used to determine the magnetic properties of the films [106]. The principle of this measurement is as follows: When a sample is placed in a homogeneous magnetic field, a magnetic moment is induced in the sample. If this sample is made to undergo sinusoidal motion, the vibration induces a magnetic flux change. This in turn induces a voltage in the pick-up coils. The magnetic moment determined by the VSM is related to the magnetization of the sample and its susceptibility. In this study, the magnetic property is characterized using a Lakeshore

736 vibrating sample magnetometer (VSM). The sample is suspended from a vibrating drive head by a non-magnetic rod and placed between two electromagnets which produce a magnetic field. The vibrator generates a vertical sinusoidal vibration, thus the sample experiences sinusoidal motion, inducing an electrical signal in the coils mounted on the pole faces to the electromagnets. The signal picked up by the coils is proportional to the frequency and amplitude of the sinusoidal motion, and the total magnetic moment of the sample at the applied magnetic field. The frequency and amplitude of the sinusoidal are maintained constant by a capacitor. By feeding the signal from the pick-up coils and the reference signal into a demodulator, the magnetic moment of the sample is extracted.

2.4.10 Magnetic force microscopy (MFM)

MFM is a variety of atomic force microscope [106] where a sharp magnetized tip scans the magnetic samplesuch that the tip-sample magnetic interactions are detected

Chapter 2 Literature Review and used to reconstruct the magnetic structure of the sample surface [22]. In MFM, the probe must have a magnetized tip. The magnetic probe is standard silicon cantilever coated with a Co thin film. The interaction of the tip with various magnetic domains on the surface results in cantilever deflection. The microscope can sense the deflection of the cantilever which results in a force image in a static mode. A mapping of the magnetic forces above the sample surface can be achieved when the sample is scanned under the tip. The force gradient ( F ' ) detected contains information of both the surface structure and surface magnetization:

'' FFF' surface magnetic (3.10)

' ' where Fsurface is the surface component of the gradient and Fmagnetic is a magnetic component of the gradient. Signals from surface topography dominate at a distance close to the surface while at a distance further away from the surface, the magnetic signal dominates. Thus, normal MFM images may contain a combination of topography and magnetic signals. MFM is capable of imaging magnetic domains of several tens of nanometers.

Chapter 3 Experimental Procedures

In this chapter, the detailed procedures of targets preparation and thin film deposition measurement are described. An overview of the techniques used to investigate the multiferroic properties of the samples is also provided.

3.1 Fabrication of Targets

In this study, pure and chemically modified BiFeO3 targets are fabricated using commercial oxide powders from Aldrich, and the fabrication process is as follows:

3.1.1 Pure BiFeO3 target

Polycrystalline BiFeO3 powders were prepared by conventional solid-state reaction using high-purity powders of bismuth oxide, Bi2O3, and iron oxide, Fe2O3 as starting compounds, and were mixed via ball milling for about eight hours using high purity alcohol as a medium. After mixing in stoichiometric proportions, powder was calcinated at 500 °C for 3 h in an alumina crucible. The calcinated material was once again ground thoroughly prior to sintering at 800 °C for 30 min to improve the synthesis uniformity.

3.1.2 Chemically Modified BiFeO3 targets

5 mol % La, Ru-doped, La, Ru codoped, 2, 5 mol% Nb doped and 0.5, 2 mol% Mn doped BiFeO3 with compositions respectively of Bi0.95La0.05FeO3, BiFe0.95Ru0.05O3,

Bi0.95La0.05Fe0.95Ru0.05O3, BiFe0.995Mn0.005O3 and BiFe0.98Mn0.02O3 ceramics with 10% excess bismuth, were prepared by conventional solid-state reaction using high-purity

(>99.99%) powders of bismuth oxide, Bi2O3, iron oxide, Fe2O3, and oxides of La2O3,

RuO2., MnO2 (provided by Aldrich) as starting compounds. The starting materials were weighed in stoichiometric proportions and thoroughly mixed via ball milling for about 12 h using 3 mm diameter Y-stabilized ZrO2 balls and high purity alcohol as

Chapter 3 Experimental Procedures media. The mixtures were dried at ~70 oC, pressed into diameter of 10 mm pellets at about 200 MPa pressure and precalcined at 600 °C for 3 h in alumina crucibles. The calcined materials were once again ground thoroughly prior to sintering at 820 °C for

o 1h for BiFeO3 and 850 C for the doped samples. Meanwhile, 2, 5 mol %

Pb(Zr0.52Ti0.48)O3 (PZT) modified BFO targets were synthesized by standard solid-state reaction using high purity Bi2O3, Fe2O3 (Aldrich, > 99.9%) and PZT with 10% excess

Bi2O3 and PbO to compensate for evaporation of Bi and Pb during sintering using the same method.

3.2 Thin film growth

In this study, the films deposited by PLD were done according to the following processes:

3.2.1. Substrate and target cleaning

Before deposition, the substrates should be cleaned in acetone and then alcohol in an ultrasonic agitation. Following that, it was cleaned in deionized water and blown via

N2 gas before being fixed onto the substrate holder. The target surface was polished by fine sandpaper before deposition.

3.2.2 Film deposition

For the film deposition, the chamber should be pre-pumped to a base pressure of < 10-5

Torr. The substrates were heated to the objective temperature via the substrate heater.

The oxygen flow to the chamber was regulated by setting the pumping valve to maintain a dynamic pressure of 25 to 100 mTorr. KrF laser with wavelength of 248 nm and laser energy of 200 mJ were used. The target-to-substrate distance was fixed at

45mm. The chemically modified BFO targets prepared as per above mentioned and

Chapter 3 Experimental Procedures

metallic oxide targets such as SrRuO3, LaNiO3, La(Sr0.5Co0.5)O3 were used. Details of the deposition parameters are shown in Table 3.1.

Table 3.1 Deposition parameters for thin films

Substrate Laser O2 pressure Energy density Targets Temperature repetition Rate (mTorr) (J/cm2) (oC) (Hz)

BFO 550 25~100 2 5

La-BFO 550 25~100 2 5

Ru-BFO 550 25~100 2 5

Nb-BFO 550 25~100 2 5

Mn-BFO 550 25~100 2 5

PZT-BFO 550 25~100 2 5

SrRuO3 650 50 2 5

LaNiO3 550 50 2 10

3.3 Thin film characterization

The as-deposited thin films were cooled to room temperature at a rate of about 5 oCmin-1. Structure of the films was determined by X-ray diffraction (XRD) with Cu

Kα radiation on a Shimadzu XRD-7000 diffractometer. Pt top electrodes of diameter of 100 μm were sputtered on the surface of the films through a shadow metal mask to form metal-insulator-metal (MIM) capacitors. The surface morphology and piezoelectric properties of the thin films were examined using a piezoresponse force microscopy (PFM) by Pt/Ir-coated Si tips (spring constant of ~ 2 N/m, resonance frequency of ~ 70 kHz, and the size of the probing tip radius is ~ 20 nm). The

Chapter 3 Experimental Procedures thickness of the films and the area of the Pt top electrode were examined using a

Hitachi S4300 field emission scanning electron microscope (FESEM). Dielectric constant and loss tangent of the films were measured without dc bias using a precision impedance analyzer (Wayne Kerr Electronics 6500B Series). X-ray photoelectron spectroscopy (XPS) analysis was performed to investigate the surface chemistry of the films using a VG ESCALAB 220i-XL system with a pass energy of 20 eV for high resolution spectrum acquisition with an Al Kα (1486.6 eV) x-ray source. The ferroelectric, leakage mechanism and magnetic properties were characterized using a

Radiant Precision Workstation (Radiant Technologies, Median, New York) with a vacuum MMR K-20 temperature controller system and a Lakeshore 736 vibrating sample magnetometer (VSM) respectively.

Chapter 4 Pure BFO thin films

4.1 Introduction

In this chapter, the fabrication and characterization of pure BFO thin films by PLD is presented. The BFO films were deposited on different substrates, such as SrTiO3 single crystal with different orientations, Pt/TiO2/SiO2/Si substrate, SrRuO3 and LaNiO3 buffered Si substrate. Thin films deposited on single crystal SrTiO3 substrate were to investigate the orientation-dependent effect on the microstructure and multiferroic properties, and polycrystalline substrate for the potential used for devices in the industry. These studies include the dependence of microstructure of substrate orientation, different bottom electrode effects, nanoscale domain switching, and the surface potential of pure BFO films.

4.2 Epitaxial BiFeO3 thin films

4.2.1 Structures of (001), (011) and (111) oriented BiFeO3 thin films

Figure 4.1 shows the XRD traces of the BFO films deposited on SRO-buffered STO

(001), (110), and (111) substrates, which only reveal the (00l) (l = 1, 2, 3), (0kk) (k = 1,

2) and (111) peaks corresponding to the substrates and SrRuO3 (SRO) bottom electrodes without any impurity phases. The values of the interplanar spacing of the films are d<001> = 4.004 Å, d<011> = 2.829 Å and d<111> = 2.313 Å corresponding to the pseudocubic out-of plane lattice parameters of 4.004, 4.000, and 4.006 Å for (001),

(110), and (111) oriented BFO films respectively. These values are larger than that for

BFO bulk lattice (3.96 Å) due to the in-plane compressive strain induced by the substrate [107].

Chapter 4 Pure BFO thin films

Figure 4.1 X-ray diffraction spectra of epitaxial BFO films deposited on SRO buffered STO substrates with different orientations.

The in-plane epitaxial arrangement of the films was studied using the pole figure measurement at a fixed 2scan of 110 reflections of the STO, as shown in Figure 4.2.

These clearly indicate the “cube-on-cube” epitaxial growth of BFO films on SRO/STO substrates and the epitaxial relationship as follows: (001)BFO||(001)SRO||(001)STO,

(011)BFO||(011)SRO||(011)STO and (111)BFO||(111)SRO||(111)STO.

Figure 4.2 Pole figure plots corresponding to BFO (110) diffraction for (a) (001), (b) (011), and (c) (111)-oriented thin films, indicating a cube-on-cube growth behavior.

Chapter 4 Pure BFO thin films

In order to further determine the symmetry changes of BFO films obtained via XRD, room temperature Raman spectra have been observed as shown in Figure 4.3, showing substrate orientation Raman spectra dependent and the prominent peaks are mainly because of the BFO normal mode of vibrations rather than that of STO substrate [108].

The spectra can be indexed as a pseudocubic structure and summarized using the

irreducible representation [109] Raman 34ABE11   , which corresponds to the XRD results. Three relatively intense peaks are observed at 140, 170, and 220 cm-1, which can be assigned to the A1 mode attributed to Bi-O bond vibration [110]. The ions into

FeO6 octahedra of pure BFO would broaden BFO bonds due to increasing dispersion in Bi-O bond lengths [111] and the structural distortion could also be varied due to the strong hybridization between Bi (6s26p3) and O (2s22p4). Furthermore, negative peak shifts by a few cm-1 could be observed from (100) to (111) substrate orientation, which can be due to the variation of polar orientation soft the Bi-O bonds, enhances the number of longer Bi-O bond and reduce the covalent character [112].

Figure 4.3 Room temperature Raman spectra of the (100), (110), and (111)-oriented rhombohedral BFO thin films and of a (111) STO substrate.

Chapter 4 Pure BFO thin films

4.2.2 Surface morphology

The surface topographies of BFO films grown on (001), (011), and (111)-oriented STO substrates measured by AFM are shown in Figures 4.4a-c. According to the

Winterbottom construction [113], the equilibrium shape of the BFO films as a function of the substrate orientation could be confirmed as below: (i) for (001) orientation, BFO follows a layer-by-layer growth along the lowest energy surfaces of {001} surfaces; (ii) for (011) orientation, BFO shows stripe growth modes on a (110) orientation; (iii) for

(111) orientation, BFO forms islands characterized with three {100} surfaces. The root-mean-square (RMS) roughness measurements show the RMS in the order of 3.62,

3.02 and 6.27 nm for the BFO film (001), (011), (111) orientation respectively, indicating that substrate orientation could change the crystallite structure of the films as well as the morphology of the films.

Figure 4.4 Atomic force microscopy (AFM) images of BFO thin films deposited at (a) (001), (b) (011), (c) (111)-oriented substrate.

4.2.3 Macroscopic electrical properties of (001), (011) and (111) oriented BiFeO3 thin films

The effect of substrate orientation on the ferroelectric properties of BFO thin films is shown in Figure 4.5. The polarization hysteresis were saturated for each orientation,

-2 -2 and the remanent polarization Pr of ~ 55 μC/cm for (001), ~74 μC/cm for (011) and

~109 μC/cm-2 for (111) substrates, are consistent with the previously reported epitaxial

Chapter 4 Pure BFO thin films

BFO films[114]. The remanent value shows a simple calculated relation as follows:

3PPP(001) 6 / 2 (011) (111) , indicating that the spontaneous polarization lies close to

<111>, and that the values measured along (001) and (011) are simply projections onto these orientations [114]. The measured remanent value is higher for the ceramic samples which could be due to the reduced leakage current in film state due to reduced pores, oxygen vacancies and dislocations in the films and the strain effect in the films.

To confirm this assumption, the leakage current density was measured.

Figure 4.5 Ferroelectric hysteresis (P-E) loops of (001), (011), (111)-oriented BFO thin films measured at room temperature.

Figure 4.6 shows the leakage current density, (J), versus electric field, (E), of the films as a function of the orientation. The BFO film shows low leakage current for all orientations, while the (011) and (111) film exhibits a lower leakage of nearly one order of magnitude compared with that of (001)-orientated film at the positive bias. It is known that in BFO, Fe ions often exist in a mixed-valance Fe2+/3+ state due to introduction of oxygen vacancies. The higher leakage current in (001)-oriented film could have stemed from the higher oxygen vacancies concentration. On the other hand,

Chapter 4 Pure BFO thin films the substrate orientation could impact the domain structures and the type of domain walls, of which the conductivity is different as discussed in Chapter 2. Thus the domain wall density could also be one of the important reasons responsible for the different conductivity.

Figure 4.6 Leakage current density for (001), (011), and (111)-oriented BFO thin films measured at room temperature.

Figure 4.7 is the current-time plot showing the characteristics of dependence of the

BFO films on orientation measured at a positive bias of 5 V. It is clear that the leakage current of the (111)-oriented BFO stabilizes immediately whereas the current relaxation for the (001) and (011)-oriented BFO film is very slow, and stabilizes after about 100 ms. The time taken to stabilize leakage current is probably carrier transfer and domain dynamics dependent, such as the polarization bound charges at the surfaces of the ferroelectric layer, free charge carriers in the conducting electrodes, and domain wall motion [115].

Chapter 4 Pure BFO thin films

Figure 4.7 Leakage current and time relationship for (001), (011), and (111)-oriented BFO thin films measured at room temperature.

4.2.4 Ferroelectric domain structure using PFM

In order to investigate the microscopic ferroelectric properties, the nanoscale ferroelectric properties were investigated using PFM. Figure 4.8 shows the microscopic ferroelectric properties as well as the surface topography. The surface topographies of the BFO films grown on (001), (011), and (111)-oriented STO substrates are shown in Figures 4.8(a), 4.8(d) and 4.8(g) respectively. The BFO films also show piezoelectrically inactive regions as a function of the substrate orientations on applying a 1 V amplitude voltage, predicting that the piezoelectricity is increased in varying the orientation from (001) to (111). The out-of plane PFM images are shown in Figures 4.8(c), 4.8(f), and 4.8(i), which indicate that the ferroelectric polarization points dominantly downward toward the SRO layer. It is believed that the characteristics of the domains and domain walls directly impact the ferroelectric switching behavior [2]. It has been proven that the role of the strain relaxor to the SRO layer is minimal and the less conductive 71o domain walls rather than 109 o walls due

Chapter 4 Pure BFO thin films to the thicker (>20nm) SRO bottom electrode [116]. The (111) oriented BFO film is along the polarization vector, therefore, the image is with no in-plane PFM image

[117]. The in-plane PFM images are still under research and the detailed domain architectures of BFO as a function of the substrate orientation will be done in future work.

Figure 4.8 (a), (d), and (g) AFM images (1×1 µm 2), (b), (e) and (h) out-of-plane piezoresponse (OP-PFM) amplitude images, and (c), (f) and (i) OP-PFM phase images of (001), (011), (111)-oriented BFO thin films. Yellow (bright) and purple (dark) on the piezoresponse image correspond to negative (upward) and positive (downward) domains respectively

For piezoelectric materials, the surface displacements could be detected via the converse piezoactuation [118]. In a typical experimental setup, the piezoelectric thin film is grounded with a conducting layer such as SrRuO3 during PFM measurement.

An AC bias is applied to the tip from a signal generator. With the AFM tip acting as the top electrode, an ac electric field is built across the piezoelectric film. Thus, the surface displacement caused by the electric field could be detected as deflections of the integrated AFM cantilever using a lock-in amplifier. The amplitude of film vibration

Chapter 4 Pure BFO thin films

normal to the surface detected during PFM (zvibration) is proportional to the converse piezoelectric coefficient and the ac applied bias Vac given by the equation

zvibration d zz V ac (4.1)

In this study, PFM data were observed for the sample with different orientations by acquiring multiple 5 m scan size images at a single location on each sample. The AC voltage was increased image by image from 0.5 to 3.0 V at a step of 0.5 V without dc voltage being applied during the scan. The overall intensity of the amplitude increases with the increase in ac voltage. Histogram data obtained from each of the PFM image denote Gaussian distribution with single narrow peaks. The average piezoresponse of each PFM images observed at distinct ac excitation amplitudes is plotted for different orientations in Figure 4.9. The piezoresponse increases linearly with increasing ac biases from 0.5 to 3 V as expected. From the slopes of the plots, the piezoelectric constant, d33 could be obtained. Table 4.1 shows d33 as a function of the three crystalline orientations.

Figure 4.9 Calculated average piezoresponse of PFM images obtained at each voltage step vs ac voltage for (001), (011), (111)-oriented BFO thin films. 95% confidence error bars are smaller than 3% and too negligible to be visible on the graph.

Chapter 4 Pure BFO thin films

eff Table 4.1 Orientation dependence of d33, Ps, and effective coefficient, Q33 .

-1 2 4 -2 Orientation d33, pmV Ps, C/cm , m C

001 100.3 105 0.47×10-2

011 62.2 87 0.48×10-2

111 20.5 68 0.23×10-2

The measured values of d33 are in good agreement with experimental results via other methods such as the voltage dependence of dielectric constant [119]. The d33 values of

(001) and (011)-oriented films are much higher than that of (111)-oriented film, which could be due to the spontaneous polarization along the (111) direction. For (111)- oriented film, only the 180o domain switching contributes to the dielectric response.

However, the electric filed is at an angle of ~54o with respect to the easy axis of the polarization for (001)-oriented film. Therefore, the polarization rotation might enhance the piezoelectric constant before the 180o domain reversal occurs. The related strain s to the square of the polarization, s QP2 , and the effective piezoelectric coefficient is defined as the derivative of the strain with respect to the electric field:

eff d33  2 QP (4.2) where Q is the electrostrictive coefficient, and  is the dielectric constant. The effective electrostrictive coefficients calculated as a function of crystalline orientations are also listed in Table 4.1. The piezoelectric constants d33 are smaller than for other perovskite ferroelectric (100~1000 pmV-1), whereas the calculated electrostrictive coeffients Q is not, which could be due to the small dielectric constant [46, 119] affecting the piezoelectricity as given in Equation 4.2.

Chapter 4 Pure BFO thin films

Figures 4.10(a) and 4.10(b) show an out of plane PFM amplitude and a phase image at room temperature respectively. By applying a dc bias to the tip during the scan, the polarization for all films can be reversibly switched. Here an external DC bias of -10V was applied to a 5×5 m2 area followed by +10 V applied to a 3×3 m2 area, with the change in contrast between the two regions marking the occurrence of ferroelectric switching, as shown in Figures 4.10(c) and 4.10(d). Furthermore, the PFM image indicates that the BFO film, which is imaged without a top electrode, is self-polarized

(imprinted) downward (purple).

Figure 4.10 Out of plane PFM images demonstrating switching of (001)-oriented films: (a) amplitude and (b) phase images before DC bias and (c) amplitude and (d) phase images after -10 and +10 V DC voltages.

To further examine the domain switching behavior of the film at nanoscale, a

Switching Spectroscopy Piezo Force Microscopy (SS-PFM) [120] was used to record the variation of local hysteresis loops at the centre of the grain boundaries of BFO film.

Chapter 4 Pure BFO thin films

During the acquisition of a hysteresis loop in SS-PFM, the tip is fixed at a given

location on the surface and the wave form Vtip V probe( t ) V ac cos t is applied to the tip [121]. As can be seen from Figures 4.11(a)-4.11(c), the tip position was denoted as

“A”. The corresponding PFM phase and amplitude are shown in Figures 4.11(d)-

4.11(g). Between each step of dc bias (state “ON”), the dc voltage is stepped back to zero (state “OFF”), and the PFM phase and amplitude loops are measured at both “ON” and “OFF” state to better understand the electric interaction between PFM tip and the film surfaces [103]. The phase hysteresis loops and amplitude butterfly loops are obtained under “OFF” state as shown in Figures 4.11(d), and (e), which shows the true ferroelectricity of the films. However, the asymmetric of the coercive field could be observed and might be due to the self-poling effect, which arises at the interface between BFO and SRO bottom electrode, suggesting that the downward state is more stable compared with the upward state. Meanwhile, the phase and amplitude loops under “ON” state were detected, the amplitude increases linearly with the applied voltage indicating that the response is dominated by electrostatic interactions [122].

Chapter 4 Pure BFO thin films

Figure 4.11 (a) AFM images (1×1 µm 2), (b) out-of-plane piezoresponse (OP-PFM) amplitude images, (c) OP-PFM phase images of (001)-oriented BFO thin films with the tip position was donated as A; (d) and (f) PFM phase voltage hysteresis loop and (e) and (g) amplitude voltage butterfly loop of (001)-oriented BFO thin films, in both “ON” (f, g) and “OFF” (d, e) states.

4.2.5 Magnetoelectric coupling of BiFeO3 thin films via PFM and MFM

Local magnetic response data acquired by MFM are shown in Figure 4.12(c) simultaneously with the AFM morphology and PFM amplitude images of Figure

4.12(a), and (b) respectively. The uniform magnetic contrast given by most of the pillars suggests that they are in a single-domain state, rather than the stripe magnetic domain. The magnetic probe was magnetized in such a direction that white (positive

Chapter 4 Pure BFO thin films phase, attractive interaction) stands for down magnetization and black (negative phase, repulsive interaction) corresponds to up magnetization [123].

Figure 4.12 (a) Topography (20 × 20 m2), (b) amplitude and (c) magnetic domain imaging of (001)-oriented BFO thin film using MFM.

In order to investigate the coexistence of magnetic and ferroelectric domain structure at microscopic level, PFM and MFM were measured simultaneously on the (001)- oriented BFO thin film. For MFM, a cobalt coated tip was used. The morphology and

MFM image before application of dc bias are presented in Figures 4.13(a) and 4.13(d).

Figures 4.13(b) and 4.13(e) show the results of simultaneous measurement of the topography and PFM phase image after +10V dc bias. The domain structures were switched by the dc electric field between the bottom electrode and the tip. To study the effects of ferroelectric polarization and electric field control of the magnetic domain structure, MFM image was scanned again after applying the dc bias as shown in

Figures 4.13(c) and 4.13(g). The topographies before and after electric field switching were not varied from the morphology images, whereas the MFM image was changed after the dc bias. The indexed magnetic domain structure in Figure 4.13(g) was changed via the electric field compared with the unswitched area in Figure 4.13(d).

The perpendicular antiferromagnetically coupled domains and the antiferromagnetic domain walls are clearly visualized by MFM for switching the ferroelectric domain zone [124]. Since the magnetic domains were changed via the electric field, it could be

Chapter 4 Pure BFO thin films emphasized that there is a significant magnetoelectric coupling between the ferroelectric ordering and antiferromagnetic ordering.

Figure 4.13 AFM and PFM images of (001) BFO thin film. (a), (d) Topography and magnetic domain imaging of BFO thin film detected using magnetic force microscopy (MFM). (b), (e) Topography and piezoresponse imaging of BFO thin film using piezoresponse force microscopy (PFM). The image shows the ferroelectric domain switching performed by the poling DC bias. (c), (f) Topography and domain imaging of BFO thin film using magnetic force microscopy after writing ferroelectric domains (e). This image gives information of both ferromagnetic and ferroelectric domain structures.

4.3 Polycrystalline BiFeO3 thin films

4.3.1 BFO on Pt/TiO2/SiO2/Si substrate

Figure 4.14 shows the XRD spectra of the polycrystalline BFO thin film grown on

Pt/TiO2/SiO2/Si substrate. It is revealed that the film exhibits polycrystalline perovskite structure with rhombohedral R3c symmetry [125] and free of impurity such as

Bi2Fe4O9. The out-of-plane lattice parameter calculated from the XRD result shows aBFO = 0.3972 nm. Comparing with a value of 0.3962 nm for the bulk BFO, it should be noted that a tensile strain (0.25%) for the pure BFO film. The film-substrate misfit strain may have impact on the spontaneous polarization of the films from the substrate materials. 48

Chapter 4 Pure BFO thin films

Figure 4.14 X-ray diffraction (XRD) spectra of BFO deposited on Pt substrate at 550 oC 50 mTorr.

Figure 4.15(a) shows the atomic force microscopy (AFM) morphologies of the as- received BFO thin film. The average grain size of the BFO film is approximately 150 nm, and the root mean square (RMS) roughness value is 10.88 nm for BFO. The out- of-plane PFM (OP-PFM) amplitude and phase images of the BFO film have been shown in Figures 4.15(b) and 4.15(c), respectively. The ferroelectric domain structures of the BFO exhibit a fractal growth habits, domain size similar to the grain size, and the polarization in up or down within each domain (Figure 4.15(c)). Domain patterns can develop with either {100} type-plane for 109o or {101} type-plane for 71o walls [126]. As shown from the XRD pattern, the polycrystalline films should be a mixture of all three types of domain walls.

Chapter 4 Pure BFO thin films

Figure 4.15 (a) AFM image (1×1 µm 2), (b) out-of-plane piezoresponse (OP-PFM) amplitude image, and (c) OP-PFM phase image of BFO thin films deposited on Pt/TiO2/SiO2/Si substrate.

Figure 4.16(a) shows the ferroelectric polarization versus electric field (P-E) curves of

2 the BFO thin films measured at 1 kHz, where remnant polarization Pr of 66 μC/cm is obtained. The values of Pr are similar to those reported previously for other BFO thin films[83], and the P-E loops of the films exhibit well-defined saturation and a good symmetrical rectangular shape with a coercive field of 700 kV/cm.

Figure 4.16 (a) Ferroelectric hysteresis loop and (b) leakage current density for BFO thin films deposited on Pt/TiO2/SiO2/Si substrate measured at room temperature.

The leakage current density (J) versus electric field (E) of the film is shown in Figure

4.16(b). Below 300kV/cm, both leakages increase quickly, after which leakage current densities increase slowly with the increase in the field. The decrease in leakage current of the BFO film compared with that reported [127] is believed to be attributed to the

Chapter 4 Pure BFO thin films

 single phase, smoother surface, reduction in VO and the decreased density of domain walls.

Figure 4.17 illustrates the dielectric constant, ε, and dielectric loss, tanδ, of the BFO films as a function of frequency. It is evident that ε and tanδ of the BFO film varied slightly from 160 to 130 and 0.015 ~ 0.43 with frequency respectively. At frequencies higher than 105 Hz, the increase in tanδ implies that the charge carriers in the films have not enough time to respond to the external electric field while the maximum electrical energy is transferred to the oscillation ions [128].

Figure 4.17 Dielectric properties of pure BFO thin films measured at room temperature as a function of frequency.

The frequency dispersion of dielectric constant, r and loss tangent tanof Pt/BFO/Pt capacitors as a function of temperature and frequency are shown in Figure 4.18. Strong frequency dispersion with a drastic decrease in the dielectric constant value at frequencies above 105 Hz can be observed. An abrupt increase in dielectric constant could be seen at 200 K, which could be due to spin-reorientation transition similar to those observed in most of the rare earth orthoferrites [129]. An associated Debye like 51

Chapter 4 Pure BFO thin films relaxation peak (a characteristic of intrinsic dipoles present in the system) in the dielectric constant data is observed at the vicinity of the step-like frequency [130]. The temperature-dependent frequency shift suggests that the relaxation could be of

Maxwell–Wagner type, which arises due to the presence of regions with different conductivities within the sample [130] or by the substrate induced stain in the film

[131].

Figure 4.18 Variation in (a) dielectric constant (r) and (b) loss tangent (tan of pure BFO thin films deposited on Pt/TiO2/SiO2/Si substrate measured as a function of temperature.

Figure 4.19 shows the magnetization hysteresis loops (M - H) of the BFO films when a maximum magnetic field of 5000 Oe is applied in-plane to the substrate surface at room temperature. The saturated magnetization of BFO is found to be 8.69 emu/cm3.

The magnetization of the BFO might be due to the balance between the antiparallel sublattice magnetization of Fe3+.

Chapter 4 Pure BFO thin films

Figure 4.19 Magnetic hysteresis loops for polycrystalline BFO thin films deposited on Pt/TiO2/SiO2/Si measured at room temperature.

The average piezoresponse of each PFM image obtained at distinct ac excitation amplitude is plotted for the polycrystalline BFO in Figure 4.20. The piezoresponse increases linearly with increasing ac voltage from 1.0 to 3.0 V, and the piezoelectric

-1 constant, d33, is calculated from the slopes of the plot closed to 30 pmV . Compared with the epitaxial BFO thin film, the piezoelectric constant for polycrystalline BFO grown on Pt/TiO2/SiO2/Si is smaller than those of the (001) and (011)-oriented materials, but higher than that of the (111)-oriented sample. This phenomenon could be because the polycrystalline BFO thin film has isotropic polarization direction and show a combination of the epitaxial one.

Chapter 4 Pure BFO thin films

Figure 4.20 Calculated average piezoresponse of PFM images obtained at each voltage step vs. ac voltage for the BFO thin films deposited on Pt/TiO2/SiO2/Si substrate.

The switching behaviors of the polycrystalline film at nanoscale were investigated via a SS-PFM [120]. The tip position is denoted as “A”, as shown in Figure 4.21(a). The corresponding PFM phase and amplitude images are shown in Figures 4.21(b) and

4.21(c). The PFM phase and amplitude loops are measured at both “ON” and “OFF” state to get insight on the electrical interaction between PFM tip and the film surface

[103]. The phase hysteresis loops and amplitude butterfly loops are obtained under

“OFF” state as shown in Figures 4.21(d) and 4.21(e) which show the true ferroelectricity of the films. However, the asymmetry of the coercive field has been observed and might be due to the self poling effect which arises at the interface between BFO and Pt bottom electrode. Meanwhile, the phase and amplitude loops under “ON” state were detected. The amplitude increases linearly with the applied voltage, indicating that the response is dominated by electrostatic interactions [122].

Compared with the epitaxial sample, the voltage coercive and amplitude of BFO/Pt are smaller than those of BFO/SRO, which could be due to the mixture of polar orientation in the polycrystalline film being made easier to switch the domain motions.

Chapter 4 Pure BFO thin films

Figure 4.21 (a) AFM image (1×1 µm 2), (b) out-of-plane piezoresponse (OP-PFM) amplitude image, and (c) OP-PFM phase image of polycrystalline BFO thin films with tip position was donated as “A”; (d) and (f) PFM phase voltage hysteresis loop and (e) and (g) amplitude voltage butterfly loop of polycrystalline BFO thin films deposited on Pt/TiO2/SiO2/Si, in both “ON” (f, g) and “OFF” (d, e) states.

The AFM morphology, PFM amplitude images, and local magnetic response data acquired by MFM are shown in Figures 4.22(a)-(c), respectively. The magnetic contrast given by most of the pillars suggests that they are in a single-domain state rather than the stripe magnetic domain. The magnetic domains are bigger than that of epitaxial BFO, which could be the antiferromagnetic properties, compressed by the epitaxial strain, were released by the anisotropic structure in the polycrystalline film.

Chapter 4 Pure BFO thin films

Figure 4.22 (a) Topography (20 × 20 m2), (b) amplitude, and (c) magnetic domain imaging of polycrystalline BFO thin film deposited on Pt/TiO2/SiO2/Si using MFM.

4.3.2 Size effect on the piezoelectric response of BFO on Pt/TiO2/SiO2/Si substrate

The magnitude of the internal strain is a function of the film thickness from the thickness-dependent relaxation of lattice-mismatch-induced stresses by the formation of misfit dislocation during film growth [132]. The thickness dependence of the crystal structure, ferroelectric domain structure and ferroelectric properties in polycrystalline

BFO have all been studied.

Films ranging from 10 to 160 nm have been grown on Pt/TiO2/SiO2/Si substrates by pulsed laser deposition at 550 oC in oxygen partial pressure of 50 mTorr. Structure, surface morphology, and local ferroelectric properties have been investigated by XRD and piezoresponse force microscopy (PFM). Meanwhile, the piezoelectric constant was investigated to the films with altering the film thickness via the PFM.

Figure 4.23 shows the XRD θ-2θ scans for the polycrystalline BFO thin films of different thickness on Pt/TiO2/SiO2/Si substrates. The XRD peaks were indexed using a distorted rhombohedral structure. All the BFO films have crystallized perovskite phase with mixed orientations of (001), (011), and (111). The intensity of the (110) peak increases with increasing thickness, indicating that perovskite distorted to the pseudocubic structure. The thickness of the thin films was determined to be about 150 nm by a surface profile (Alpha-step IQ, KLA-Tencor).

Chapter 4 Pure BFO thin films

Figure 4.23 XRD spectra for BFO films of varying thickness on Pt/TiO2/SiO2/Si substrate.

The out-of-plane lattice parameters, as determined from the (002) diffraction peaks, are plotted in Figure 4.24 as a function of thickness on Pt/TiO2/SiO2/Si substrate. Lattice misfit between BFO and Pt(aPt = 0.392nm) results in an in-plane isotropic compressive strain that causes an elongation in the out-of plane direction, and this strain gradually decreases with increasing film thickness [91]. From the out-of-plane lattice parameters, it should be noted that the BFO films are under compressive strain for thinner thickness (< 80 nm) in contrast to a tensile strain for the thicker ones (> 80 nm).

Chapter 4 Pure BFO thin films

Figure 4.24 Out of plane lattice parameter as a function of thickness of BFO thin films on Pt/TiO2/SiO2/Si substrate.

Figures 4.25 (a), (d) and (g) show respectively the surface morphologies of the 10, 80 and 160 nm-thick BFO films on Pt/TiO2/SiO2/Si substrate. The RMS values were approximately 6.4, 19.3 and 17.2 nm for 10, 80, and 160 nm-thick films respectively, predicting that the roughness of the film increases with increasing thickness and decreases with increasing thickness. Meanwhile, the piezoresponse amplitude images shown in Figures 4.25(b), (e), and (h). reveal that the majority of domains are limited by the grain boundary. The piezoelectrically inactive regions (> 50 pm) under the application of a 2 V amplitude voltage are 12.4%, 13.9% and 12.6% for the 10, 80 and

160nm-thick-film, respectively, predicting that the larger piezoelectric property corresponding to the 71 and 109o ferroelastic domain types has increased due to the increase in thickness to 80nm and then reduced in the 160nm-thick film. The increased piezoelectric properties for thickness below 80nm could be caused by the enhanced density of the 71o and 109o ferroelastic domain walls. Figures 4.25(c), (f) and (i) show the piezoresponse phase images. The majority of the domains of the BFO film are oriented with polarization upward (yellow bright), and only small portion of domains

Chapter 4 Pure BFO thin films exhibits polarization oriented downward (purple dark) (corresponding to smaller concentration of the 180º domain walls). The out-of-plane PFM phase images show a ratio close to 5.5: 4.5, 7.5: 2.5, and 8.0: 2.0 for upward and downward polarized domains in the 10, 80 and 160 nm respectively. Furthermore, the volume density of domain walls also decreased due to the increase in thickness. It is commonly believed that the characteristics of the domains and domain walls directly impact the ferroelectric switching behavior. Energy of the domain walls is proportional to the density of the walls, and conductivity of certain domain walls of the thinner BFO is much higher than that of the thicker one. This could be related to the density of misfit dislocation density ρ. A further increase in the film thickness will result in loss of coherency strains of the film and relaxation starts, for instance, by formation of misfit and/or threading dislocations or by breaking into different mutually relaxing domains

[133]. The appeareent misfit dislocations would impact the domain structure and further vary the piezoelectric properties.

Figure 4.25 (a), (d), and (g) AFM images (2×2 µm 2); (b), (e) and (h) out-of-plane piezoresponse (OP-PFM) amplitude images, and (c), (f) and (i) OP-PFM phase images for 10, 80, 160 nm BFO thin films on Pt/TiO2/SiO2/Si substrate respectively. Yellow (bright) and purple (dark) on the piezoresponse images correspond to negative (upward) and positive (downward) domains respectively.

Chapter 4 Pure BFO thin films

The piezoelectric constants, d33 of each thickness obtained from the corresponding

PFM amplitude image at distinct ac excitation amplitudes are plotted for the BFO film

-1 in Figure 4.26. d33 increases with increasing thickness from 40 pmV of a 10 nm film to 120 pmV-1 of a 80 nm film due to the relaxation of the compressive stress originated from the Pt bottom electrode [134]. The decrease with further increasing thickness to

-1 25 pmV at 160 nm thickness might be attributed to the misfit strain becoming tensile strain. The result here is a little different from that calculated theoretically, which describes that the piezoelectric constant d33 increases as the film thickness is increased.

However, there might exist a critical thickness (here is 80 nm) for the highest d33 since the experimental conditions could have introduced other parameters which have not been included in the calculation process, such the misfit dislocation, strain amplitude.

The existence of such a critical thickness is due to microstresses which may be the result of the periodic deviation of the actual misfit strain from the average misfit on the interface between the polydomain structure and the substrate [135].

Figure 4.26 d33 as a function of thickness of BFO film deposition on Pt/TiO2/SiO2/Si substrate.

Chapter 4 Pure BFO thin films

4.4 Effects of bottom electrode on switching behavior at nanoscale

One key factor affecting ferroelectric performances is the bottom electrode which should neither chemically interact with the multiferroic materials nor form a low permittivity compound at the interface [136]. Many metals such as Pt, Ru [137] and perovskite conducting electrodes such as SrRuO3(SRO) [56], YBa2Cu3O7-x(YBCO)

[138], LaNiO3(LNO) [139], and RuO2 [140] have been used as bottom electrodes.

Although there are many experimental results proving that bottom electrode materials can significantly affect the electrical and magnetic properties of the memory materials, systematic studies on the effect of the metal and oxide metallic electrode on multiferroic films are still limited.

LNO film has been chosen as bottom electrode and template. It is an isotropic n-type metallic oxide with a resistivity of 150–210 μΩcm at 300K. Conductive LNO (3.83 Å) bottom electrode with small lattice misfit (-3.3%) to the bulk BFO (3.96 Å) provides a high possibility for the preferred [141] (100)-textured growth of BFO. Meanwhile, oxide SrRuO3 (SRO) with a pseudocubic perovskite structure having the space group of Pbnm is of great interest due to its ferromagnetism and high electrical resistivity of

280 μΩcm [142] for future electronic device applications [143, 144]. SRO shows the lattice parameter of the pseudocubic cell as 3.93 Å at room temperature, similar to that of BFO (3.96 Å) with smaller lattice misfit (-0.7%) than that of LNO layer. The bottom electrode could therefore induce different strain states for the BFO thin film.

4.4.1 Structure and multiferroic properties of BFO on LaNiO3/Si and SrRuO3/Si substrates

Figures 4.27(a) and 4.27(b) show the X-ray diffraction spectra of BFO films deposited on LaNiO3 (LNO) and SrRuO3 (SRO)-coated Si substrates respectively. Usually, the

Chapter 4 Pure BFO thin films growth habits of the BFO thin films depend not only on the substrates but also on the deposition conditions. With the same deposition conditions for the BFO thin film, the growth habit only depends on the substrates [145]. LNO bottom electrode, owning a perovskite structure, provides a better possibility for the textured growth of the perovskite BiFeO3 than the polycrystalline SRO bottom electrode. The thickness of the

LNO and SRO layer is about 50 nm. The orientations of the BFO films, preferred 100 orientation, are inherited from those of LNO while that of BFO deposited on SRO layer shows a polycrystalline state. Such inheritance renders the LNO films to act as the seed layers. The out-of-plane lattice parameters calculated from the (002)- diffraction peak is approximately 3.878 Å, indicating a compressive strain (-2.0%) in the BFO film compared with the bulk BFO lattice parameter (3.96 Å) according to

misfit strain  calculated from equation:  ()/af a s a s , where af are the lattice parameter of the film, and as is the lattice parameter of the substrate. As for the films deposited on the SRO-coated Si substrate, the BFO phase could be indexed according to the pseudocubic structure (P4mm) as shown in Figure 4.26(b). The calculated out- of-plane lattice parameter is about 3.955 Å, suggesting a considerably smaller compressive strain (-0.16%) than that of the LNO buffered BFO thin film.

Figure 4.27 XRD spectra for polycrystalline BFO thin films deposited on (a) LaNiO3 (LNO) and (b) SrRuO3 (SRO)-coated Si substrate.

Chapter 4 Pure BFO thin films

Figures 4.28(a) and 4.28(d) show the surface morphologies of the BFO on LNO and

SRO-coated Si substrate respectively. It can be seen that the grains and root-mean- square (RMS) roughness are respectively changed from the order of 80 nm and 2.43 nm for the BFO film on LNO to 100 nm and 6.53 nm for the BFO film on SRO, indicating a crystallite structure with clearly resolved morphological features on both electrodes. LNO bottom electrode could induce a smoother morphology and decrease the grains size compared with the SRO bottom electrode, which could be due to LNO layer having a preferred (100) texture. The substrate orientation could control the domain variants [146]. The resolved piezoresponse amplitude images of the BFO grown on LNO and SRO bottom electrode as shown in Figures 4.27(b) and 4.27(e), reveal that most domains are limited by the grain boundary that raring lower piezoelectric properties, attributed to their higher resistivity than that of the grains inside. In addition, the BFO film on the SRO bottom electrode shows piezoelectrically inactive regions (> 50 pm, 20.7%) less than that of BFO film on LNO bottom electrode

(>60 pm, 47.5%) with a 1 V amplitude voltage applied. This predicts that the BFO film grown on LNO bottom electrode has a larger piezoelectricity property corresponding to the 71 and 109o ferroelastic domain types due to its (100)-textured growth behavior.

Figures 4.27(c) and 4.27(f) show the piezoresponse phase images. The majority of the domains of the BFO film grown on LNO are oriented with polarization downward

(purple dark, 86.1%), and only a small portion of domains exhibits polarization oriented upward (yellow bright, corresponding to smaller concentration of the 180º domain walls). In contrast, the phase image of BFO grown on SRO bottom electrode shows a ratio close to 1: 4 for upward and downward polarized domains. Furthermore,

BFO film deposited on LNO bottom electrode has a lower volume density of domain walls than that of BFO film grown on SRO bottom electrode. The characteristics of the

Chapter 4 Pure BFO thin films domains and domain walls directly impact the ferroelectric switching behavior, and energy of the domain walls is proportional to the density of walls. Conductivity of certain domain walls of BFO is much higher than that of the domains themselves, and

180º domain walls are the most conductive ones [46]. Therefore, BFO film grown on

LNO bottom electrode should be expected to have a lower leakage current than that of

BFO film grown on SRO bottom electrode. Furthermore, film-substrate misfit strain may result in a drastically different thermodynamic stability of two parallel domain walls with the same orientation, and thus the domain stability and domain-wall motion in BFO films [147]. As mentioned above, the misfit strain induced by LNO bottom electrode in BFO film is larger than that of SRO bottom done. The domain switching behavior of (001)-textured BFO grown on LNO bottom electrode might be similar to

(001)-epitaxial BFO thin film on STO single crystal substrate, where there are eight possible ferroelectric polarization directions corresponding to the four structural variants of the rhombohedral phase [146]. However, for the randomly oriented BFO thin film deposited on SRO bottom electrode, the possible ferroelectric polarization directions should be reduced due to the isotropic crystal structure. Therefore, the energy to drive the domain motion in the BFO grown on LNO bottom electrode should be smaller than that for BFO grown on SRO bottom electrode, and the polarization should be higher in BFO grown on LNO bottom electrode than that of BFO grown on the SRO bottom electrode. To further investigate the ferroelectric properties, macroscopic ferroelectric test has been performed as follows.

Chapter 4 Pure BFO thin films

Figure 4.28 (a) and (d) AFM images (1×1 µm 2), (b) and (e) out-of-plane piezoresponse (OP-PFM) amplitude images, and (c) and (f) OP-PFM phase images for BFO thin films on LNO and SRO-coated Si substrates, respectively. Yellow (bright) and purple (dark) on the piezoresponse image correspond to negative (upward) and positive (downward) domains respectively.

Figure 4.29(a) shows the ferroelectric polarization versus electric field (P-E) curves of the BFO films deposited on LNO and SRO bottom electrodes measured at 1 kHz with a Pt top electrode. The P-E loops of the films show well-defined saturated rectangular shapes. BFO film deposited on the LNO bottom electrode shows the remanent

2 polarization, Pr, to be 53.2 C/cm , whereas that of the BFO film grown on SRO

2 bottom electrode is ~ 40C/cm . Meanwhile the coercive field, Ec, is slightly decreased via LNO bottom electrode comparing with that of SRO bottom electrode. This means that LNO as an oxide electrode for BFO thin film provides a better interface between the ferroelectric thin film and the electrode in terms of reducing leakage current [145].

Note that the P-E loops measured for the two samples are shifted toward the positive field. The shift may be caused by the existence of an ultrathin interfacial layer which results in an internal bias due to the accumulation of free charges at the interface between the ferroelectric and the nonswitching thin layer [148]. The different bottom and top electrodes produce different interfacial layers close to the top and bottom

Chapter 4 Pure BFO thin films electrodes, which causes the thin film capacitor to have an asymmetric structure and thus leads to asymmetric P-E loops [145]. This was however not observed for the BFO thin film with Pt as both top and bottom electrodes mentioned previously.

Figure 4.29(b) shows the leakage current density, (J), versus electric field, (E), of both films. The BFO film with LNO bottom electrode shows lower leakage current in the electric field, < 200 kV/cm, than that of the BFO film with SRO bottom electrode. It is known that the 109o domain walls were the domain leakage path rather than the matrix and that they prevent the complete ferroelectric switching of (001) BFO domains, whereas the intrinsic ferroelectric properties of BFO can be observed from two-variant

BFO films with 71o domain walls [149]. Thus, the BFO with LNO bottom electrode with an (001)-textured structure has a lower density of 109o domain walls than that of the BFO film with SRO bottom electrode. There is additionally, the possibility of formation of point defects, for example, oxygen vacancies, in both films. It is generally believed that LNO bottom electrode could be formed under a lower deposition temperature (> 550 oC) compared with that of the SRO bottom electrode (>

650 oC). Thus, the interface between ferroelectric films near the film/electrode interface could accommodate more oxygen vacancies at lower temperature for BFO film deposition (550 oC) and the following annealing process.

Chapter 4 Pure BFO thin films

Figure 4.29 (a) Ferroelectric hysteresis loops and (b) leakage current density for BFO thin films deposited on LNO and SRO-coated Si substrates measured at room temperature.

Figures 4.30(a) and 4.30(b) illustrate the dielectric constant, ε, and dielectric loss, tan, of the BFO films with LNO and SRO bottom electrode as a function of frequency respectively. It is evident that ε and tanδ of the BFO film with SRO bottom electrode varied largely from 900 to 170 and 0.95 ~0.1 with frequency respectively. The dielectric properties of BFO with SRO bottom electrode show dispersion like the

Debye relaxation behavior, which could be due to the large concentration of dipoles between the BFO and SRO interface. Furthermore, the dielectric anomalies could be due to other interfaces such as grain boundaries. In contrast, the BFO film with LNO bottom electrode changed slightly from 230 to 100 for dielectric constant and 0.3 to

0.2 dielectric loss, respectively. The BFO film with SRO bottom electrode has a higher

ε than that of the BFO film with LNO bottom electrode, which may be ascribed to the larger grain size in the BFO film with SRO bottom electrode due to the size effect

[150].

Chapter 4 Pure BFO thin films

Figure 4.30 (a) Dielectric constant (r) and (b) tangent loss (tan dependence on frequency of pure BFO thin films deposited on LNO and SRO coated Si substrate measured at room temperature. The piezoelectric thin film is grounded with a conducting layer, here is LNO and SRO during PFM measurement, and thus the surface displacements could be detected via the converse piezoactuation [118]. PFM data were observed for both samples by acquiring multiples of 5×5 m2 scan size images at a single location on each sample.

The overall intensity of the amplitude increases with increasing the ac voltage.

Histogram data obtained from each of the PFM images denote Gaussian distribution with single narrow peaks. The average piezoresponse of each PFM images observed at distinct ac excitation amplitudes is plotted for different orientations in Figure 4.31.

From the slopes of the plots that the piezoelectric constant, d33 is is calculated to be

66.8 and 348.2 pmV-1 for BFO thin film with SRO and LNO bottom electrode respectively. The distinct enhanced piezoelectric constant for BFO deposited on LNO bottom electrode could be attributed to that the (001)-textured crystal structure exhibited a larger d33 than the polycrystalline film, and higher compressive misfit strain at the BFO/LNO interface than BFO/SRO interface. Furthermore, the residual stresses are closely related to the crystallographic orientation of the films, thus the residual stresses in films with (001)-textured orientation (LNO) might be smaller than those in films with random orientation (SRO). This leads to a higher piezoelectric

Chapter 4 Pure BFO thin films response, lower dielectric dissipation and higher value of saturation polarization in

BFO film with LNO bottom electrode [151]. Meanwhile, the smaller grain size would also contribute to the higher piezoelectric of the BFO with LNO bottom electrode.

Figure 4.31 Calculated average piezoresponse of PFM images obtained at each voltage step vs ac voltage for the BFO thin films deposited on the LNO and SRO coated Si substrate.

The AFM morphology and local magnetic response data acquired by magnetic force microscopy (MFM) are shown in Figures 4.32(a) and 4.32(b), respectively. The magnetic contrast given by most of the pillars suggests that they are in a single-domain state, rather than the stripe magnetic domain. The magnetic domains are bigger in the

BFO film with SRO bottom electrode than that of the BFO with LNO bottom electrode, which could be the antiferromagnetic properties compressed by the misfit strain being released.

Chapter 4 Pure BFO thin films

Figure 4.32 (a) Topography (5×5 m2), (b) magnetic domain imaging of polycrystalline BFO thin film deposited on LNO/Si and SRO/Si substrate using magnetic force microscopy (MFM) respectively.

4.5 Effect of bottom electrode on the surface potential of polycrystalline BFO

Figures 4.33(a) and 4.33(d) show the morphologies of the BFO film deposited respectively on LNO and Pt bottom electrode. Figures 4.33(b) and 4.33(e) show the out-of plane PFM phase images after a series of dc bias was applied to the BFO with

LNO and Pt bottom electrode, respectively. The eleven 7.0 × 0.5 m2 rectangular areas were scanned by a contact mode with the dc bias ranging from -10 to 10 V with a 2 V step to the cantilever tip as shown in Figures 4.33(b) and 4.33(e). Films are almost entirely upward polarized in the as-grown state with the exception of small downward pointing domains due to electrostatic relaxation [117] and self-poling effect as a consequence of the bottom electrode [152]. Furthermore, sharp piezoelectric responses

Chapter 4 Pure BFO thin films with complete reversible switching were observed at +8, and +10 V for BFO with

LNO bottom electrode and -8, -10, and +10 V for BFO film with Pt bottom electrode, indicating that the coercive fields are larger for BFO film with LNO bottom electrode than that for Pt bottom electrode, in good agreement with the SS-PFM hysteresis loops discussed previously. Figures 4.33(c) and 4.33(f) indicate KPFM images over the biased areas with an ac voltage of 1 V to record the potential distribution on both BFO thin films. The white and black rectangle indicates the positive and negative surface potential compared with the unbiased surfaces, whereas the gray area suggests that magnitudes of ferroelectric polarization and screened charges are analogous [153].

While applying the dc bias, the charges are injected to the BFO films surfaces from the conductive tip to screen the polarization charges of the oriented dipoles [154].

Figure 4.33 (a) and (d) AFM images, (b) and (e) out-of-plane PFM phase image and (c) and (f) KPFM surface potential distribution of the area scanned with the applied biases from -10 ~ +10V with a 2V step to the cantilever tips for BFO deposited on Pt/TiO2/SiO2/Si and LNO/Si substrate respectively.

Chapter 4 Pure BFO thin films

The surface potential line profile as a function of applied biases for both BFO films are shown in Figure 4.34. The initial absolute surface potential values of negative biases are much larger than that of positive ones for both thin films. The asymmetry of the surface potential behaviors might be attributed to the presence of an internal built-in electric field near the BFO/electrode interface, different screen charge of as-grown rate and different amounts of trap sites [155, 156]. Meanwhile, in increasing the positive voltage, more electrons are injected and a saturated screen charge occurs at about +8V for BFO film with LNO bottom electrode. Further increasing the positive biased voltage to +10V, the amplitude of surface potentials decreases gradually due to the overscreen charges [157]. However, as for the BFO with Pt bottom electrode, the surface potential gets a saturated screen charges at about +4 V, only about half of that of BFO with LNO bottom electrode, which could be due to the higher work function of

Pt bottom electrode (5.4 eV) than that of LNO bottom electrode (4.3eV). For the negative biased regions, positive surface potentials are obtained due to hole trapping effect [153]. With decreasing biases to -8V, the highest surface potential was observed for both films, which could be because the positive screen charges are dominant over the positive polarization charges. The surface potential saturates as increase in biases applied would associate with the Coulombic repulsion between screen charges [156].

The surface potential was found to keep positive to 0.05V for BFO deposited on LNO bottom electrode at 0 V bias, which is less than that of reported ferroelectric

PbZr0.53Ti0.47O3 (1.0V) [158] and BiFe0.95Ca0.05O3 (0.5V) [153]. This reduced bias suggests that electron trapping is easier in BFO films. In contrast, BFO deposited on Pt bottom electrode presents a negative surface potential at -0.025 V at 0 V dc bias, which may be because the electrons transfer from the Pt bottom electrode is easier than that of the LNO bottom electrode. Hence, the oxide metallic bottom electrode such as the

Chapter 4 Pure BFO thin films

LNO, or SRO is better than Pt electrode to apply in data storage application since oxide metallic bottom electrode could be better in preventing electron trapping.

Figure 4.34 Surface potential line profiles obtained from KPFM images.

Chapter 5 Chemically modified BFO thin films

5.1 Introduction

Applications of BFO are still hindered by challenges such as serious electrical leakage, ferroelectric reliability, chemical fluctuation, large coercive field, and inhomogeneous magnetic spin structure [145, 159]. Great efforts have been devoted to improve the multiferroic properties of BFO via site-engineering concept to introduce chemical pressure in this system [125]. The “chemical pressure” to change the electronic or crystalline structure destroys the spin cycloid structure, vary the ferroelectric Curie

 temperature and eliminate the oxygen vacancies (VO ).

Lanthanum (La) was used to alter the electric and magnetic properties via various techniques. However, ferroelectric and ferromagnetic characterization fluctuated distinctly in many research reports. As for magnetization, the inhomogeneous magnetic spin structure of BFO films can be suppressed by La substitution to improve the weak ferromagnetic ordering [50].As switching behavior of domains and domain walls is of crucial importance for ferroelectric and magnetic properties, La-doping is expected to have effects on the switching of domains and domain walls in BFO films.

Raveau et al. [160] have demonstrated that whatever the nature of the antiferromagnetic data of the undoped perovskite, such as Ln0.4Ca0.6MnO3,

Pr0.5Sr0.5MnO3 or CaMnO3, ferromagnetism can be induced by Ru doping in a very effective manner which leads to colossal magneto resistance (CMR) properties.

However, few works have focused on Ru-doped BiFeO3 thin films, even the

theoretically expected effective moment 4  2.83  is the Bohr magneton . eff, Ru  , 

Comparing to Cr, Sc or Nb ions, it is believed that Ru can most effectively stabilize the

Chapter 5 Chemically modified BFO thin films ferromagnetic state [161], which may be attributed to the 4d metal character of Ru. It is known that the 3d electrons are more localized compared to 4d electrons, therefore, the exchange interaction between Fe-O-Ru could be favored much more than any other transition metals [162]. The positive influence of Ru doping on magnetic ordering could be interpreted based on the mixed valence states of Ru, because Ru has the ability to exhibit three different oxidation states (Ru3+, Ru4+ and Ru5+), while the low- spin configuration of Ru ions are suitable for a ferromagnetic coupling with Fe ions through the empty eg orbital. Therefore it is expected that through Ru doping in

BiFeO3 thin film, high ferromagnetic property can be achieved while still keeping the high ferroelectric properties.

Additionally, synthesis and characterization of codoped BFO thin films have been proven to be a useful way to improve ferroelectricity and magnetization simultaneously. (Mn, Ti), (La, Ni)-codoped films showed greatly reduced leakage current density and well saturated hysteresis loops [145, 163, 164]. Thus, (La, Ru)- codoped BFO thin films are expected to show the similar results. Meanwhile, other

ABO3 perovskite materials have also been introduced to form solid solutions with

BiFeO3. For example, ferroelectric and ferromagnetic properties were enhanced in

PbTiO3, PbZrO3, and BaTiO3 modified BiFeO3 ceramics and thin films [64-66]. The insertion of other ABO3 perovskite into BFO not only stabilizes the BFO perovskite phase but also forms a morphotropic phase boundary (MPB) with BFO due to their different crystal symmetries. Core-shell BFO–Pb (Zr0.52Ti0.48) O3 (PZT) composite films have been reported that insulating PZT can increase the ferroelectricity and significantly enhance the magnetization of BFO [165]. However, the effects of adding small doses of PZT to modify BFO have little been investigated. The structural distortion, leakage current, concentration of oxygen vacancies, Fe-O-Fe bond angle,

Chapter 5 Chemically modified BFO thin films and Bi-O bond lengths of BFO might be varied by PZT insertion. Thus, high ferromagnetic and ferroelectric properties could be achieved simultaneously through

PZT modification in BFO thin film.

Therefore, La, Ru-doped and (La, Ru) codoped BFO thin films should be studied. The possible causes for the changed ferroelectric and ferromagnetic characteristics should be understood via defect chemistry and valence effect. The codopping effecs on the switching of domains structure and domain walls density are also to be investigated.

Finally, the effects of Pb(Zr0.52Ti0.48)O3 (PZT) on the crystalline structure, surface topography, domain structure, ferroelectric and magnetic behavior of BiFeO3 films grown on Pt/TiO2/SiO2/Si substrates will also be investigated. The overall ferroelectric and magnetic properties have been studied for PZT modified BFO thin films.

Meanwhile, the deep insight to the enhancement of ferroelectric and ferromagnetic characteristics is discussed.

5.2 La doped BFO

Figure 5.1(a) shows the XRD spectra of the as-deposited BFO and BLFO thin films.

The film exhibits a single-phase perovskite structure with no secondary phase being observed. The XRD spectra also suggest that the films are polycrystalline with rhombohedral symmetry [125]. The diffraction pattern of the BLFO shows a relatively broad single peak at 2θ = 32.8° and no (104) reflection appears, indicating that the rhombohedral distortion is reduced toward the orthorhombic or tetragonal structure with La doping. There is a slight increase in lattice parameter after La doping because the ionic radii of La3+ (0.1032 nm) is slightly larger than that of Bi3+ (0.1030 nm) [166].

The out-of-plane lattice parameters obtained from the XRD spectra are aBFO =

0.3971nm and aBLFO = 0.3972nm. Note that BLFO film is under more tensile strain

(0.312%) than that of BFO film (0.262%) in comparison with bulk BFO (0.396 nm).

Chapter 5 Chemically modified BFO thin films

The tensile strain and the film-substrate misfit may be the cause for the increased tetragonality of the BLFO film. Moreover, the strain could change the domain behavior and further change the ferroic behavior and magnetic interaction [133].

Figure 5.1(b) shows a typical FE-SEM cross-sectional microstructure of the BFO thin film on Pt/TiO2/SiO2/Si substrate, suggesting that the BFO film has dense columnar grains with thickness of about 350 nm.

Figure. 5.1. (a) X-ray diffraction spectra for BFO and BLFO thin film, (b) a typical FE-SEM cross-sectional image of the film.

Figures 5.2(a) and 5.2(d) show the AFM morphologies of the as-received BFO and

BLFO thin film respectively. The average grain size of the BLFO film is approximately 200 nm bigger than that of the BFO. The presence of La dopant is

 effective in controlling the volatility of Bi atoms to suppress VO concentration

'''  according to Null VBi 3 2V O since bond energy of La-O bond (789.6 ± 8 kJ/mol) is stronger than that of Bi-O bond (337.2±12 kJ/mol), which also hinders the oxygen ions motion and consequently decreases grain growth rate [167, 168]. The root mean square (RMS) roughness value is 10.88 and 5.97 nm for BFO and BLFO respectively, indicating that La-doping can also smoothen the film surface morphology. The out-of-

Chapter 5 Chemically modified BFO thin films plane PFM (OP-PFM) amplitude and phase images of BFO and BLFO film are shown in Figures 5.2(b), 5.2(c) and 5.2(e), 5.2(f), respectively. The ferroelectric domain structures of BFO exhibit a fractal growth habits, domain size similar to the grain size, and the polarization in up or down within each domain (Figure 5.2(c)). Upon introducing La dopant, the domain structures of BLFO show more homogeneous domains and the domains are pinned at the ground boundaries (Figure 5.2(f)). The density of domain walls in BLFO is less than those in BFO, therefore it is expected to have a lower leakage current in BLFO since certain domain walls in BFO are much more conductive than the domains themselves [99]. There exists three types of domain walls, namely those that separate domains 71, 109, and 180o difference in polarization

[169]. Domain patterns can develop with either {100} type-plane for 109o or {101} type-plane for 71o walls [170]. As shown from the XRD pattern, the polycrystalline films should be a mixture of all three types of domain walls. The fraction of the different types of ferroelectric domain walls is under study. Meanwhile, the domain size in BLFO film is larger than that of BFO, which may have originated from the enhanced tensile strain by La doping, bigger grain size and even the misfit dislocation between the film and the substrate [133]. The domain structure and density of domain walls of BLFO also suggests that BFLO can undergo a higher electrical bias and easier to pole than that of BFO, or even a different magnetic property.

Chapter 5 Chemically modified BFO thin films

Figure. 5.2. (a), (d) AFM images, (b), (e) out-of-plane PFM amplitude and (c), (f) phase images for BFO and BLFO thin film(1 × 1 µm 2) respectively. The height data are the topography in the image, and the phase data (piezoresponse) are mapped on as color. Bright and dark indicate the upward and downward domain orientation in Figure 5.2(c) and 5.2(f).

The leakage current density (J) versus electric field (E) of the films is shown in Figure

5.3. Below 300kV/cm, leakages of both capacitors increase quickly, after which leakage current densities increase slowly with the increase in electric field. The BFO film shows a low leakage current only when the electric field is below 200 kV/cm, while the leakage of the BLFO film is nearly two orders of magnitude lower than that of the BFO film. Decrease in leakage current of the BLFO film is believed to be attributed to the single phase, smoother surface, reduction in V  and the decreased O , density of domain walls. The lower concentration of oxygen vacancies in BLFO film could restrain the strain relaxation such that the remaining strain is higher, which is in line with the XRD result.

Chapter 5 Chemically modified BFO thin films

Figure.5.3 Leakage current density of Pt/BFO/Pt and Pt/BLFO/Pt capacitors at room temperature.

Figure 5.4 shows the ferroelectric polarization versus electric field (P-E) curves of the

BFO and BLFO thin films measured at 1 kHz, where remnant polarization Pr of 66 and

2 102 μC/cm are obtained respectively. The values of Pr are similar to those reported previously for other BFO thin films [166]. The P-E loops of the films exhibit well- defined saturation and good symmetrical rectangular shapes due to the low leakage. A lower coercive field of 380 kV/cm for BLFO compared with 700 kV/cm for BFO could arise from the change in domain switching dynamics because the presence of La ions might break the symmetry between orientations and extend to other low states of symmetry [171]. An alternative mechanism for the low coercive filed could be associated with the domain wall density of BLFO film [172]. Furthermore, the 6s2 lone-pair orbital of Bi3+ is stereochemically active and responsible for the ferroelectric distortion, and the distortions induced by La doping are therefore likely to be caused by the turning off of the lone-pair activity.

Chapter 5 Chemically modified BFO thin films

Figure 5.4 Ferroelectric hysteresis loops of Pt/BFO/Pt and Pt/BLFO/Pt capacitors at room temperature.

Figure 5.5 illustrates good saturation characteristic of Pr and Ec of BLFO as a function of electric field. Pr and Ec at the maximum electric field of 1300 kV/cm were approximately 102 μC/cm2 and 370 kV/cm, comparable with those of (111)-oriented epitaxial BFO film on STO single crystal substrate.

Figure 5.5 Electric field dependence of remanent polarization and coercive field of Pt/BLFO/Pt capacitors at room temperature.

Chapter 5 Chemically modified BFO thin films

Figure 5.6 illustrates the dielectric constant, ε, and dielectric loss, tanδ, as a function of frequency. It is evident that ε and tanδ of the films vary slightly with frequency (102 ~

106 Hz). Dielectric constant varies in the range from 320 to 260 for the BLFO and from 220 to 150 for the BFO. It is obviously that La dopant can significantly enhance the dielectric constant from 190 to 290 at 1 kHz, whereas it only has little influence on tanδ , which may be due to the bigger grain size of the BLFO, and the low charged

 ''' '''  defects or defect dipole complexes of the BLFO, such as VO and VBi or ()'VVBi O .

When the frequency is higher than 105 Hz, a slight increase in tanδ is observed, implying that the charge carriers in the films cannot follow the external electric field and the maximum electrical energy is transferred to the oscillation ions [171].

Meanwhile, there is a noticeable increase in tanδ at low frequency which might be due to space charges in the interface of the films [56].

Figure 5.6 Dielectric properties of Pt/BFO/Pt and Pt/BLFO/Pt capacitors at room temperature.

Chapter 5 Chemically modified BFO thin films

Magnetization characteristics of both thin films measured at a maximum magnetic field of 5000 Oe applied parallel to the substrate surface at 300 K are shown in Figure

5.7. The magnetic hysteresis loops shows weak ferromagnetic ordering in nature with

3 average saturated magnetization (Ms) from Ms ≈ 8 emu/cm (corresponding to a

3 magnetization of about 0.055 µB/Fe) for the BFO to 15 emu/cm (0.102 µB/Fe) for the

BLFO thin film. Enhancement in magnetization could be attributed to several possible reasons: i) the spatial homogenization of spin arrangement [173], ii) the disturbed spin cycloid structure when La ions positions in Bi-site [159], iii) the formation of partial

Fe2+ due to Bi ions volatilization or oxygen vacancies, iv) varied canting angle of Fe-

O-Fe bond due the distortion of FeO6 octahedron caused by introduction of La ions

[174], or v) the increased tensile strain changes the balance between the antiferromagnetic and ferromagnetic interactions, such as the net magnetization or the ferromagnetism originates at the domain walls [133, 174].

Figure 5.7 Magnetic properties of BFO and BLFO thin film measured at room temperature.

Chapter 5 Chemically modified BFO thin films

5.3 Ru doped BFO

Figure 5.8(a) shows the XRD spectra of the BFO and BFRO thin films grown on

Pt/TiO2/SiO2/Si substrate. It is revealed that all the films exhibit polycrystalline perovskite structure with rhombohedral R3c symmetry [125] and free of impurity such as Bi2Fe4O9. The rhombohedral distortion is reduced toward an orthorhombic or tetragonal structure with Ru doping deduced from a broadened single peak at around

32.7° (FWHM from 0.465° for BFO to 0.513 for BFRO as shown in inset in Figure

5.8(a)). The out-of-plane lattice parameters calculated from the XRD results are aBFO =

0.3972 nm and aBFRO = 0.3960 nm. Slight lattice reduction of the BFRO film is observed in comparison with the BFO film, which is caused by the smaller ion radius of Ru4+ ion (0.062 nm) than that of Fe2+ (0.078 nm) and Fe3+ (0.065 nm) ions.

Compared to a value of 0.3962 nm for the bulk BFO, it should be noted that the BFRO film is under compressive strain (0.05%) in contrast to a tensile strain (0.25%) for the pure BFO film. The film-substrate misfit strain may affect the spontaneous polarization of the films, which is via the doping effect on the films rather than from the substrate materials. From the FESEM cross-sectional images of BFO and BFRO thin films, it can be seen that the BFRO film has smoother surface and interfaces, and the thicknesses of the BFO and BFRO layers are estimated to be 250 and 260 nm respectively.

Chapter 5 Chemically modified BFO thin films

Figure 5.8 (a) X-ray diffraction spectra of BFO and BFRO thin films with inset of magnified patterns showing a diffraction at 2θ = 31.76°; (b) and (c) cross-sectional FESEM images of BFO and BFRO respectively.

Figures 5.9(a) and 5.9(d) show the surface morphologies of the BFO and BFRO thin films respectively. It can be seen that the grains and the root-mean-square (RMS) roughness are respectively changed from the order of 100 nm and 10.7 nm for the BFO film to 300 nm and 6.4 nm for the BFRO film, revealing the crystallite structure of the films with clearly resolved morphological features and that Ru dopant can induce a smoother morphology and increase the grains size, correlating well to the FESEM results (not shown). The piezoresponse images present a much more complex variation of contrast that reflects the perplexing arrangement of domains in the ferroelectric films [175]. The resolved piezoresponse amplitude images of the BFO and BFRO films as shown in Figures 5.9(b) and 5.9(e), reveal that most domains are limited by the grain boundary. Quite often, the BFRO film shows piezoelectrically inactive regions (> 400 pm) than that of BFO film (>300 pm) under the application of a 500 mV amplitude voltage, predicting that the BFRO film has a larger piezoelectricity property corresponding to the 71 and 109o ferroelastic domain types. Figure 5.9(c) and

5.9(f) show the piezoresponse phase images. The majority of the domains of the BFRO

Chapter 5 Chemically modified BFO thin films film are oriented with polarization upward (yellow bright), and only small portion of domains exhibit polarization oriented downward (corresponding to smaller concentration of 180º domain walls). In contrast, the BFO phase image shows a ratio close to 6.5:3.5 for upward and downward polarized domains. Furthermore, the BFRO film has a lower volume density of domain walls than that of BFO film. It is commonly believed that the characteristics of the domains and domain walls directly affect the ferroelectric switching behavior. Energy of the domain walls is proportional to the density of walls, and conductivity of certain domain walls of BFO is much higher than that of the domains themselves, and the 180º domain walls are the most conductive ones [94, 99]. Therefore, the BFRO film should be expected to have a lower leakage current.

Figure 5.9. (a) and (d) AFM images (2×2 µm2), (b) and (e) out-of-plane piezoresponse amplitude images, and (c) and (f) piezoresponse phase images of BFO and BFRO thin films. Yellow (bright) and purple (dark) on the piezoresponse image correspond to positive (upward) and negative (downward) domains respectively.

Figure 5.10 shows the leakage current density, (J), versus electric field, (E), of both films. The BFO film shows lower leakage current in the electric field < 200 kV/cm, while the BFRO film exhibits lower leakage nearly one order of magnitude compared

Chapter 5 Chemically modified BFO thin films with that of the BFO film in the electric filed > 200 kV/cm. It is known that in BFO, Fe ions often exist in a mixed-valance Fe2+/3+ state due to introduction of oxygen vacancies. As the valence fluctuates, holes, h , in the BFO film will be created in order to compensate for the change in charge at a relatively high oxygen partial pressure, which can be expressed as

3 2  Fe  Fe  h (5.1)

At a low oxygen partial pressure, the change in the oxidation state will lead to the formation of oxygen vacancies according to

2 '   Null 2BiBi  2 Fe 3  5O O V O Fe (5.2)

Consequently, BFO often shows a large leakage current density. By doping tetravalance Ru4+ ions, the valance fluctuation of Fe2+/3+ can be compensated according to

4  2 ' 4   Ru 2BiBi  Fe 33  Ru  6O O Fe Fe (5.3)

Due to reduction in the concentration of oxygen vacancies caused by doping high valance Ru, the leakage current density is reduced.

Figure 5.10 Leakage current density for BFO and BFRO thin films measured at room temperature. 87

Chapter 5 Chemically modified BFO thin films

Figure 5.11 shows the ferroelectric polarization versus electric field (P-E) curves of the

BFO and BFRO films measured at 1 kHz. The P-E loops of the films show well- defined saturation and good symmetry rectangular shapes. The BFO film shows the

2 remnant polarization, Pr, to be as high as 64 C/cm while that for the BFRO film is 99

2 C/cm . Meanwhile the coercive field, Ec, is slightly decreased via Ru doping comparing with that of the BFO film. The enhanced ferroelectric properties of the

BFRO film are coincidental with their domain structure as shown in Figure 5.11 and the compressive file-substrate strain.

Figure 5.11 Ferroelectric hysteresis loops for BFO and BFRO thin films measured at room temperature.

Figure 5.12 illustrates the variation of dielectric constant, ε, with dielectric loss, tan, of the BFO and BFRO films as a function of frequency. It is evident that ε and tanδ of the Ru-doped BFO film varied slightly from 210 to 160 and 0.012 ~0.023 with frequency respectively. At frequencies higher than 105 Hz, the increase in tanδ implies that the charge carriers in the films do not have time to respond to the external electric

Chapter 5 Chemically modified BFO thin films field [75] and the maximum electrical energy is transferred to the oscillating ions [128].

The BFRO film has a higher ε than that of the BFO, which can be ascribed to the larger grain size in the BFRO film due to the size effect [176].

Figure 5.12 Dielectric properties for BFO and BFRO thin films measured at room temperature.

Figure 5.13 shows the current-time plot characteristics of BFO and BFRO films measured at a positive bias of 5 V. It is clear that the leakage current of BFO stabilizes immediately whereas the current relaxation for the BFRO film is very slow, and stabilizes only after about 80 ms. The time taken to stabilize leakage current is probably carrier transfer and domain dynamic dependent, such as the polarization bound charges at the surfaces of the ferroelectric layer, free charge carriers in the conducting electrodes, and domain wall motion [115]. For the BFRO film, since it has more complicated defect chemistry and larger domain and smaller concentration of domain walls than that of BFO film, it will need more time to be stabilized.

Chapter 5 Chemically modified BFO thin films

Figure 5.13 Leakage current and time relationship for BFO and BFRO thin films measured at room temperature.

Figure 5.14 shows the magnetization hysteresis loops (M - H) of BFO and BFRO films when a maximum magnetic field of 5000 Oe is applied in-plane to the substrate surface at room temperature, where the saturated magnetization of BFO and BFRO is

8.69 and 16.53 emu/cm3 respectively. Compared with the BFO film, an increase in magnetization in BFRO films is observed. The increase in magnetization of the BFRO might be due either to the balance between the antiparallel sublattice magnetization of

Fe3+ and a destroyed spin cycloid due to Ru ions substitution [46, 72], or an increase in canting angle [56].

Chapter 5 Chemically modified BFO thin films

Figure 5.14 Magnetic hysteresis loops for BFO and BFRO thin films measured at room temperature.

To further investigate and understand the role of Ru doping on the electrical and magnetic properties, the chemical state of Fe and Ru ions in the films were investigated via XPS analysis as shown in Figure 5.15. The Fe 2p3/2 peaks correspond to Fe2+ and Fe3+. By fitting the peaks for the valence state of Fe ions, the rations of

Fe2+/ Fe3+ in BFO and BFRO films are calculated as 1:2.5 and 1:5 respectively, indicating the presence of Fe2+ ions being less in the BFRO. This could also be an evidence for the decrease in oxygen vacancies which corresponds to the lower leakage current density in the BFRO film.

Figures 5.15(b) and 5.15(c) show the Ru XPS spectra in BFRO film. It is found that the Ru3d3/2 and Ru3p3/2 peaks are positioned at the binding energy of 285.1 and 464.0 eV respectively. The Lorentzian dividing peaks corresponding to 285.1 eV and 464 eV of Ru suggest the presence of two possible redox couples involving Ru4+/ Ru5+ with a

Chapter 5 Chemically modified BFO thin films ratio of about 1:2, and Ru3+/ Ru4+ with a ratio of 1:4. It may be argued that the variation of Ru valence in the BFRO film is presumably due to the comparable oxidation reduction potential (ORP) Ru4+/ Ru5+ (1.07 eV) and Fe2+/Fe3+ (0.77 eV). As

3+ 2+ 4+ 5+ a result, Fe gets reduced to Fe state, while Ru gets oxidized to Ru state, which is a source of ferromagnetic interaction between Ru and Fe ions in the BFRO film [177].

Additionally, part of Fe3+ is also formed as a result of the internal redox reaction

2 4  3  3  Fe Ru  Fe  Ru (5.4)

In the oxidizing conditions, the system tends to keep the oxidation state of Ru to 4+ and to reduce the sate of iron. The introduction of Ru5+ will induce ferromagnetic coupling and extra Fe2+ can also participate in the enhancement of ferromagnetism via

2 5  3  forming a electronic fluctuation along the “ Fe O  Ru  O  Fe ” bond [160].

Furthermore, an alternative mechanism of charge compensation of Fe reduction from

Fe3 to Fe2 may be ascribed to the higher Ru doping which induces a small fraction

 of RuFe defect. The Ru dopant will distort the long-range cycloidal spin arrangement of BFO, and enhance the magnetic moments via valence effect. Ru ions also promote ferromagnetism by developing ferromagnetic superexchange interactions with the surrounding Fe ions.

Chapter 5 Chemically modified BFO thin films

Figure 5.15. (a) XPS spectra of the Fe2p lines of BFO and BFRO films, the insets are the peak fitting simulations. (b) and (c) XPS spectra of Ru 3d and 3p showing variable oxidation state for the BFRO film.

5.4 La, Ru codoped BiFeO3

Figure 5.16 shows XRD spectra for BFO, BLFO, BFRO, and BLFRO films deposited on Pt/TiO2/SiO2/Si substrates. All films are polycrystalline with no second phases. The spectra were indexed by the distorted rhombohedral structure (R, R3c) for BFO, BLFO,

BFRO, and BLFRO films [58, 178] (Diffraction peaks in the films in the vicinity of 2

= 32º are shown in the left inset of Figure 5.16(a)). The (200) peaks near 46º are shown in the right inset of Figure 5.16(b), which become broadened and shifted to a higher angle for BLFO and lower angle for BFRO in comparison with BFO. However, the

Chapter 5 Chemically modified BFO thin films

(200) peak for BLFRO showed no shift which is attributed to the fact that the ionic radius of La3+ ion (0.1032nm) is a little bigger than that of Bi3+(0.1030 nm), while the ionic radius of Ru4+ ion (0.0620 nm) is smaller than that of Fe2+ (0.0780 nm) and Fe3+

(0.0650 nm). The codopping of La and Ru ions into BFO thin film shows no modification in lattice parameters, which might exclude the strain effect on the reduced leakage current in the film and enhanced magnetization, as discussed later.

Figure 5.16.XRD spectra of BFO, BLFO, BFRO and BLFRO thin films with inset of magnified patterns showing diffraction at (a) 2θ = 32° and (b) 46°.

Figure 5.17 depicts the Raman spectra in unpolarized configuration of the four samples.

Positive shift in the vibration (A1) at 220 cm-1 for all doped films could be attributed to structural distortion and the additional La and Ru result in stiffening of the Bi-O bonds with a possible change to a more covalent character [179]. BFO in normal Raman spectrum 13, 8, and 27 modes are expected for rhombohedral, tetragonal, and monoclinic symmetries respectively [109, 110]. According to the selection rule and fitting results, BFO is rhombohedral, while BLFO, BFRO, and BLFRO could be

Chapter 5 Chemically modified BFO thin films monoclinic[180]. Two wide bands appear at 620cm-1 and 450~550cm-1, and are assigned to symmetric and antisymmetric stretching of the basal oxygen ions of the

FeO6 octahedra associated with the Jahn-teller distortion [110, 180]. By comparing the integrated intensity of the peak at 150cm-1 (Bi-O vibration) with regard to the band appearing at 620 cm-1, it was found that the octahedral distortion is the strongest with the introduction of Ru ions into Fe sites with regard to La doping. This phenomenon could be attributed to the polar displacement of Bi3+ ions being far more sensitive to doping concentration than that of Fe cations [181]. Furthermore, the number of longer

Bi-O band is reduced via La doping in comparison with that of Ru doping since the

Raman frequency of Bi-O bond is inversely proportion to Bi-O bond length [179]. A strong Jahn-teller distortion on substitution of Ru could be attributed to the strong hybridization between Ru4d and O2p, and Ru doping increases the activation energy of formation of Jahn-teller polarons [182]. In this case, La and Ru should influence the magnetic properties of BLFRO thin film. The charge disproportion between Fe and Ru:

Fe3 Ru 4   Fe 2   Ru 5  , where the formation of Ru5+ should increase the Jahn-

Teller effect and the super-exchange interactions between the “Fe2+-O-Ru5+-Fe3+” network[183].

Chapter 5 Chemically modified BFO thin films

Figure 5.17 Raman spectra with inset of magnified patterns showing composition dependent shift in wave number between (a) 200 to 240 cm-1, and (b) 550 to 700 cm- 1for BFO, BLFO, BFRO and BLFRO thin films.

The surface morphologies of BFO and BFRO thin films are shown in Figures 5.18(a) and 5.18(e) respectively. The grain size and root-mean-square (RMS) roughness are changed from the order of 100 nm and 10.7 nm for the BFO film to 500 nm and 18.2 nm for the BLFRO film respectively, suggesting that La and Ru ions are able to produce a rougher morphology and increase the grain size greatly. The PFM images present a much more complex variation of contrast that reflects the perplexing arrangement of domains and local piezoelectric response in the ferroelectric films

[175]. Figures 5.18(b) and 5.18(f) show the resolved PFM amplitude images of the

BFO and BLFRO films. The BLFRO film shows piezoelectrically inactive regions (>

300 pm) more than that of BFO film (>75 pm) under the application of 1V amplitude voltage, predicting that the BLFRO film owns a larger piezoelectricity property caused by the 71 and 109o ferroelastic domain types. Figures 5.18(c) and (g) show the PFM phase images. Most domains exhibit fractal characteristic and are limited by the grain 96

Chapter 5 Chemically modified BFO thin films boundary in BFO, while large domains of micron size are obtained in BLFRO. The majority of the domains of the BLFRO film are oriented with polarization downward

(yellow bright in Fig. 5.18(g)), and only a small portion of domains exhibits polarization oriented upward (white bright in Fig. 5.18(g)). This difference in domain behavior might be related to defect chemistry in BFO and BLFRO. For BFO film, the

'''  defect dipole complexes, ()'VVBi  O , induce an internal field and lead to a preferred direction of the ferroelectric domain. Thus the polarization is also randomly oriented up and down due to the randomly located nature of the defect-dipole complexes [184].

The large domain of BLFRO film might have arisen from the breaking of some complex of due to La and Ru doping. The characteristics of the domains and domain walls directly affect the ferroelectric switching behavior [2]. Energy of the domain walls is proportional to the density of the walls, and conductivity of certain domain walls of the BFO is much higher than that of the domains themselves. A tapping cantilever with cobalt coated tip was used for MFM for BFO and BLFRO, as shown in the Figs 5.18(d) and 5.18(h), where the magnetic images dark and bright regions are considered as magnetic domains. However, the MFM image in BFO shows an irregular magnetic domain structure, while that of BFLRO film shows a

“labyrinthine” domain structure [185], suggesting that the codopping of La and Ru induces a ferromagnetic spontaneous magnetization indicating that macroscopically, the magnetic properties of BLFRO are enhanced.

Chapter 5 Chemically modified BFO thin films

Figure 5.18 (a) and (e) AFM images (2×2 µm2), (b) and (f) out-of-plane piezoresponse amplitude images, (c) and (g) piezoresponse phase images, and (d) and (h) magnetic domain images (20×20 µm 2) via MFM of BFO and BFRO thin films. Purple (dark) and yellow (bright) on the piezoresponse image correspond to positive (upward) and negative (downward) domains respectively.

The chemical state of Fe ions in the films via XPS analysis are shown in Figure 5.19(a), where the mixed Fe valence state may be the main cause of the leakage current in BFO

2+ 3+ system [186]. The Fe 2p3/2 peaks correspond to Fe and Fe . By fitting the peaks for the valence state of Fe ions, the rations of Fe2+/ Fe3+ in the BFO and BLFRO films could be calculated as 1: 3 and 1: 6 respectively, indicating the presence of Fe2+ ions is less in BFRO. This could also be an evidence for the decrease in oxygen vacancies and the lower leakage current density in the BLFRO film. Furthermore, the alternation in

Fe 2p orbit spectra also predicts the influence on Fe site by the doping of Ru4+, which might affect the binding state of Fe-O-Fe/Ru bond in the Fe–O octahedral [183]. The rotation of the oxygen octahedral is suggested to play an important role in the origin of the weak ferromagnetism in BFO [187], and consequently the magnetization of the doped-BFO films was affected somewhat [63].

Chapter 5 Chemically modified BFO thin films

Figure 5.19 (a) XPS spectra of the Fe2p lines of BFO and BLFRO films, the insets are the peak fitting simulations and (b) XPS spectra of La3d5 lines and fitting simulation of BLFO and BLFRO thin film.

In addition, the binding energy obtained for the La3d5 spectra of BLFO and BLFRO are shown in Figure 5.19(b). The BLFO film has a nearly slight line indicating that

La3+ ions prefer to replace Bi-sites. However, the observed double-peak structure of spin-orbit split component (j = 5/2) in BLFRO film, where peak and satellite show a comparable intensity, is typical for oxygen-rich La(III) species[188]. The splits of XPS

La3d5 might have originated from the La3+ occupying the Fe-site with the exception of

3+ Bi-sites, like the La ions occupying the W-site in PbWO4 compound [189]. The

3+ substitution of La in the Fe-site could enhance the distortion of FeO6 octahedral and destroy the cycloid spin structure, which could also enhance the magnetization of the

BLFRO film as shown below.

Figure 5.20(a) shows the magnetization hysteresis loops (M - H) of BFO and BLFRO films. A maximum magnetic field of 5000 Oe is applied in-plane to the substrate surface at room temperature. The saturated magnetization (Ms) is 8.69 and 17.5

3 3 emu/cm and the remnant magnetization (Mr) is 0.75 and 6.13emu/cm for BFO and

Chapter 5 Chemically modified BFO thin films

BLFRO respectively. The improvement in magnetization in the BLFRO films is observed as expected, which might be ascribed to the breakdown of the balance between the antiparallel sublattice magnetization of Fe3+ due to partial La substitution

[190], destroyed spin cycloid and homogeneous spin structure [191] via the introduction of Ru into Fe sites (PRL), alerted canting angle of Fe-O-Fe/Ru bond due to the Jahn-teller distortion [110] , and even the variation in the balance between the antiferromagnetic and ferromagnetic interactions via Ru and La ions codopping.

Figure 5.20 (a) Magnetic hysteresis loop (M-H), (b) polarization-electric field hysteresis loops of BLFRO thin film as a functions of electric field recorded in the 100~350 K temperature range.

The ferroelectric hysteresis loops (P-E) of the BFO and BLFRO films measured at 1 kHz are shown in Figure 5.20(b). The P-E loops of the films show well-defined saturation and no contribution from the freely movable charges[178]. The BFO film

2 shows the remnant polarization (Pr) of 57.6 C/cm while that of the BLFRO film is

92.5 C/cm2 at an applied electrical field of 1000 kV/cm. Meanwhile, the coercive field, Ec, is slightly decreased via La and Ru doping comparing with BFO film, This

 phenomenon could be associated with the gradual release of Vo from the defect

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Chapter 5 Chemically modified BFO thin films

'''  complexes in the BLFRO[192], such as ()'VVBi  O since partial defect complexes might align along the polarization orientation. Thus the driving force for domain switching in BLFRO film should be weaker than that for BFO film due to the lower content of defect complex[192]. The enhanced ferroelectric properties and reduced coercive field of the BLFRO film are also coincident with their defect atmosphere and

 domain structure, such as the compensated oxygen vacancies (Vo ), restrained varying valence of Fe ions due to Ru valence effect, structural distortion raised from doping effect and even the different domains and domain walls behavior [193].

5.5 Pb(Zr0.52Ti0.48)O3 (PZT) codoped BFO

Figure 5.21(a) shows the XRD spectra of only pure and PZT-modified BFO thin films grown on Pt/TiO2/SiO2/Si substrates. The pure BFO film exhibits a polycrystalline perovskite structure with rhombohedral (R) R3c symmetry[58] and free of impurity

[62]. However, the PZT modified BFO thin film shows coexistence of rhombohedral

(R, R3c) and tetragonal (T, P4mm) symmetry, since 110 reflections are split into 110R and 110T. The rhombohedral distortion is reduced toward the tetragonal structure with increasing PZT concentration deduced from the broadened peaks in the vicinity of 2

= 32.34° (shown in the inset of Figure 5.21(a)). The out-of-plane lattice parameters calculated from the XRD results are aBFO = 3.955 Å, aPZT2 mol % = 3.964 Å and aPZT 5mol%

= 3.974 Å. The increased lattice parameters via PZT substitution could have stemed from that the radius of Zr4+ (0.72 Å) and Ti4+ (0.68 Å) being larger than that of the Fe3+

(0.64 Å) and the radius of Pb2+ (1.49 Å) being greater than that of the Bi3+ (1.30 Å).

The lattice parameters calculated from the diffraction data indicate that the

101

Chapter 5 Chemically modified BFO thin films tetragonalities (ca / 1) are 3.1% for PZT2%BFO and 3.56% for PZT5%BFO films.

The increase in tetragonality might be due to the shorter bond lengths for the first

Ti(Fe)-O (1.67Å) in PZT modified BFO as compared to that for the first Fe-O (1.953

Å) in pure BFO [66]. Figure 5.21(b) shows the FESEM cross-sectional images of the

PZT 2% modified BFO thin film. The modified films have smoother surface and interfaces, and the thickness of the films is estimated to be 150 nm corresponding to the surface profile results.

Raman spectra analysis for the structure of the modified films was performed to gain some insight into how the BFO symmetry changes. Raman spectra in unpolarized configuration of the three samples measured at room temperature are shown in Figure

5.21(c). The Raman active-modes for the rhombohedral pure BFO with R3c symmetry

are summarized using the irreducible representation Raman, R 3 c 49AE 1  [58, 110].

Two sharp bands are visible at around 140 and 170 cm-1 in PZT modified BFO films, but the intensity is larger than that of pure BFO while the intensity of the peak at 170 cm-1 increases with the concentration of PZT. This difference reflects the presence of regions with different symmetries in the modified film [195], with coexisting tetragonal and rhombohedral in good agreement with the average structure of the films characterized by XRD. Thus, the selection rules for the Raman active-modes of the

PZT modified BFO films could be described as Raman, P 4 mm 34ABE 1  1  . The Raman spectra between 120 and 250 cm-1 are attributed to Bi-O vibrations [111], which is because the Bi ions substituted by Pb ions have changed the number of longer Bi-O bonds that produce this band. The positive shift in the vibration mode (A1) situated at

220 cm-1 (shown in the inset of Figure 5.21(c) and Figure 5.21(d) position shift) is sensitive to the concentration of PZT, leading to some stiffening of the Bi-O bonds and increased dispersion in Bi-O bond lengths [181]. Two wide bands appear at around

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Chapter 5 Chemically modified BFO thin films

480 and 620 cm-1 in the PZT modified BFO films may be which ascribed to the Jahn-

Teller distortions of the FeO6 (and perhaps Fe/Zr/TiO6) octahedra [111]. The intensity of Figure 5.21(d) indicates the integrated intensity of the most intense peak in the spectra of the films at 154 cm-1 with regard to the band appearing at 620cm-1 plotted as a function of PZT concentration x. Octahedral distortions are the greatest for the

PZT5%BFO film while the Bi-O vibration is the most intense in pure BFO film, which could be because the hybridization between Bi and O was enhanced via Pb insertion.

Figure 5.21(a) XRD spectra of the pure PZT2% and PZT5%BFO thin films with inset of magnified patterns showing diffraction at 2θ = 32.34°; (b) typical cross-sectional FESEM images of PZT 2mol% modified BFO film; (c) Raman spectra of pure PZT 2% and PZT 5%BFO thin films with inset of magnified patterns showing composition dependent shift in wave number between 200 to 240 cm-1; and (d) the peak intensity -1 ratio I620/I154 and 220 cm line peak position versus PZT mol percent. Solid lines are guides for eyes only.

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Chapter 5 Chemically modified BFO thin films

Figures 5.22(a), (b) and (c) show the AFM images of the pure, PZT2% and

PZT5%BFO thin films respectively. It can be seen that the grain sizes are of the order of 100, 200, and 240 nm and RMS roughness are 10.88, 9.81, and 7.31nm respectively for the pure, PZT2% and PZT5%BFO thin films. This behavior indicates the crystallite structure of the films with clearly resolved morphological features and that PZT can induce a smoother morphology as well as increase the grains size. This might be mainly caused by a rise in crystallization temperature due to the chemical bond

o strength of Ti-O and Zr-O because the melting temperature of TiO2 (1750 C) and

o o ZrO2 (2677 C) are higher than that of Fe2O3 (1570 C) [62].

The piezoresponse images present a much more complex variation of contrast that reflects the arrangement of domains in the ferroelectric films [175]. The OP-PFM amplitude images of the pure, PZT2% and PZT5%BFO films are shown in Figures

5.22(d), (e) and (f). The domains are limited by the grain boundaries. Some of the grain boundaries act as barriers limiting the propagation during the poling process. The piezoelectrically inactive regions (corresponding to the 71 and 109o ferroelastic domain walls) in the amplitude images predict that the PZT modified BFO films have larger piezoelectricity property than that of pure BFO film. Figures 5.22(g), (h) and (i) show the OP-PFM phase images. The majority of the domains of the PZT2% and

PZT5% BFO films are oriented with polarization upward (purple dark), and only small portion of domains exhibits polarization oriented downward (corresponding to smaller concentration of the 180º domain walls), and the ratio for dark to bright area is 9.5:0.5 and 7.5:2.5 respectively. In contrast, the BFO phase image shows a ratio close to

4.5:5.5 for upward and downward polarized domains. A lower volume density of 180o domain walls than that of the BFO film exists in the PZT modified BFO thin films. It is believed that the characteristics of the domain structures and domain walls directly 104

Chapter 5 Chemically modified BFO thin films affect the ferroelectric switching behavior [2]. Therefore, the PZT modified BFO thin films should be expected to have a lower leakage current and higher ferroelectric polarization.

Figure 5.22 (a), (b), and (c) AFM images (2×2 µm2), (d), (e) and (f) out-of-plane piezoresponse (OP-PFM) amplitude images, and (g), (h) and (i) OP-PFM phase images of pure, PZT2% and PZT5%BFO thin films respectively. Yellow (bright) and purple (dark) on the piezoresponse image correspond to negative (downward) and positive (upward) domains respectively.

Figure 5.23 shows the magnetization hysteresis loops (M - H) of the pure, PZT2% and

PZT5%BFO films when a maximum magnetic field of 5000 Oe is applied in-plane to the substrate surface at room temperature. The saturated magnetization (Ms) of the pure and PZT2% and PZT5%BFO films is 8.5, 11.4 and 14.3 emu/cm3 respectively. The partly zoomed in M-H curves (shown in the inset in Figure 5.23) exhibit the remanent

3 magnetization (Mr) of 1.04, 2.42 and 4.28 emu/cm , suggesting a weak ferromagnetic nature of the films. Compared with pure BFO film, increases in magnetization in the

105

Chapter 5 Chemically modified BFO thin films

PZT modified films are observed. The increases in the magnetization of the PZT modified thin films might be due to several origins: i) Pb2+ ions doping in the Bi3+ sites can suppress the spiral spin structure of the BFO, and permits a canting of the antiferromagnetic sublattices to produce the weak ferromagnetism [59, 187]; ii)Zr4+ insertion in the antiferromagnetic ordering of the Fe3+ sites, which introduces a chemical pressure making the “-Fe-O-Fe-O-Fe-” bonds to “-Fe-O-Zr-O-Fe-“, and changes the average Fe-O-Fe bond angle, then enhances the ordering between Fe3+ and

Zr4+ and improves the magnetization value [194]; iii) substitution of the Fe site with nonmagnetic Ti4+ reduces the oxygen vacancies and the magnetization partially originated probably from the coexistence of Fe2+ and Fe3+ ions, and produces a large local structural distortion caused by the existence of Fe2+ [63]; and iv) the destroyed spin cycloid order originates from the Zr/Ti ion substation of Fe ions and Bi instituted by Pb ions.

Figure 5.23 Magnetic hysteresis loop (M-H) measured at room temperature for pure, PZT2% and PZT5%BFO films.

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Chapter 5 Chemically modified BFO thin films

Figure 5.24(a) shows the ferroelectric polarization versus electric field (P-E) curves of the films measured at 1 kHz. The P-E loops of the films show well-defined saturation and good symmetrical rectangular shapes. The BFO film shows a remnant polarization,

2 Pr, to be as high as 61.3 C/cm while that of PZT2% and PZT5%BFO films are 76.8 and 96.7 C/cm2 respectively. However, there is a slight imprint effect for pure BFO film shifts to the negative field corresponds to a preferred polarization “down” state and PZT2% and PZT5%BFO shift to the positive field corresponds to the “up” state, which may be explained by the stress induced by the film/Pt lattice mismatch or

 clamping, domain pinning induced by oxygen vacancies (VO ) and grain boundaries, or an internal bias caused by the accumulation of free charges at the interface between the films and the non-switching thin layer [195, 196]. Bi ions displacement effects significantly the spontaneous polarization of BFO while the polarization will decrease

''' when more Bi vacancies (VBi ) are formed [184]. Introducing Pb ions into the Bi site will recover the oxygen vacancies from the Bi evaporation since bond energy of Pb-O bond (3.9 eV) is the higher than that of Bi-O bond (1.7eV). Meanwhile, Pb substitution

'   '''  can form defect dipole complex (PbBiV O ) or ()'VVBi O defect dipole complexes, which could act as pinning points for domain motion. The enhanced ferroelectric properties of the PZT modified film are coincident with the increased rattling space of the oxygen octahedron due to larger Pb2+ ions [197]. The reduced coercive field of the

PZT modified BFO films may be partially ascribed to the increased tetragonalities from the rhombohedral symmetry, the compositions around the MPB, and higher concentration of defect dipole complexes.

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Chapter 5 Chemically modified BFO thin films

Figure 5.24 (a) Polarization-electric field hysteresis loops, (b) Leakage current density, (c) dielectric properties of pure, and (d) leakage current vs. time relationship of pure and PZT2% and PZT5%BFO thin films measured at room temperature.

Figure 5.24(b) shows the leakage current density (J) as a function of applied electric field (E) of all films. The BFO film displays a leakage current density of 2.29×10-3

A/cm2 at 300 kV/cm, while the PZT2%BFO and PZT5%BFO films exhibit lower leakage current density of 6.83×10-4 and 3.68 ×10-6 A/cm2 at 300kV/cm respectively.

The leakage current densities via PZT modification are lowered by nearly one to three orders of magnitude in comparison with pure BFO film. Note that the leakage current density in PZT5%BFO film decreases by nearly two orders of magnitude comparing with that of PZT2%BFO film. The bigger grain size, lower RMS of PZT5%BFO films and greater reduced oxygen vacancies could be attributed to reduced leakage current.

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Chapter 5 Chemically modified BFO thin films

Fe ions often show mixed-valance Fe2+/3+ state in BFO due to introduction of oxygen vacancies. At a low oxygen partial pressure, a change in oxidation state will lead to the formation of oxygen vacancies according to

2 '   Null 2BiBi  2 Fe 3  5O O V O Fe (5.5),

Thus, the BFO often shows a large leakage current density. By introducing Pb2+, Zr4+ and Ti4+ into BFO, the valance fluctuation of Fe2+/3+ can be compensated by the electrons from the donor dopants instead of oxygen vacancies according to

4X 2 ' 4   Zr 2BiBi  Fe 33  Zr  6O O Fe Fe (5.6)

4  2 ' 4   Ti 2BiBi  Fe 33  Ti  6O O Fe Fe (5.7)

The leakage current density is reduced due to reduction in the concentration of oxygen vacancies caused by doping high valance Ti and Zr, and the Pb ions substitution for Bi in BFO.

Figure 5.24(c) illustrates the frequency dependence of the dielectric constant, ε, and dielectric loss, tan, of the pure and PZT2% and PZT5%BFO films measured at room temperature. It is found that the PZT modified BFO films present an intense drop in the dielectric constant in the low frequency region between 100 and 1000 Hz. This decrease is possibly caused by space charge polarization or Maxwell–Wagner type interfacial polarization [198]. PZT modification may decrease tan, in particular in the lower frequency range due to the enhanced resistivity of the PZT modified BFO films.

The PZT-BFO films have higher ε than BFO, which could arise from the larger grain size in the PZT-BFO films due to size effect [176].

Figure 5.24(d) reveals the current-time characteristics of the pure, PZT2% and

PZT5%BFO films measured under + 5 V applied to the bottom Pt electrodes. It is clear that the leakage current of BFO stabilizes after about 50 ms whereas the current

109

Chapter 5 Chemically modified BFO thin films relaxation for the PZT 2% and PZT5% BFO films is very slow, and stabilizes only after about 200 and 400 ms. The longer time taken to stabilize leakage current for PZT modified BFO thin films may probably be due to the more complicated carriers transfer and domain dynamics dependence, such as polarization bound charges at the surfaces of the ferroelectric layer, free charge carriers in the conducting electrodes, larger domain structures and smaller concentration of domain walls [199]. As the PZT concentration is increased from 2mol% to 5mol%, the leakage current is reduced by nearly two orders of magnitude, which may have originated from the higher resistivity of PZT and lower oxygen vacancy density and the contribution of defect-dipole complexes.

5.6 Conclusions

The role of doping is discussed and the modified multiferroic properties are investigated.

1. Pure and La-doped BFO thin films are deposited on Pt/TiO2/SiO2/Si substrate by pulsed laser deposition. The films show a pure and polycrystalline phase with uniform surface morphologies. The domain growth habits are investigated via PFM, indicating that the domain structure is changed while domain wall density decreased by La doping. The leakage current density of BLFO decreases by two orders of magnitude compared with that of BFO. The BLFO thin film exhibits coexistence of greatly enhanced ferroelectric properties with large remanent polarization and ferromagnetic orderings in comparison with BFO film. The enhanced ferroelectric properties of La doped BFO are analyzed by domain engineering. High dielectric constant and low dielectric loss of BLFO ensure observation of good ferroelectric and piezoelectric behaviors.

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Chapter 5 Chemically modified BFO thin films

2. BFO and BFRO films have been fabricated on (111) Pt/TiO2/SiO2/Si substrate by pulsed laser deposition system. The films have a pure and polycrystalline phase with uniform surface morphologies. The electrical and magnetic properties are affected remarkably by the different nanoscale domain structure for both the films. The leakage current of the BFRO film decreases by nearly two orders of magnitude at high electric filed compared with that of the BFO film as expected. The BFRO film also reveals high dielectric constant and low dielectric loss. Defect chemistry due to introduction of

Ru ions is achieved. As a result, a better coupling between ferroelectric and ferromagnetic ordering has been obtained by doping Ru ions into Fe site of BFO.

3. The (La, Ru) codoped BFO films are fabricated on (111) Pt/TiO2/SiO2/Si substrate by pulsed laser deposition system. The XRD and Raman spectra indicate that the films have a pure and polycrystalline phase with phase transition from rhombohedral to monoclinic. The Ru ions induce strong Jahn-Teller distortions. The formation of oxygen vacancies and bivalent Fe ion have been constrained via Ru dopant and partial

La ions introduced into the Fe sites as evidenced by the reduced leakage current and the XPS results. The ferroelectric domain structure and wall density, the magnetic domains of the BFO film are significantly modified by the introduction of La and Ru

2 ions. The saturated ferroelectric hysteresis loops with Pr and Ec values of 92.5 C/cm and 320 kV/cm respectively at an electric field of 1000 kV/cm, while the Ms is 17.5

3 3 emu/cm and Mr 6.13 emu/cm for BLFRO film. The electrical and magnetic properties are affected by the different nanoscale domain structure for the films and the different defect chemistry atmospheres in BLFRO film. Consequently, a better coupling between ferroelectric and ferromagnetic ordering has been obtained via codoping La and Ru ions into BFO film.

111

Chapter 5 Chemically modified BFO thin films

4. Structural characterization by means of X-ray diffraction and Raman spectra analysis reveal the existence of a MPB region between tetragonal and rhombohedral symmetries in the multiferroic 2mol% and 5mol%PZT modified BFO films on (111)

Pt/TiO2/SiO2/Si substrate prepared using the pulsed laser deposition system. The PZT modified BFO films exhibit a pure and polycrystalline phase with uniform surface morphologies. The ferromagnetic properties are enhanced by effects such as Pb ions suppressing the spiral spin structure, Zr, Ti ions introduced chemical pressure changing the Fe-O-Fe bond angle, or even substituents causing destruction of a cycloidal spin structure. The ferroelectric properties are affected by the different nanoscale domain structure for the films. The leakage current of the PZT2%BFO and PZT5%BFO films decreases by about one and two orders of magnitude at high electric filed in comparison with that of the pure BFO film. The PZT modified BFO films also have higher dielectric constant and lower dielectric loss than BFO. As a result, a better coupling between ferroelectric and ferromagnetic ordering has been obtained by introducing PZT into BFO system.

112

Chapter 6 Leakage mechanism of BFO

Chapter 6 Leakage mechanisms of BFO

6.1 Introduction

Leakage current analysis in ferroelectric materials has been an important study for the use of ferroelectric materialism functional devices [200]. A number of leakage mechanisms have been presented to describe the possible conduction mechanism and can be divided into two main categories: bulk limited and interface limited [201].

The first mechanism is the bulk space-charge-limited current (SCLC) mechanism [202,

203], where

2 9r 0  E JSCLC  3 8d (6.1)

where μ is drift mobility of charge carrier, εr, permittivity of dielectric, ε0, permittivity of free space, E, applied electric field, d, thickness of dielectric and θ, the ratio of the total density of free electrons to the trapped electrons. The physics of this mechanism can be explained through charge injection effect [204]. At low voltages, the current is limited by thermionic emission from the cathode. As the voltage is increased, the charge injected increases and the space between the cathode and the anode becomes filled with charge. At high currents, it is the current limiting effect [201].

The second mechanism is the bulk limited Poole-Frenkel (PF) mechanism [201], resulted from positively charged, defect or impurity generated electron traps in the materials. The mechanism involves the consecutive hopping of charges between trap centers [204]. The interaction between the positively charged trap and electron gives rise to a Coulombic barrier. The ionization of the trap charges is a result of both thermal and field activation. The Poole-Frenkel (PF) mechanism may be written as

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Chapter 6 Leakage mechanism of BFO

E q qE Jcexp{  [I  ]} PF kT kT   0 r (6.2)

where EI is the ionization energy of the traps in the film and c is a constant.

The third mechanism is the interface limited Schottky effect arises from the Schottky barrier at the interface of the metal electrode and ferroelectric materials. Schottky emission effect is described by

 q qE J AT 2 exp{  [b  ]} Schottky kT kT 4  0 r (6.3)

where JSchottky is current density across a Schottky Barrier, A is the Richardson constant,

T is the temperature in Kelvin, q is the electron charge, b is the Schottky Barrier height, and k is the Bolzmann constant. This arises from the Schottky barrier at the interface of the metal electrode and dielectric films.

Finally, the fourth mechanism for leakage current is interface limited Fowler-

Nordheim (FN) tunneling behavior, which can be described by

3/2 2 Ci JFN  BE exp[ ] E (6.4)

where B and C are constants and i is the potential barrier height. The FN mechanism described the injection of charge carriers into an insulator layer from electrodes taken place by tunneling through an interfacial energy barrier.

Leakage currents can affect the long-term reliability of storage elements. The leakage currents can also be used as a scientific tool to understand the basic properties of the capacitor. A systematic study of the mechanism of conduction could provide a wealth

114

Chapter 6 Leakage mechanism of BFO of information about interface and bulk space charge [205]. The influences of parasitic phases on the properties of BFO epitaxial thin films have been investigated with results indicating that the growth windows of single-phase BFO films are narrow. In this study, the deposition temperature and oxygen pressure on the leakage current mechanisms in Pt/BFO/LNO capacitors are reported. In addition, the La and Ru codopping effect on leakage current is one of the main problems in the case of chemically modified BFO films. Hence, the conduction mechanism and potential barrier at film/electrode interface are investigated for the chemically modified thin film.

6.2 Deposition temperature dependent leakage mechanism

Figure 6.1 shows the leakage current density as a function of electrical field for the

BFO thin films at different growth temperatures. The difference in leakage current measured with positive and negative electric field is due to the asymmetric electrodes in Pt/BFO/LNO capacitors for 450 oC. However, the others are relatively symmetrical although different top and bottom electrodes. This could suggest that the leakage current mechanism is bulk limited and the electrode interfaces would have less influence [119]. It is evident that by increasing the growth temperature, the leakage current of the thin films increases except for that of film deposited at 500 oC.

The large leakage current in BFO thin films are often attributed to the existence of Fe2+

 ions and commensurated oxygen vacancies (VO ) in contrast to Pb based ferroelectrics such as PbTiO3 thin films in which volatilization of PbO is considered to be the

 primary cause. In addition toVO , leakage current is also affected by porosity, phase purity and crystallinity of the films. Therefore, BFO films grown at 500 oC having a homogeneous microstructure, low porosity and preferred orientation have the lowest leakage current.

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Chapter 6 Leakage mechanism of BFO

Figure 6.1. Typical J – E characteristics of Pt/BFO/LNO capacitors for different deposition temperatures at a fixed 50 mTorr oxygen pressure.

Figure 6.2 shows the log(J) vs log(E) plots for the BFO films deposited at different temperatures. At low electric field, the leakage current of all BFO films except for 600 oC appear linear with slopes near unity, indicating Ohmic conduction. According to

SCLC theory, the log-log plots of current density vs. voltage can be mathematically fitted with several linear segments with different gradients as shown in Figure 6.2.

With increasing electric field, electrons are injected into the film, and the density of the free electrons becomes greater than that of the thermally stimulated electrons. The slopes of the plots are also indicated in Figure 6.2. The film deposited at 550 oC shows

Ohmic conduction in all bias. Slopes of around 2show agreement with SCLC behavior, suggesting that discrete traps are embedded in the background of continuously distributed traps [206]. After the trap-filling process in all films except for film deposited at 550 oC, slopes of more than 2 correspond to trap-free conduction in the

SCLC.

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Chapter 6 Leakage mechanism of BFO

Figure 6.2 Log (J) vs log (E) plots at positive bias for the SCLC mechanism.

Figure 6.3 shows the plot of ln(J) vs E0.5 indicating a Schottky mechanism. A straight line fit to the data results in a relative dielectric constant close to the optical relative dielectric constant indicating the presence of an interface controlled Schottky effect.

This suggests that the Schottky mechanism could be ruled out for the films deposited at 450 oC

Figure 6.3 Log (J) vsE0.5 plots at positive bias for the Schottky mechanism.

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Chapter 6 Leakage mechanism of BFO

0.5 Figure 6.4 presents the plots of log(J/E) vs E for the films deposited at different temperatures. The fitting is reasonably good for a wide range of electric fields for films deposited at higher than 550 oC. Since the PF emission is a bulk limited process the emission of charge carriers trapped in the defect centers contribute to the conduction process. The trap centers could be distributed in the forbidden regions between the valence and conduction band of the BFO films. The carriers in the traps could be activated either thermally or electrically. Thus, it could be concluded that the trap centers are more in the film prepared at higher deposition temperature because more defect centers such as the oxygen vacancies, bismuth vacancies or the misfit dislocations could be found at the higher depiction temperature [200].

Figure 6.4 Log (J/E) vsE0.5 plots at positive bias for the PF mechanism.

Figure 6.5 shows the plot ln (J/E2) vs (1/E) at positive electric field for BFO films deposited at different deposition temperatures. The leakage current shows FN tunneling in high electric field and the onset electric field of FN tunneling decreasing with increasing deposition temperature except for the film deposited at 550 oC.

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Chapter 6 Leakage mechanism of BFO

Figure 6.5 Log (J/E2) vs1/E plots at positive bias for FN mechanism.

6.3 Oxygen pressure dependent leakage mechanism

Figure 6.6 shows the J-E characteristics of the BFO thin films fabricated at different oxygen pressure measured at room temperature. It can be seen that the leakage current density of the films decrease with the increase in partial oxygen pressure except for the film deposited at 50 mTorr. The understanding of the leakage mechanism of BFO deposited at different pressure and the asymmetric structural configuration dominated by the asymmetric electrode-material interface is crucial. The increased leakage current as the partial oxygen pressure is decreased could be related to the defect chemistry in the films. The oxygen vacancies could be enhanced at lower oxygen

Null2 BiXx  2( Fe2 )'  5 O  V pressure due to BiFe3 O O

At a relatively high oxygen partial pressure, holes, h , in the BFO films will be created to charge compensation due to Fe32 Fe  h  .

119

Chapter 6 Leakage mechanism of BFO

Figure 6.6 Leakage current density-electric field characteristics of Pt/BFO/LNO capacitor films deposited at different oxygen pressures at 500 oC

Figure 6.7 shows the log (J) vs log (E) plot for BFO films deposited at different partial oxygen pressures. At low electric field, the leakage current of all BFO films appears linear with slopes of nearly unity, predicating Ohmic contact behavior. The log-log plots of current density vs. voltage are mathematically fitted with several linear segments with different gradients according to SCLC theory as shown in Figure 6.7.

The slopes of the plots are also indicated in Figure 6.7. Film deposited at oxygen pressure smaller than 75 mTorr did not show plot close to 2, suggesting that no discrete traps are embedded in the background of continuously distributed traps [206].

Slopes more than 2 would correspond to trap-free conduction in the SCLC.

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Chapter 6 Leakage mechanism of BFO

Figure 6.7 Log (J) vs. log (E) for BFO films deposited at different oxygen pressures.

Figure 6.8 shows the plot of ln(J) vs E0.5 for Schottky mechanism at low oxygen pressure. A straight line fit indicates the presence of an interface controlled Schottky effect. This suggests that the Schottky mechanism could be ruled out for the films deposited at 75 mTorr.

Figure 6.8 Log (J) vsE0.5 plots at positive bias for Schottky mechanism for films deposited as a function of oxygen pressure.

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Chapter 6 Leakage mechanism of BFO

0.5 Figure 6.9 presents the plots of log(J/E) vs E for the films deposited at different oxygen pressures. The plots show a linear relation for electric field of 160 kV/cm. PF emission is a bulk limited process in which the emission of charge carriers trapped in the defect centers contributes to the conduction process. The optical dielectric constant

could be obtained for the slope of e e//opt 0 kT . The films deposited at low oxygen pressure show PF mechanism.

Figure 6.9 Log (J/E) vsE0.5 plots at positive bias for the PF mechanism.

Figure 6.10 shows in the plot of ln (J/E2) vs (1/E) at positive electric field for BFO films deposited at different oxygen pressures. The leakage current shows FN tunneling in high electric field. The film deposited at 75 mTorr shows the lowest onset of FN tunneling, indicating that the film is strongly controlled by the FN tunneling mechanism. It should be pointed out that the conduction mechanism is mainly dominated by Ohmic contact in the low electric field range.

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Chapter 6 Leakage mechanism of BFO

Figure 6.10 Log (J/E2) vs1/E plots at positive bias for the FN mechanism at different oxygen pressures.

6.4 Leakage mechanism in La and Ru codoped BFO

Figure 6.11(a) shows the diffraction spectra of the BLFRO thin film on Nb-STO substrate. The figure reveals a high quality epitaxial film that appears to be single phase and only (00l) diffraction peaks from BLFRO film and substrate. The out-of-

plane lattice parameter ( a f ) obtained from the (002) diffraction is about 4.057 Å, indicating an in-plane compressive misfit strain induced by the substrate

( f  as a f a s = -3.89%, where as is the substrate lattice parameter = 3.905 Å). The in-plane epitaxial arrangement of the films was assessed using pole figure measurement at a fixed 2scans of 110 reflections of the Nb-STO as shown in Figure

6.11(b). This clearly shows the fourfold symmetry along the <001> direction and

“cube-on-cube” epitaxial growth of BLFRO films on Nb-STO substrates [207]. The surface morphology of BLFRO film shown in Figure 6.11(c), suggests very smooth films (root mean square roughness of 0.907 nm) with atomic terraces present. Out-of- plane piezoelectric phase image is presented in Figure 6.11(d). The ferroelectric

123

Chapter 6 Leakage mechanism of BFO domains show a fractal characteristic with average domain size of ~ 60 nm, and the ferroelectric polarization prefers downwards (55%, dark purple) in the as-deposited film compared with upwards (45%, bright yellow) due to the self-poling effect of the bottom electrode [208].

Figure 6.11 (a) X-ray diffraction spectra for BLFRO thin film on a (001) Nb-STO substrate, (b) Pole figure image of the BLFRO thin film on (011) plane. (c) AFM image and (d) out-of-plane piezoresponse phase image of BLFRO thin film.

Figure 6.12(a) shows the ferroelectric hysteresis loops (P-E) of the BLFRO films measured at 1 kHz. The P-E loops of the films show well-defined saturation and no contribution from the freely movable charges[178]. The BLFRO film shows a remnant

2 polarization (Pr) of 70.4 C/cm and coercive field of 320 kV/cm. Figure 6.12(b) shows the I-V characteristics of epitaxial BLFRO film as function of electric field at various temperatures. The leakage current reduces monotonically with decreasing

124

Chapter 6 Leakage mechanism of BFO temperature, and is seen to be symmetrical between 100 and 200 K while asymmetrical between 250 and 350 K albeit with different bottom electrodes. The symmetrical plots may suggest that the leakage mechanism is bulk limited and the electrode interface could be less influenced [209] at a relatively low temperature. Figure 6.12(c) shows the frequency-dependent dielectric constant,r and dielectric loss, tan of BLFRO thin film measured at room temperature. The effective dielectric constant (10 kHz~1MHz) is about 150, which is higher than that of pure BFO ceramics and film reported [210,

211], while the low frequency dielectric constant is about 176. Figure 6.12(d) shows the temperature dependence of dielectric constant of BLFRO film between 100 and

400 K. The dielectric constant increases with increasing temperature due to the enhanced electrical conductivity and exhibites less frequency dielectric relaxational effects and no Maxwell-Wagner-type contribution to the permittivity [212].

Figure 6.12 (a) Polarization-electric field hysteresis loops at room temperature, (b) leakage current density as a function of temperature, (c) dielectric properties as a function of frequency, and (d) temperature dependence of dielectric constant measured at different frequencies of BLFRO thin films. 125

Chapter 6 Leakage mechanism of BFO

To understand the leakage mechanism, several models have been used to describe the carrier transport behavior in the ferroelectric film. Figure 6.13(a) shows J-E plots for

BLFRO film at low field (E < 400 kV/cm) at a temperature range between 100 and 350

K. The plots are found to fit well with the modified Langmuir-Child Law [213] described as

JEE= 2 (6.5) where the values of  and  as a function of temperature simulated from the fitting are depicted in the inset in Figure 6.13(a). Both and decrease as temperature is reduced, suggesting that both the Ohmic and space charge limited contributions (SCLC) decrease due to the reduced density of free carriers caused by carrier injection and volume-generated free carriers as temperature decreases [214]. However,  and at

100 K show a strange behavior which may be contributed by other mechanisms.

Figure 6.13(b) shows plots of ln(J/E) vs E1/2 at positive bias and at temperature range of 100~350 K at high field (E > 400 kV/cm). The plots present a linear relationship from 200 to 350 K. According to the mechanism of PF effect, which is associated with a thermal emission of electrons overcoming the Coulomb barrier between electrons and positively charged traps with the help of applied electric field, a higher electric field is required to activate the trapped electrons to produce an obvious contribution on

1/2 the conduction [213]. The inset in Figure 6.13(b) presents the  I as a function of E , indicating a zero-field trap ionization energy of 0.393 eV. The value might be associated with shallow oxygen vacancies (VVV,,   ), because the vacancy state O OO12 lies 0.2 ~ 0.4 eV below the conduction band in BFO, rather than the Fe2+ ions trap centre as the Fe2+ ions are likely trap center in BFO with = 0.65~0.8 eV [215, 216].

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Chapter 6 Leakage mechanism of BFO

It is expected that the introduction of Ru ions into the Fe-site could suppress Fe3+ valence fluctuation and the shallow trap level will be affected by temperature.

Figure 6.13(c) shows the log(J) vs. log(E) plots at the temperatures of 100, 140, and

170 K. The slopes of the plots are listed in the inset table of Figure 6.13(c). The slopes at low electric field is in the vicinity of 1, indicating Ohmic contact behavior. The slopes at higher electric field are around 8, which represent the trap filling process of

SCLC [217]. The defect chemistry could be as follows: i) at low T and E,

Ru4+substituted Fe3+ site releases electrons:

 X Ru O Ru3  e '  2 O (6.6) 2 Fe O ii) at high T and E, the vacancies captured release electrons and form trapped electron center (F-center) and further affects the shallow trap emission rate:

V e' V  , (6.7) OO21

V  e' OX (6.8) O1 o

The injection of charge carriers into a ferroelectric layer from electrodes may take place by tunneling through an interfacial energy barrier [218], which is named as

Fowler-Nordheim tunneling (FN). Fig 6.13(d) shows the plots of ln(J) vs 1/E, where a linear relation for 200 ~ 350 K at high electric field can be observed, indicating the formation of a partial depletion layer between the as-deposited thin film and the electrodes [219]. The onset electric field of FN tunneling decreases with elevated temperature suggesting a decrease in potential barrier height and increase in leakage current density with increasing temperature. Meanwhile, the leakage current is mainly controlled via Ohmic contact in the low electric field range over the whole temperature range. 127

Chapter 6 Leakage mechanism of BFO

Figure 6.13 (a) J-E curves of BLFRO film at low field, fitting according to modified Langmuir-Child law; inset presents the fitting parameters  and  as function of temperature, (b) Log (J/E) vs E1/2 curves of BLFRO film at high fields at various temperatures; inset shows trap ionization energy as a function of E1/2, (c) Log(J) vs Log(E) plots at positive bias and temperature of 100, 140 and 170K; and (d) ln(J/E2) vs (1/E) at positive bias and temperature of 200~350K.

The other interface limited mechanism is Schottky effect due to ionized oxygen vacancies accumulated near the interface. This mechanism arises from the Schottky barrier at the interface of the electrode and ferroelectric film, which could be described

by factorizing the polarization term 2qNeff V / o  st P /  o  st as [220]

q qP q2 N V 0 eff (6.9) J ~ exp{ [ B 2 )  ]} kT4800 op  st   op P

128

Chapter 6 Leakage mechanism of BFO where 0 is the apparent potential barrier,  is the low-frequency dielectric constant, B st

P is the ferroelectric polarization, and Neff is the effective density of fixed charge in the depletion region. The apparent potential barrier, could be obtained from the

2 0 slope of ln(J/T ) as a function of 1000/T, and the zero field potential barrier ( B, app ) could be given by the intercept of vs V1/2. As shown in Figure 6.14, is about 0.480 eV. The true potential barrier can be obtained from

00 qP B   B, app  2 . The low frequency (100 Hz) dielectric constant (  st ) is 40 op  st

176 as shown in Fig 6.12(c), while the polarization value is 70.4 C/cm2. Considering

0 the optical dielectric constant being 6.25 [216], the true potential barrier ( B ) can be determined as 0.803 eV, which is higher than the value of 0.62 eV for (100) epitaxial

BFO film deposited on SRO buffered STO (001) substrate [119]. The enhanced potential barrier of BLFRO film could indicate that the leakage current of BLFRO should be reduced by nearly one order of magnitude compared with that of pure BFO film due to a difference of about 0.2 eV [119]. Meanwhile, the band gap of the

BLFRO film is assumed to be reduce because the La dopants could reduced the band gap of BFO by first-principle calculation [221]. The enhanced potential barrier and reduced band gap could be the reason for the reduced leakage current density of

BLFRO in comparison with BFO film.

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Chapter 6 Leakage mechanism of BFO

Figure 6.14. V1/2 dependence of apparent potential barrier for BFLRO film deposited on Nb-STO substrate.

6.5 Conclusions

1. Deposition temperature is an important parameter for the leakage mechanisms

in Pt/BFO/LNO thin film capacitors. At the deposition temperature from 450 to

550 oC, leakage current shows bulk limited SCLC behavior. For 550 to 650 oC,

the leakage current mechanism show bulk limited PF emission and FN

tunneling emission together.

2. Partial oxygen pressure has been proven as an important parameter affecting

the leakage mechanisms in Pt/BFO/LNO thin film capacitors. At the deposition

partial oxygen pressure from 25 to 50 mTorr, the leakage current shows bulk

limited PF behavior. For 75 mTorr, the leakage current mechanism shows

interface FN tunneling emission together. For 100 and 150 mTorr, the leakage

current shows Schottky behavior.

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Chapter 6 Leakage mechanism of BFO

3. (La, Ru) codoped BFO films have been deposited on (001) Nb-STO substrate

using the pulsed laser deposition system. The electrical and magnetic properties

are affected remarkably by the different nanoscale domain structure for the

films and the different defect chemistry atmospheres in BLFRO film. As a

result, a better coupling between ferroelectric and ferromagnetic ordering has

been obtained via codopping La and Ru ions into BFO film. Additionally,

leakage current mechanisms of BLFRO film have been investigated at various

temperatures. PF and SCLC mechanisms are found to be dominant in high and

low electric field, while the SCLC behavior might be dominant at low

temperature and PF emission and FN tunneling are preferred at high

temperature.

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Chapter 7 EPR study of BFO

7.1 Introduction

The magnetic order of BFO exhibits a long-range modulation associated with a cycloidal spiral with a period length of approximately 62 nm [222]. The spiral

  propagation is along the [101] direction with a spin rotation within (121) plane[223].

The study of magnetic and dynamic properties of BFO using Raman scattering, infrared, or ellipsometric spectroscopy is still in its infancy [223-225]. Scott et al. [224,

226, 227] reported that new magnetic phase transition in bismuth ferrite single crystal is unexpectedly found at 140 and 200 K due to spin-reorientation via Raman scattering study. Furthermore, there have been some reports about the spin glass behavior in BFO under 120 K [228]. There are also reports on the existence of spin-modulated magnetic structures at Fe-sites in BFO via nuclear magnetic resonance (NMR) [229].

Electron paramagnetic resonance (EPR) is a useful tool for exploring spin dynamics in ferromagnetism and antiferromagnetism [230]. Measurements of the magnetic moment provide information about the local magnetic properties, the nature of spin-spin interactions, the distribution of internal field and spin-lattice correlations. In this chapter, the magnetic phase transition of polycrystalline pure and Mn doped BFO is observed via EPR, to identify the magnetic phase changes at low temperature and the effect of chemical modification on magnetic phase transition.

The free trivalent iron ion possesses five unpaired electrons in the 3d shell 3d 5 , hence

5 the discussion on the corresponding high spin 3d53 Fe  EPR spectra ( S  ), in 2 principle, should be based on the most general form of spin Hamiltonian, including

132

Chapter 7 EPR study of BFO fourth-rank tensor component [231, 232] An approximate spin Hamiltonian can be written as [233]

k  q  k qq H ge eB0 S+ B k O k ( S x , S y , S z ) (7.1) k2,4,6

Here, the g-matrix is taken as isotropic with ge  2.0023, the free electron g-value, e

q denotes the Bohr magneton, B0 is the external magnetic field, Bk is the fine-structure

q (FS) spin-Hamiltonian coefficient, and Ok are the extended Stevens spin operators

[234]. The first term represents the electronic Zeeman interaction while the second term is the effective FS Hamiltonian describing the interaction of the crystal field with the paramagnetic ion [234, 235]. The electronic g-matrix can be treated as isotropic and is expected to be close to that of the free electron For Fe3+, this g value is typically

giso  2.002 [236, 237]. Meanwhile, the g factor could be obtained via g hv/  B A g factor gives an important clue as to the ion responsible for the intense line.

The temperature dependence of the linewidth has been found to be described by the simple formula

HTTTH()[ 0 ()/()]  ()  (7.2)

where HT()denotes the linewidth at temperature T, 0 ()T is the free spin susceptibility, ()T is the measured susceptibility, and the symbol H() is a temperature-independent constant.

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Chapter 7 EPR study of BFO

7.2 EPR study of pure BFO

The X-band (9.5 GHz) continuous wave EPR measurements were carried out on a

Bruker EMX EPR spectrometer with a flow nitrogen cryostat, the accuracy is about 0.1

K. Additionally, a standard field marker (polycrystalline 2, 2-diphenyl-1- picrylhydrazyl, DPPH, with g = 2.0036) was used in the calibration of the resonance magnetic field values and also in the determination of the exact g-factor of the samples.

Figure 7.1(a) shows the powder X-ray diffraction spectra of the calcined sample of

BFO, confirming the single phase BFO. The powder spectra are phase perovskite structure with a rhombohedral space group of R3c. The normalized X-band EPR spectra recorded at room temperature for BFO powders are shown in Figure 7.1(b).

The spectra are dominated by two resonance parts: one is at low field resonance shoulder (LF) near g = 4.11, characteristic of magnetically isolated high spin S = 5/2

Fe3+ ions in a low symmetry environment (tetragonal coordinated Fe3+ ions),

3  corresponding to ()Fe Vo defect dipoles. Meanwhile, the other resonance part is a much broader resonance at high field (HF) near 340 mT as well as g = 2.05, which occur due to Fe3 ions in octahedral. From the normalized EPR spectra measured at different microwave energy, it can be concluded that only one phase rather than two phases exists in the sample.

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Chapter 7 EPR study of BFO

Figure 7.1 (a) X-ray diffraction spectra of BiFeO3 polycrystalline powder. (b) normalized X-band EPR spectra of BiFeO3 polycrystalline powders measured under different microwave energy at room temperature.

To conclude the interpretation for the EPR signal at LF and HF, the samples were annealed in an oxygen and argon atmospheres respectively. BFO is a mixed-valent

2+  system with fraction of Fe due to oxygen vacancies (VO ), which can be described as

32X   2Fe OOO  2 Fe  V  1 2 O2 . The X-band EPR spectra recorded at room temperature for the annealed samples and the original samples are shown in Figure 7.2.

The LF resonance shoulder disappears after annealing in an oxygen atmosphere, which means that the LF line originates from the oxygen vacancies ( ) or defect dipole

3  3+ ()Fe Vo in the sample, while the HF line is enhanced because the Fe concentration is increased via the oxidization. The spectrum annealed in an argon atmosphere, the LF resonance shoulder at around 150 mT is enhanced due to the signal come from or . The spin counting results for these three samples annealed under different atmospheres are listed in Table 7.1. From the spin counting results, it was found that the electron number of powder annealed in the oxygen atmosphere increased slightly, while that of powder annealed in argon atmosphere has

135

Chapter 7 EPR study of BFO increased sharply, indicating that reducing atmosphere can produce a lot of electrons, reducing the Fe3+ to Fe2+ and forming many oxygen vacancies and defect dipoles.

Figure 7.2 X-band EPR spectra for BFO powders treated under different annealing atmospheres recorded at room temperature.

Table 7.1. Spin counting for samples annealed in different atmospheres

Annealing Air O2 Ar atmosphere

Electron Number/g 6.06E+18 7.92E+18 2.10E+19

Figure 7.3 shows the X-band EPR spectra recorded in the 120 ~ 300 K temperature range. It can be observed that the low field shoulder of spectrum (LF) and the high- field main resonance (HF) are both temperature dependent. The observed broaded

Lorentzian line at high-field along the entire temperature range is due to the spin of the

Fe3+ ions. Scott et al. [80, 224, 226, 227] reported the presence of new magnetic phase transitions in BFO at 140 and 200 K due to a spin-reorientation transition as in orthoferrites. To further analyze the observed EPR spectra as a function of temperature,

136

Chapter 7 EPR study of BFO the present research work explored the EPR parameters related to the magnetic structure dependence of temperature: the resonance shift (Bc), the g-factor, the peak-to-

peak width ( BPP ), the integrated intensity ( IEPR ), and the asymmetry parameter.

Figure 7.3. Temperature dependence of X-band EPR spectra of polycrystalline BFO powder.

The temperature dependence of the shifts in position in the EPR spectra is shown in

Figure 7.4(a). The positions of the LF and HF resonances fall off with decreasing temperature range, and the resonances show a linear behavior running almost parallel to each other [238]. However, there are breaks in the shift versus T curve at around

140, 200, and 260 K for the HF resonance and 200 K for the LF resonance, which might be due to spin-reorientation, arise to the charge localization in the BFO occurs at low temperatures, Fe2+-O-Fe3+ clusters and thermally activated hopping-inducing ferromagnetic interactions [239].

Temperature dependence of the g-factor of the EPR spectra is plotted in Figure 7.4(b).

The g-factor shows a value larger than free electron of 2.0023 and DPPH of 2.0036, and it increases with temperature fall-off. This may be caused by a gradual build-up of

137

Chapter 7 EPR study of BFO the orbital ordering at low temperature [240], which can change the spin–orbit coupling as well as the crystal-field parameters and hence can lead to an increase in the value of g-factor as the temperature is lowered. The g-factor of HF and LF resonance shows a sharp decrease at 140, 200 and 260 K. It is crucial to consider the displacement of the ions from their centrosymmetric positions in multiferroic materials

[241]. The shift of Bi3+ and Fe3+ ions from their ideal perovskite structure contributes to their multiferroic properties. The shift of Bi3+ ions gradually decreases with temperature [242, 243] while the shift of Fe3+ ions remains constant within statistical error of 0.252(6) Å. Meanwhile, the short and long Bi-Fe distances, Fe-O distance, and the Fe-O-Fe angler inside the octahedra decrease with decreasing temperature below

550 K.The oxygen octahedra rotation angle however increases with temperature fall- off [242]. The ion displacements will break the centrosymmetry of BFO and induce a

 spin-reorientation in the 140, 200 K and 260 K. The defect chain Fe Vo Fe angle in the oxygen octahedra may vary with temperature, which is responsible for the canting of the magnetic moments.

The EPR intensity IEPR is expected to be proportional to the dc susceptibility  of the spins, as can be seen in Figure 7.4(c). The EPR intensity decreases with increasing temperature having a minimum at 140 K and then increases steadily with the increase in temperature. Perhaps being induced by magnetic fluctuations at 140 K as slows down the spin reorientation transition [226]. However, similar behavior cannot be seen at 200 and 260 K, which may be attributed to the transitions near 200 and 260 K having much greater strain coupling than that at 140 K [226] or/and the existence of competition between spin-glass and ferromagnetic ordering below 150 K [244].

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Chapter 7 EPR study of BFO

Figure 7.4. Temperature dependence of the resonance field, (a) BC , (b) g-factor, and (c) the intensity of EPR, IEPR indicate that magnetic phase transitions near 140, 200, and 260 K for BFO powder. The solid lines are guides for eyes.

The temperature dependence of the peak-to-peak linewidth BPP is shown in Figure

7.5(a). It can be seen that the linewidth decreases with increasing temperature, which is found to be common in antiferromagnetic materials[245, 246]. The solid line in Figure

7.5(a), can be fitted via the equation [247]

A B( T )   B  exp(  E / k T ) (7.3) PP0 T a B

139

Chapter 7 EPR study of BFO

where Ea , is the thermal activation energy for the antiferromagnetic spin cluster.

The linewidth curve may correspond to the conductivity curve, and is thus able to explain the mechanism for the magnetic phase transition. The dominant process of relaxation of the spin is the spin-lattice relaxation rather than spin-spin relaxation. The observed narrowing of the EPR peaks with the increase in temperature in the present case can be attributed to an exchange interaction between unlike spins.

As mentioned above, some Fe2+ ions and oxygen vacancies coexist in the BFO

2+ 3+ polycrystalline samples. Therefore, BiFeO3 can be written as Bi(Fe Fe )O3-x with the bracketed Fe2+ and Fe3+ ions occupying the octahedrally coordinated sites with d5 and d6 electron configurations respectively, and the formation of Fe2+ on the sites adjacent to the vacancy coordinated sites [187]. According to the charge ordered module, it can be predicated that Fe2+/Fe3+ charge order occurs at low temperature and the

Fe23 Fe charge fluctuation persists even in the charge ordered states [248].

Figure 7.5(b) shows the linewidth data plotted as ln(BTPP ) verse 1000/T. The activation energy of 1.288 ± 0.02 eV deduced from the present EPR linewidth for BFO, corresponding to the high-temperature conductivity data showed an activation energy of 1.03 ±0.05 eV for BFO [249] .

140

Chapter 7 EPR study of BFO

Figure 7.5 (a) Temperature dependence of the peak-to-peak EPR linewidth BPP and

(b) ln(BTPP ) as a function of 1000/T. The solid lines are guides for the eyes only.

The temperature dependence of the asymmetric parameter A/B is shown in Figure 7.6.

This asymmetric parameter is an indication of the extent of anisotropy in the magnetic structure of the sample [250]. As can be seen in Figure 7.6, the spin-reorientation behavior occurs at 140, 200 and 260 K. The large value of A/B for BFO indicates a large magneto crystalline anisotropy in the magnetic interactions.

141

Chapter 7 EPR study of BFO

Figure 7.6. Variation of A/B ratio with temperature in BFO. Inset illustrates the method adopted to calculate the A/B ratio.

Recent EPR experiment investigated the magnetic properties of multiferroic BiFeO3.

Broad resonance at low field shoulder and high field are observed, indicating oxygen vacancy and Fe3+ defect respectively. The work demonstrates that the magnetic phase transition occurs at 140 and 200 K due to spin-reorientation via temperature dependent

EPR signals like resonance field, g-factor, linewidth, intensity, and asymmetric parameter. The magnetic phase transition may be related to the displacement of ions with variation in temperature. Furthermore, the temperature-dependent linewidth shows an activation energy close to that of electrical conductivity.

142

Chapter 8 Conclusions and Future Work

Chapter 8 Conclusions and Future work

8.1 Conclusions

In this work, pure and chemically modified BFO thin films have been systematically grown via a pulsed laser deposition system. The conclusions of these studies could be summarized as follows:

1. Epitaxial BFO thin films as a function of substrate orientation have been fabricated

on the SrTiO3 substrates. The surface morphology and local piezoresponse

properties and magnetic domain structure have been investigated via the AFM,

PFM and MFM, indicating that the substrate orientation could remarkably affect

the grain size and domain structure. The average piezoresponse constant of BFO at

different orientated films have been calculated as a function of ac voltage. The

local swishing behavior has been determined via the SS-PFM, indicating the

nanoscale domain switching dynamics. The effect of bottom electrode on surface

potential of polycrystalline BFO grown on the Pt/TiO2/SiO2/Si and LNO/Si

substrate were compared via the KPFM.

2. La-doped BFO (BLFO), Ru doped BFO (BFRO), and (La, Ru) codoped BFO

(BLFRO) thin films on Pt/TiO2/SiO2/Si substrate by pulsed laser deposition were

investigated. The films show a pure and polycrystalline phase, and possess uniform

surface morphologies. The domain growth habits are investigated via PFM,

indicating that the domain structure is changed and domain wall density is

decreased by La, Ru doping and (La, Ru) codoping.

3. The structural characterization by means of X-ray diffraction and Raman spectra

analysis revealed the existence of a MPB region between tetragonal and

rhombohedral symmetries in the 2mol% and 5mol% PZT modified BFO films on

143

Chapter 8 Conclusions and Future Work

the (111) Pt/TiO2/SiO2/Si substrate grown by pulsed laser deposition system. The

ferromagnetic properties were enhanced by effects such as Pb ions suppressing the

spiral spin structure, Zr, Ti ions introducing chemical pressure and change the Fe-

O-Fe bond angle, or even substituents causing destruction of a cycloidal spin

structure. The leakage current of the PZT2%BFO and PZT5%BFO films decreases

by about one and two orders of magnitude at high electric field comparing to pure

BFO film. As a result, a better coupling between ferroelectric and ferromagnetic

ordering has been obtained via introducing PZT into the BFO system.

4. The deposition temperature and partial oxygen pressure are two important

parameters for the leakage mechanisms in Pt/BFO/LNO thin film capacitors. At a

fixed oxygen partial pressure (50 mTorr) and deposition temperature from 450 to

500 oC, the leakage current shows bulk limited SCLC behavior. For 550 to 650 oC,

the leakage current mechanism shows bulk limited PF emission and FN tunneling

emission together. At a fixed deposition temperature (500 oC) and partial oxygen

pressure from 25 to 50 mTorr, the leakage current shows a bulk limited PF

behavior. For 75 mTorr, the leakage current mechanism shows bulk limited PF and

interface limited FN tunneling emission together.

5. The magnetic phase transition temperatures in pure and Mn doped BFO were

observed via low temperature EPR measurement. Broad resonances at low field

shoulder and high field are observed, indicating presence of oxygen vacancy and

Fe3+ defect respectively. It is demonstrated that the magnetic phase transition

occurs at 140 and 200 K for pure BFO due to spin-reorientation via temperature

dependence EPR signals like resonance field, g-factor, linewidth, intensity, and

asymmetric parameter.

144

Chapter 8 Conclusions and Future Work

8.2 Future Work

Although this work has succeeded in presenting a clearer picture on the enhanced magnetic and electric properties of the chemically modified BFO films fabricated via

PLD, more works should concentrate on the clarification on some existing issues mentioned below:

1. In-plane domain structure in the epitaxial pure and chemically modified BFO

thin films have not been observed. This is important for the electric properties

and should be investigated in detail.

2. The doping concentration affects the structure of the BFO films, and the La, Ru

and codoped La and Ru doping concentration should be paid more energy to

image the phase image of the modified BFO films.

3. Further research should be carried out on the defect chemistry of the

chemically modified BFO films on different substrates without Fe

contamination since the work on BFO film on SrTiO3 cannot produce a clear

image of the defect chemistry of BFO films due to the difficulty in dividing the

Fe origin in BFO from the STO substrate. Other substrate could be helpful to

prevent the Fe signal from STO substrate, such as sapphire and glass substrate.

145

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