Structure and Properties of Lead Zirconate Titanate Thin Films by Pulsed Laser Deposition
GOH WEI CHUAN (B.Sc (Hons), NUS)
A THESIS SUBMITTED FOR THE DEGREE OF Doctoral of Philosophy Department of Physics National University of Singapore 2005
Acknowledgements
I would like to express my deepest gratitude to my supervisors, Prof. Ong Chong Kim
and Dr. Yao Kui. I would like to thank Prof. Ong for giving me the opportunity to
study and perform research work in the Center of Superconducting and Magnetic
Materials (CSMM). His passion and enthusiasm in the search for understanding the
underlying physics of the experiments have deeply influenced my mindset in
conducting experiments and will continue to be my source of inspiration and guidance.
Without Prof. Ong’s constant guidance and criticism, I would have lost my bearing in
the vast sea of knowledge, and would not have reached this far.
I would also like to express my greatest appreciation to Dr. Yao Kui in Institute of
Material Research and Engineering (IMRE). His constant advice and meticulous
attention to the theoretical and experimental details had deeply influenced my way of
research both in designing experiments and interpreting the results. Without his
supervision and encouragements in countless hours of his time, it would not be
possible for me to complete my publications and thesis. For that I am in debt to him
and will forever remember his advice when pursuing my future endeavors.
I am indebted to my fellow colleagues in CSMM, IMRE and Department of Physics,
NUS, including A/P Sow Chorng Haur, Xu Sheng Yong, Wang Shi Jie, Li Jie, Yang
Tao, Tan Chin Yaw, Rao Xue Song, Chen Lin Feng, Yan Lei, Kong Lin Bing, Liu
Hua Jun, Lim Poh Chong, Yu Shu Hui, Gan Bee Keen and all those have shared their time helping me and discussing with me in this project. Their help are greatly appreciated.
i
I would also like to acknowledge the financial support from the National University of Singapore for providing scholarship during this course of study.
Last but not least, I would like to thank my family, especially Jin Yu, for supporting
me and helping me both spiritually and financially throughout the long years in pursuing my dream in doing research in the scientific field. None of this would be possible without their love and concern.
ii
Table of Contents
Page Acknowledgements i Table of Contents iii Summary vi List of Publications ix List of Figures x List of Tables xiii List of Symbols xiv
1. Introduction 1 1.1 History and applications of ferroelectric materials 1 1.2 Motivation, aim and objective of the thesis 2 1.3 Ferroelectricity 4 1.4 Piezoelectricity 8 1.5 Lead Zirconate Titanate Oxide (PZT) 15 1.6 Current status and problems in PZT research 16 1.7 References 19 2. Apparatus and experimental procedures 23 2.1 Samples preparation 23 2.1.1 Pulsed laser deposition 23 2.1.2 Sol-gel 26 2.2 Crystal and microstructure characterizations 29 2.2.1 X-ray diffractions 29 2.2.2 Scanning electron microscope and atomic force 31 microscope 2.3 Electrical and ferroelectric characterizations 34 2.3.1 Impedance, dielectric and loss tangent measurement 34 2.3.2 Ferroelectric loop measurement 35 2.3.3 Piezoelectric constant measurement 37
iii
2.4 References 39 3. Early growth stage of PLD-derived PZT film 41 3.1 Introduction 41 3.2 Experimental procedures 42 3.3 Results and discussion 44 3.3.1 Crystal structure 44 3.3.2 Microstructure 46 3.4 Conclusions 55 3.5 References 57 4. Effects of microstructure on the properties of 59 polycrystalline PZT thin film 4.1 Introduction 59 4.2 Experimental procedures 59 4.3 Results and discussion 61 4.3.1 Crystal structure 61 4.3.2 Microstructure 61 4.3.3 Electrical properties 65 4.3.4 Piezoelectric properties 68 4.4 Conclusions 71 4.5 References 72
5. Epitaxial La0.7Sr0.3MnO3 conductive film as bottom 74 electrode for PZT film 5.1 Introduction 74 5.2 Experimental procedures 75 5.3 Results and discussion 76 5.3.1 Crystal structure 76 5.3.2 Microstructure 79 5.3.3 Resistivity and magnetoresistance 81 5.3.4 Discussion 82 5.4 Conclusions 86 5.5 References 88
iv
6. Pseudo-epitaxial lead zirconate titanate (PZT) thin 89 film on silicon substrate 6.1 Introduction 89 6.2 Experimental procedures 89 6.3 Results and discussion 90 6.3.1 Crystal structure 90 6.3.2 Microstructure 92 6.3.3 Electrical properties 94 6.3.4 Piezoelectric properties 97 6.4 Conclusions 98 6.5 References 100 7. Overall conclusions and future work 102
7.1 Conclusions 102 7.2 Future work 104
v
Summary
Ferroelectric materials have attracted great research interests due to their
extraordinary electrical and electromechanical properties. As a result of these
properties, we had seen the applications of ferroelectric materials in ferroelectric
random-access memories (FERAMs), dynamic random-access memories (DRAMs),
gate oxides, piezoelectric sensors, actuators and micro-electro-mechanical systems
(MEMS).
In this thesis, we chose Lead Zirconate Titanate (PZT) for investigation, in
consideration of its excellent piezoelectric properties and a variety of potential applications. Pulsed laser deposition (PLD) was used as the main fabrication method for growing the PZT thin films. Sol-gel deposition method was also used in some of the experiments to produce PZT films with different morphology. We focused our research on investigating the microstructure of the PZT thin films and the correlation between the structure and performance properties of the PZT films.
The early growth stage of the PZT film on SrTiO3 (STO) substrate using PLD
was investigated. The PZT film deposited onto STO underwent a three dimensional
island growth mode. A two layer growth structure was observed for the PZT film with
a thickness of about 40 – 50 nm. As the PZT film increased, small grains start to
merge into large grains. Further increase in PZT film thickness finally led to column-
like growth mode. This growth structure was favorable because it would help
maintaining an acceptable surface roughness while the film thickness is further
increased.
The effect of microstructure on the performance of PZT film was investigated
by using two types of PZT films with different morphologies fabricated via PLD and
vi
sol-gel methods respectively. We observed that difference in microstructure would
significantly offset the electrical properties of the films. The PZT film with denser
microstructure would have a significantly higher dielectric constant and remnant
polarization with lower coercive field. However the microstructural difference
resulted only in relatively smaller difference in the loss tangent and piezoelectric
properties. As a result, PZT films with looser microstructure could have a higher
piezoelectric voltage constant due to the lower dielectric constant.
Currently, microelectronic technologies are mainly based on silicon
technology. Thus it is important to fabricate PZT film on silicon substrate so that they
can be integrated into silicon technology. Due to the large difference in thermal
expansion coefficient and lattice constant between silicon and PZT, and the diffusion
of silicon into PZT, we had to use buffer layers to address these problems. In this
thesis, we selected Yttria-Stabilized-Zirconia oxide (YSZ) and Yttrium Barium
Copper oxide (YBCO) as buffer layers and found that they well compensated for the
difference in lattice constant and provided an effective diffusion barrier to prevent
silicon from diffusing into PZT film. Platinum is commonly used as the bottom electrode for PZT on silicon substrate. But PZT film grown on platinum is usually polycrystalline and has poor electrical and piezoelectric properties. In this thesis, we chose La0.7Sr0.3MnO3 (LSMO) as bottom electrode because it has good lattice
matching with PZT and can be used as a buffer layers at the same time.
We had successfully fabricated a pseudo epitaxial PZT film on silicon substrate using LSMO/YBCO/YSZ heterostructure. The pseudo epitaxial PZT film
had a good crystallographic orientation but with granular microstructure and with
nano-sized pores distributed all over the film. Although the epitaxial quality of the
film was imperfect, we found that the remnant polarization of the film was
vii
substantially larger than that of the high quality epitaxial PZT film directly deposited on silicon substrate. We attributed the enhanced ferroelectric property of our PZT film to the partial relief of tensile stress by virtue of the granular pseudo epitaxial feature with nano-sized pores. It is therefore concluded that only improving epitaxial quality without considering the tensile stress effects may not be sufficient in achieving optimal ferroelectric properties for a ferroelectric film on silicon substrate.
viii
List of Publications
1. W. C. Goh, K. Yao and C. K. Ong, Pseudo-epitaxial lead zirconate titanate thin film on silicon substrate with enhanced ferroelectric polarization, Applied Physics Letters, 87, 072906 (2005).
2. W. C. Goh, K. Yao and C. K. Ong, Epitaxial La0.7Sr0.3MnO3 thin films with two in-plane orientations on silicon substrates with yttria-stabilized zirconia
and YBa2Cu3O7-δ as buffer layers, Journal of Applied Physics, 97, 073905 (2005).
3. W. C. Goh, K. Yao and C. K. Ong, Effects of microstructure on the properties of ferroelectric lead zirconate titanate (PZT) thin films, Applied Physics A: Material Science & Processing, 81, 1089 (2005).
4. W. C. Goh, S. Y. Xu, S. J. Wang, and C. K. Ong, Microstructure and growth
mode at early growth stage of laser-ablated epitaxial Pb(Zr0.52Ti0.48)O3 films
on a SrTiO3 substrate, Journal of Applied Physics, 89, 4497 (2001).
5. L. Yan, W. C. Goh, and C. K. Ong, Magnetic and electrical properties of
La0.7Sr0.3MnO3–Zn0.8Co0.2Al0.01O junctions on silicon substrates, Journal of Applied Physics, 97, 103903 (2005).
6. L. Yan, L. B. Kong, T. Yang, W. C. Goh, C. Y. Tan, C. K. Ong, Md. Anisur Rahman, T. Osipowicz, and M. Q. Ren, Enhanced low field magnetoresistance
of Al2O3-La0.7Sr0.3MnO3 composite thin films via a pulsed laser deposition, Journal of Applied Physics, 96, 1528 (2004).
7. S. T. Tay, C. H. A. Huan, A. T. S. Wee, R. Liu, W. C. Goh, C. K. Ong, and G.
S. Chen, Substrate temperature studies of SrBi2(Ta1 - xNbx)2O9 grown by pulsed laser ablation deposition, Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films, 20, 125 (2002).
ix
List of Figures
Pg Figure 1.1 Polarization versus electric field loop. The solid line indicates a perfect ferroelectric crystal; the dashed line shows a typical 5 ferroelectric material loop. (Adapted from reference 59)
Figure 1.2 Crystal structure of BaTiO3 (Adapted from reference 59) 6
Figure 1.3 Various phase transitions in barium titanate. (Adapted from 7 reference 59)
Figure 1.4 The piezoelectric effect: (a) longitudinal effect; (b) transverse 10 effect. (Adapted from reference 60)
Figure 1.5 Strain characteristic for a typical piezoelectric ceramic. 11 (Adapted from reference 60)
Figure 1.6 Schematic illustration of strains induced by pole reversals in 13 ferroelectric ceramic materials. (Adapted from reference 60)
Figure 1.7 Phase diagram of lead zirconate titanate (PZT). (Adapted from 16 reference 57)
Figure 2.1 A schematic drawing of the pulsed laser deposition (PLD) 25 system.
Figure 2.2 X-ray diffraction (XRD) θ-2θ scan 30
Figure 2.3 Schematic drawing of a scanning electron microscope (SEM). 32
Figure 2.4 Schematic drawing of an atomic force microscope (AFM) 34 operating in tapping mode. (Adapted from reference 17)
Figure 2.5 Equivalent Circuit of Sawyer Tower Mode 36
Figure 2.6 Equivalent Circuit of Virtual Ground Mode 37
Figure 2.7 Schematic drawing of the piezoelectric constant measurement 38 experimental setup.
Figure 3.1 XRD θ−2θ scan of the PZT film deposited on SrTiO (100) 3 45 substrate for 20 minutes.
Figure 3.2 C-axis lattice constant (nm) and full width at half maximum (FWHM) of PZT (002) rocking curve against different 46 thickness.
x
Figure 3.3 AFM surface images (1μm x 1μm) of SrTiO3 substrate and the PZT films deposited on SrTiO3 substrate with different thickness. (a) SrTiO3 substrate surface, (b)-(h) PZT films on 48 SrTiO3 substrate with different deposition times, from 1 to 3, 6, 10, 20, 30 and 40 min respectively.
Figure 3.4 AFM surface profile of the PZT film with a thickness of about 49 50 nm, the size of the bottom image is 3μm x 3μm.
Figure 3.5 Root mean square roughness and average in-plane grain size 51 versus film thickness.
Figure 3.6 A schematic diagram of the PZT film growth structure. 52
Figure 3.7 HRTEM cross-section view of the PZT film deposited on 54 SrTiO3 substrate for 3 min.
Figure 3.8 High magnification image of the interface structure between 54 PZT film and the SrTiO3 substrate.
Figure 3.9 HRTEM cross-section view of the PZT film deposited on SrTiO3 substrate showing a structural defect on the SrTiO3 55 substrate.
Figure 4.1 XRD patterns of the PZT thin films deposited by (a) PLD 61 method and (b) sol-gel method.
Figure 4.2 Surface SEM images of the PZT thin films deposited by (a) PLD method and (b) sol-gel method; Cross-sectional SEM images of 63 the PZT thin films grown by (c) PLD and (d) sol-gel.
Figure 4.3 P-E hysteresis loops of the PZT thin films grown by PLD and 65 sol-gel.
Figure 4.4 Dielectric constant and loss tangent of the PZT films grown by 67 (a) PLD and (b) sol-gel.
Figure 4.5 Displacement measurement results in the frequency domain for 70 the PZT films grown by (a) PLD and (b) sol-gel.
Figure 5.1 XRD patterns (θ-2θ) for the LSMO thin films deposited on 77 YBCO/YSZ/Si and YSZ/Si substrates.
Figure 5.2 (a) XRD patterns (φ scan) for (I) Si {113}, (II) YSZ {113} and (III) LSMO {113} peaks for the LSMO/YBCO/YSZ/Si 78 multilayer. The inset figure shows the relative φ angle relationship between YBCO {103} and LSMO {103}.
xi
Figure 5.3 Cross-sectional SEM images of (a) LSMO/YBCO/YSZ/Si and 80 (b) LSMO/YSZ/Si.
Figure 5.4 Temperature dependence of resistivity for the LSMO thin films deposited on (a) LAO, (b) YBCO/YSZ/Si and (c) YSZ/Si 81 substrates.
Figure 5.5 Magnetoresistance (MR) vs. magnetic field for the LSMO thin 82 films deposited on (a) YBCO/YSZ/Si and (b) YSZ/Si substrates.
Figure 5.6 A schematic diagram of the proposed epitaxial growth mechanisms of the YBCO thin film for the two observed in- plane orientations on the YSZ buffer layer: (a) YBCO [100](001)//YSZ [100](001), in which the Cu-O plane of the 83 YBCO is grown on the O-O plane of the YSZ (3/4 height of the unit cell of YSZ); (b) YBCO [110](001)//YSZ [100](001), the Cu-O plane of the YBCO is turned by 45º in-plane on the Zr-Zr plane of YSZ.
Figure 5.7 Structure of (a) orthorhombic YBCO crystal and (b) perovskite 85 LSMO crystal.
Figure 6.1 XRD patterns (θ-2θ) for the PZT thin films deposited on 91 LSMO/YBCO/YSZ/Si.
Figure 6.2 XRD patterns (φ scan) for (I) Si {113}, (II) LSMO {113} and (III) PZT {113} peaks for the PZT/LSMO/YBCO/YSZ/Si 91 multilayer.
Figure 6.3 SEM images of (a) surface and (b) cross-section of 93 PZT/LSMO/YBCO/YSZ/Si heteromultilayer structure.
Figure 6.4 P-E hysteresis loop of a PZT thin film sample, deposited on 95 YBCO/YSZ/Si heteromultilayer structure.
Figure 6.5 Dielectric and loss tangent of PZT film, deposited on 95 YBCO/YSZ/Si heteromultilayer structure.
Figure 6.6 XRD pattern (θ-2θ) of PZT (001) peaks for the target and film. 97
Figure 6.7 Displacement measurement results in the frequency domain for the PZT films on LSMO/YBCO/YSZ/Si at AC peak-to-peak 98 voltage of 9 V at 1500 Hz.
xii
List of Tables
pg Table 3.1 Measurement results for the PZT films grown on STO 45 substrates.
Table 4.1 The properties of the PZT films grown by PLD and sol-gel 66 deposition.
Table 5.1 Summary of PLD processing parameters for the YSZ, YBCO, 76 and LSMO thin films.
xiii
List of Symbols
Ps Spontaneous polarization
Pr Remnant polarization
Ec Coercive field
εο Permittivity of free space
Τ εr Relative dielectric constant
σΕ Poisson’s ration d33 Piezoelectric strain constant in the z-direction (for longitudinal effect) g33 Voltage output constant
Å 10-10 m
xiv Chapter 1 Introduction
Chapter 1 Introduction
1.1 History and applications of ferroelectric materials
Ferroelectric materials are widely used as capacitors, transducers, sensors, and
actuators. The history of the discovery of ferroelectricity can be traced back to the
famous work of physicists Weiss, Pasteur, Pockels, Hooke Grroth, Voigt and brothers
Vurie. The early work on ferroelectrics was mainly focused on Rochelle salt1 from
1890 to 1935 and later on potassium dihydrogen phosphate2-3 after 1935. However the
research on ferroelectricity at that time was confined to a temperature of lower than
−150 °C. Only after the World War II, the study on ferroelectricity was accelerated
rapidly with the discovery of barium titanate.4-6
From 1940 to 1950, it was considered as the early barium titanate era. During this
period major experimental features of this fascinating crystal were first studied.
7 Capacitor and transducer applications of BaTiO3 were also developed during this
period of time. From 1950 to 1960, huge influxes of different ferroelectric materials
were discovered. The number of known ferroelectric materials grew from Rochelle
Salt, potassium dihydrogen phosphate (KDP) and several ferroelectric materials in
1950 to a vast twenty-five firmly established families of ferroelectrics with more than
twenty perovskite compound and innumerable solid solutions in the early 1960. Some
of the early work on other materials such as Pb(Zr,Ti)O3 were also discovered during
this period of time.8-12
The most significant theoretical development in ferroelectrics occurred in 1960
with the introduction of the soft-mode description of ferroelectric transition made by
Cochran.13-14 The theoretical model was further extended after it was further verified
by various experiments such as Raman, Brillion and Rayleigh scattering, as well as nuclear magnetic resonance (NMR), nuclear quadruple resonance (NQR) and electron
1 Chapter 1 Introduction
paramagnetic resonance (EPR). During the years from 1970 to 1980, the research was
mainly focused on doped BaTiO3 which led to the discovery of remarkable change in
resistance at the Curie point. From 1980 to 1990, the research direction turned to the
integration and applications of ferroelectric materials.
Current research has led to the integration of ferroelectric films for various industrial applications. In the semiconductor industry, we have seen the integration of ferroelectric films for ferroelectric random-access memories (FERAMs)15-29 and
dynamic random-access memories (DRAMs).30-32 Basic field effect transistor (FET)33-
38 that utilizes the remnant polarization of ferroelectric field has been developed and
fabricated. Since ferroelectric films such as PZT have extremely high dielectric
constant, there are researches that focus on integrating PZT onto silicon substrate as
39-41 the gate oxide to replace the SiO2.
Besides the applications in electronics, piezoelectric films, especially PZT,
have been widely used in electro-mechanical devices. PZT is one of the most promising materials used in the development Micro-Electro-Mechanical Systems
(MEMS).42-48 PZT materials have been used as actuators in various devices, such as micro-pumps, micro-valve, micromotors49 and precision positioning etc.50-55
Ferroelectric materials have also found many other applications for instances, optical
waveguides, pyroelectric infrared sensors and surface acoustic filter, most of which
utilized PZT.56-58
1.2 Motivation, aim and objective of the thesis
The main motivation for me in this thesis was to develop ferroelectric films
using appropriate fabrication method such that the ferroelectric films exhibit the
properties required for device applications. PZT was selected in this study due to its
2 Chapter 1 Introduction various advantages over other ferroelectric materials, such as stable crystal structure, good ferroelectric and piezoelectric properties. Pulsed Laser Deposition (PLD) was used to deposit our PZT film and buffer layers, mainly due to its high deposition efficiency as well as excellent control over the stiochiometry of the deposited films.
The aim of this thesis was to fabricate epitaxial PZT films with superior electrical and piezoelectric properties on silicon substrates using PLD and then to understand the relationship between structure and performance properties. However, due to the large thermal expansion mismatching, large lattice mismatching and inter- diffusion between PZT and silicon substrate, appropriate buffer layers and bottom electrode have to be carefully selected and developed so that PZT films with excellent properties can be obtained on silicon substrate. In this thesis, Yttria-Stabilized-
Zirconia (YSZ) and Yttrium Barium Copper Oxide (YBCO) were chosen as buffer layers, while La0.7Sr0.3MnO3 was selected as bottom electrode. It has been shown that
PZT thin films grown on conductive perovskite oxide have better electrical properties and improved fatigue characteristics, compared to those grown on metallic platinum bottom electrode.
By understanding the crystallization process and physical properties, such as the film growth mechanism, crystallographical orientation, microstructure, for our
PZT fhin films and their effects on the properties of the films, such as the dielectric, ferroelectric and piezoelectric characteristics, we will be able to establish the optimized fabrication process to obtain PZT films with promising properties required for different applications.
3 Chapter 1 Introduction
1.3 Ferroelectricity
In dielectric materials, the constituent atoms are considered to be ionized to a
certain degree and are either positively or negatively charged. In such ionic crystals,
when an electric field is applied, cations are attracted to the cathode and anions to the
anode due to electrostatic interaction. The electron clouds also deform, causing
electric dipoles. This phenomenon is commonly known as electric polarization of the
dielectrics. The polarization is expressed quantitatively as the sum of the electric
dipoles per unit volume.
Depending on the crystal structure, the centres of the positive and negative charges may not coincide even without the application of an external electric field.
Such crystals are said to posses a spontaneous polarization. When the spontaneous polarization of a dielectric material can be reversed by an electric field, it is called ferroelectrics.
In ferroelectric materials, the domain states differ in orientation of
spontaneous electric polarization, and the ferroelectric character is established when it
is evident that the states can be transformed from one to another by suitable
application of electric field. The ability to re-orientate the domain state polarizations
separates these materials from the larger class of pyroelectric crystals in the 10 polar-
point symmetries. Saturated polarization (Ps), remnant polarization (Pr) and coercive
field (Ec) are defined by analogy with corresponding magnetic quantities. A
ferroelectric crystal would have a polarization loop as shown in Figure 1.1.
4 Chapter 1 Introduction
+Ps
+Pr
-Ec +Ec
-Pr
-Ps
Figure 1.1 Polarization versus electric field loop. The solid line indicates a perfect ferroelectric crystal; the dashed line shows a typical ferroelectric material loop. (Adapted from reference 59)
The structure of a typical ceramic ferroelectrics, barium titanate (BaTiO3), is shown in Figure 1.2. BaTiO3 has a perovskite crystal structure. In the high temperature paraelectric phase (non-polar phase), there is no spontaneous polarization.
Below the transition temperature Tc (Curie temperature), spontaneous polarization occurs, and the crystal structure becomes slightly elongated, that is, tetragonal. Figure
1.3 schematically shows the temperature dependence of the spontaneous polarization
Ps and permittivity ε. Ps decreases with the increase of temperature and vanishes at
Curie temperature, while ε tends to diverge near Tc. Also, the reciprocal permittivity
1/ε is known to be linear with respect to the temperature over a wide range in the paraelectric phase (Curie-Weiss law),
5 Chapter 1 Introduction
C ε = (1.1) ()−TT 0
where C is the Curie-Weiss constant and T0 is the Curie-Weiss temperature. T0 is
slightly lower than the exact transition temperature Tc.
Ti4+
Ba2+
O2-
Figure 1.2 Crystal structure of BaTiO3. (Adapted from reference 59)
It is also known that the spontaneous polarization Ps and the spontaneous strain Xs
follow the relationship
2 S = QPX S (1.2)
where Q is the electrostrictive coefficients, Xs decreases almost linearly with
increasing temperature. With decreasing temperature from room temperature, BaTiO3
undergoes a series of complicated phase transitions. Figure 1.3 illustrates theses successive phase transitions.
6 Chapter 1 Introduction
Rhombohedral Orthorhombic Tetragonal Cubic
εa Dielectric Constant Dielectric Constant
εc
Temperature ºC
Figure 1.3 Various phase transitions in barium titanate, εa and εc are the permittivity along the a and c axis respectively (Adapted from reference 59)
An internal structure of spontaneously electrically polarized domains is a
characteristic feature of the ferroelectric phase. The planes along which individual
domains conjoin are termed domain walls, and the process of polarization reversal or
re-orientation under high fields is accomplished by the motion of existing walls, or by
the creation and motion of new domain walls.
For each ferroelectric species, conditions for the permitted orientations of Ps in the domain structure are dictated by the prototype symmetry, which also determines whether the state will be fully, partially, or nonferroelastic. If contiguous domains are strain-distinct, the conjoining ferroelectric-ferroelastic wall has only a limited family of possible orientations which are rigorously prescribed by the conditions for dimensional compatibility between the two domains along the wall plane. Usually the permitted wall is a common high-symmetry plane in the prototype, which is lost on the appearance of Ps.
The switching behavior of ferroelastic-ferroelectric walls depends markedly
on the magnitude and the nature of the spontaneous domain strains, which can be
7 Chapter 1 Introduction
different even among ferroelectrics in the same structure family. In the cases where
the strain is large, large switching can cause stress concentration that can facture the
crystal and fatigue effects are often evident in repeated cycle of switching of the
ferroelectric domains. This problem is especially acute in polycrystalline ceramic
ferroelectrics, which have to be poled to a high saturation remanence for piezoelectric
applications.
It is noted that since the domain wall is very narrow, the energy gained by
moving the wall just one lattice constant is necessarily very much smaller than the
wall energy itself, and thus true continuous sidewise motion of the wall, as occurs for
in many magnetic domain walls, is most unlikely. The model that does appear to
explain the behavior is that the apparent sidewise motion of walls is caused by the
nucleation and growth of step-like protrusions on the existing 180º walls. The narrow domain wall of ferroelectrics has especially important consequences in the major application of ferroelectric in capacitor dielectrics. Since critical size nuclei occur only very infrequently at low fields, reversible wall motion does not contribute in a major way to the dielectric response, and the capacitor engineer must manipulate the soft single-domain ferroelectric and the paraelectric permittivity to satisfy the application requirements.
1.4 Piezoelectricity
Crystals can be classified into 32 point groups according to their
crystallographic symmetry, and these point groups can be divided into two classes,
one with a center of symmetry and the other without. There are 21 point groups which
do not have a center of symmetry. In crystals belong to 20 of these point groups,
8 Chapter 1 Introduction
positive and negative charges are generated on the crystal surfaces when appropriate
stresses are applied. Theses materials are known as piezoelectrics.
The piezoelectric effect in quartz was discovered in 1880 by the brothers J.
Curie and P. Curie. When certain types of crystals are subjected to external
mechanical stress, the resulting strain causes a polarized state in the crystals and an
electric field is created. Conversely, if a crystal is polarized by an electric field, strains
along with the resulting stresses are created. Together, these two effects are known as
piezoelectric effect. The two aspects are sometimes distinguished as the positive and
reverse effects. In crystals that show piezoelectric properties, mechanical quantities
such as stress or strain, and electrical quantities such as electric field, electric
displacement (flux density) or polarization, are interrelated. This phenomenon is
called electromechanical coupling.
Among various piezoelectric phenomena, longitudinal and transverse effects
are particularly important. Longitudinal effect means that deformations take place parallel to the electric axis; whereas transverse effect means that deformations occur perpendicular to the electric axis. In practice the two types of effect take place at the
same time (Figure 1.4). One is proportional to the field strength and the other is
proportional to the square of the field strength. The former is the piezoelectric effect
in a strict definition, while the latter is sometimes distinguished as eletrostrictive
phenomenon.
9 Chapter 1 Introduction
T
+ - V P P V - +
T
P
- T + V T V + -
P (a) (b) P: Pressure T: Tension Figure 1.4 The piezoelectric effect: (a) longitudinal effect; (b) transverse effect. (Adapted from reference 60)
Piezoelectric ceramics have a complex multi-domain structure and exhibit
quite complex behavior. Figure 1.5 shows a plot of strains in the direction of the
applied electric field for a typical piezoelectric material. The state of the materials is
determined by its history, and this property is referred to as hysteresis. As the
hysteresis loop is very complex, new terminology is required to explain it. A model
shown is in Figure 1.6, which illustrates how ‘poling’ is undertaken on a multidomain
piezoelectric material.
Some of the terminologies that are commonly used in discussion of
piezoelectricity and their respective explanations are as follow:
1. Polarization. Polarization, denoted by P, is related to electric displacement (or electric flux density) D through the linear expression
+= ε EPD ioii (1.3) where the subscript i represents any of the three coordinates x, y and z, and εο is the
10 Chapter 1 Introduction permittivity of free space. In most piezoelectric materials, D and P are non-linear functions of E and depend on the history of the material. When the term εοE in the above expression is negligible compared with P (as in most cases for ferroelectrics), D is nearly equal to P.
Strain (S)
Electric field strength (E)
Figure 1.5 Strain characteristic for a typical piezoelectric ceramic. (Adapted from reference 60)
2. Permittivity. This parameter, denoted by ε, is defined as the incremental change
in electric displacement per unit electric field when the magnitude of the
measuring is very small compared with the coercive electric field, Ec.
3. Remnant polarization. The value of the polarization that remains after an
applied electric field is removed is defined as remnant polarization.
4. Poling and switching. As-fabricated piezoelectric ceramics have a
polycrystalline structure consisting of randomly distributed domains. Poling is a
process to align these randomly distributed domains at a d.c. electric field that is
higher than the coercive field (Ec). In order to explain this in details, let us look
at the grains of a crystal as shown in Figure 1.6. The crystal has been initially
11 Chapter 1 Introduction polarized in negative direction, with each domain being polarized more or less in downward direction. Thus the originally square ceramic block has become elongated vertically. If an electric field in positive upward direction is gradually applied, the block will contract firstly, since the field opposes the polarized direction. As the electric field is increased, some of the poles in the grains will begin to reverse in direction. At a certain voltage, the block will no longer be able to contract any further. This electric field is called the coercive field, Ec.
If the field strength is further increased, the ceramic block will start to expand. When all the poles have been reversed, the block can’t expand further; the field for this condition is indicated as Emax. If the electric field is then reduced, the strain will keep decreasing until the electric field reaches zero. In the final state shown, the poles in all grains are reversed when compared with the initial state, and the block has been polarized in the positive direction.
12 Chapter 1 Introduction
(1) E = 0 (2) E = Ec
(3) E = Emax
(4) E = E (5) E = 0 c
Polarization
Figure 1.6 Schematic illustration of strains induced by pole reversals in a ferroelectric ceramic materials. (Adapted from reference 60)
Characterization of piezoelectric effects is usually based on several parameters, which are described as below:
1. Piezoelectric strain constant. As mentioned above, the strain and applied voltage
are proportional in a polarized crystal. This relationship, if we ignore the
hysteresis effect, can be expressed as
Δl ⋅= Ed for positively polarized state (1.4a) l
Δl ⋅−= Ed for negatively polarized state (1.4b) l
The proportionality constant d is called the piezoelectric strain constant.
2. Poisson’s ratio. Poisson’s ratio is a parameter which indicates relative
deformations in longitudinal and transverse directions. Specifically, it is the ratio
of transverse elongation to longitudinal contraction when a pressure is applied to a
solid at a constant voltage, or
13 Chapter 1 Introduction
S σ E = 31 (1.5) S33
The superscripts and subscripts in the terms are commonly used when tensor
expressions one necessary to describe piezoelectric phenomena. A superscript
denotes a non-varying parameter during state changes; thus σE is the Poisson’s
ratio when the applied voltage is kept constant. Subscript indicates axis direction
for cause and effect. The number 1, 2 and 3 corresponds to axis x, y and z
respectively. Thus a pressure in the z-direction (cause) creates a strain Δz/zo in the
z-direction (effect), represented by S33. Similarly, S31 is the strain in the x-
direction caused by a pressure in the z-direction. Since normally S32 = S31, S31 also
E represents S32 when it exists. The typical Poisson’s ratio σ of piezoelectric
materials is ~0.3.
3. Directionality of piezoelectric strain constant. If deformations are caused by an
electric field, but can’t be determined by the Poisson’ ratio, it is necessary to use
the piezoelectric strain constants, which possess directional qualities as well. The
strain constant in the z-direction (for longitudinal effect) is usually represented by
d33, defined by
Δz = 33 Ed Z (1.6) z0
The elongation in the x-direction (for transverse effect) is given by d31,
Δx ⎛ Δy ⎞ ⎜ ⎟ ⎜= ⎟ = 31Ed Z (1.7) x0 ⎝ y0 ⎠
4. Voltage output coefficients. The reverse piezoelectric effect is formulated as
z −= 33TgE z (1.8)
z −= TgE ,31 yx (1.9)
14 Chapter 1 Introduction
When measuring the induced voltage, from which the electric field is computed
with the sample dimension, care must be taken in order not to generate a current.
The proportionality constants g33 and g31 are called the voltage output constants.
The g constants are closely related to the d constants, and have simple
relationships:
d33 g33 = T (1.10) ε 33
d31 g31 = T (1.11) ε 31
where ε is the permittivity (dielectric constant) of the ceramic.
1.5 Lead Zirconate Titanate Oxide (PZT)
Piezoelectric Pb(Ti,Zr)O3 (PZT) solid solutions have been widely used because of their superior ferroelectric and piezoelectric properties. A phase diagram for PZT system (PbZrxTi1-xO3) is shown in Figure 1.7. Lead titanate is a tetragonal ferroelectric of perovskite structure. With increasing Zr content, x, the tetragonal distortion decreases and at x > 0.52 the structure changes from the tetragonal phase to another ferroelectric phase of rhombohedral symmetry. The line dividing these two phases is called the morphotropic phase boundary (MPB). The boundary composition is considered to have both tetragonal and rhombohedral phases.
15 Chapter 1 Introduction
Cubic C º
Tetragonal Rhombohedral
Temperature Morphotropic phase boundary
Mol %, PbZrO3
Figure 1.7 Phase diagram of lead zirconate titanate (PZT). (Adapted from reference 59)
1.6 Current status and major problem in PZT research
The problem of PZT used in MEMS devices and thin-film actuators61 is integration difficulties. The main impetus for it integration onto silicon was the prospect of non-volatile, radiation-robust memories.
Use of PZT in ferroelectric memory also suffered the same problem. The initial barrier to the development of ferroelectric memories was the necessity of making them extremely thin films because the coercive voltage is typically of the order of several kV/cm. With today’s deposition techniques this is no longer a problem, and now high-density arrays of nonvolatile ferroelectric memories are commercially available. However, reliability remains a key issue.
16 Chapter 1 Introduction
The main growth phenomena of PZT can be roughly understood in terms of a few key features of the PbO-ZrO2-TiO2 system, independently of the deposition method.
(1) Nucleation and growth of the perovskite phase require a rather precise
stiochiometry, otherwise competing phases with fluorite and pyrochlore
structures nucleate.62
(2) Lead ions or PbO molecules that are not incorporated into the perovskite
lattice exhibit high diffusivities and volatility above 500°C. The PbO vapor
pressure above PbO is approximately 100 times larger than PZT and amounts
to 1.1 Pa at 600°C.63
(3) The activation energy for nucleation of the perovskite phase (4.4 eV/unit cell)
is considerably larger than for its growth (1.1 eV).64
The growth of good quality PZT thin films still needs more efforts. Reproducible film quality is certainly possible if an industrial approach is adopted. The electrodes below the PZT thin film play a very important role for seeding the correct phase and the film texture. Without reproducible electrode quality, no reproducible PZT quality is achieved. Good piezoelectric properties have only been obtained for deposition temperatures higher than 550 - 600°C. This could become a big issue if direct integration of PZT thin films onto integrated circuits is the goal. There are still challenges in achieving high piezoelectric properties for PZT thin films. One problem of judging piezoelectric performance is the fact that we do not know exactly how to derive thin-film properties from known bulk ceramic properties. Microstructural differences, defects due to lower growth temperature and interface effects at the electrode certainly play a role. Lattice and domain wall contributions need to be
17 Chapter 1 Introduction considered differently. For a final statement, there is also not enough experience with doping of PZT thin films. In bulk ceramics, piezoelectric properties can be very much improved through doping. The question is open whether this also works with thin films.
18 Chapter 1 Introduction
1.7 References
1. D. Brewster, Edinburgh J. Sci. 1, 208 (1824).
2. G. Bush and P. Scherer, Naturwiss 23, 737 (1935).
3. G. Bush, Helv. Phys. Acta 11, 169 (1939).
4. E.Wainer and N. Salomon, National Lead Co. Reports. No. 8, 9, 10 (1938-1943).
5. S. Ogawa, J. Phys. Soc. Jpn., 1, 32 (1946).
6. B.M. Wul and I.M. Goldman, Doki. Akad. Nauk. SSSR 46, 154 (1945).
7. B. Gray, U.S. Pat. No. 2 486 560, September 20, 1946
8. G. Shirane and A. Takeda, J. Phys. Soc. Jpn. 7, 5 (1952).
9. G. Shirane, K. Suzuki and A. Takeda, J. Phys. Soc. Jpn. 8, 615 (1954).
10. E. Sawaguchi, J. Phys. Soc. Jpn. 8, 615 (1953).
11. B. Jaffe, R.S. Roth and S. Marzullo, J. Appl. Phys. 25, (1954).
12. B. Jaffe, R.S. Roth and S. Marzullo, J. Res. Natl. Bur. Stand. 55, 239 (1955).
13. W. Cochran, Phys. Rev. Lett. 3, 412 (1959).
14. W. Cochran, Adv. Phys. 9, 387 (1960).
15. K. Kim, C. Kim and S. Lee, Micro. Eng. 66, 662 (2003).
16. T. kiguchi, N. Wakiya, K. Shinozaki and N. Mizutani, Micro. Eng. 66, 708
(2003).
17. J. Junquera and P. Ghosez, Nature 422, 506 (2003).
18. H. Cheng, Y. Chen, C. Chou, K. Chang, C. Hou and I. Lin, J. Appl. Phys. 87,
8695 (2000).
19. H. Funakubo, M. Aratani, T. Oikawa, K. Tokita and K. Saito, J. Appl Phys. 92,
6768 (2002).
20. J. Scott and C. Paz de Arujo, Science 246, 1400 (1989).
19 Chapter 1 Introduction
21. K. Saito, T. Kurosawa, T. Akai, T. Oikawa and H. Funakubo, J. Appl. Phys. 93,
545 (2003).
22. S.R. Gilbert, S. Hunter, D. Ritchey, C. Chi, D.V. Taylor, J. Amano, S. Aggarwal,
T.S. Moise, T. Sakoda, S.R. Summerfelt, K.K. Singh, C. Kazemi, D. Carl and B.
Bierman, J. Appl. Phys. 93, 1713 (2003).
23. M. Tsukada, H. Yamawaki and M. Kondo, Appl. Phys. Lett. 83, 4393 (2003).
24. J. F. Scott, Ferroelectric Memories (Springer, Berlin, 2000).
25. Z. G. Zhang, D. P. Chu, B. M. McGregor, P. Migliorato, K. Ohashi, K. Hasegawa,
and T. Shimoda, Appl. Phys. Lett. 83, 2892 (2003).
26. F. Yan, Y. Wang, H.L.W. Chen, C.L. Choy, Appl Phys. Lett. 82, 4325 (2003).
27. K. Tokita, M. Aratani and H. Funakubo, Appl. Phys. Lett. 82, 4122 (2003).
28. Auciello, J.F. Scott and R. Ramesh, Phys. Today 51, 22 (1998).
29. F. Mcnally, J.H. Kim and F.F. Lange, J. Mater. Res. 15, 1546 (2000).
30. Y. Zhu, M. Shen, R. Xu and W. Cao, Surf. Coat. Tech. 160, 277 (2002).
31. H. Yang, B. Chen, K. Tao, X. Qiu, B. Xu and B. Zhao, Appl. Phys. Lett. 83, 1611
(2003).
32. S.T. Chang and J.Y. Lee, Appl. Phys. Lett. 80, 655 (2002).
33. N. Maffei and S.B. Krupanidhi, Appl. Phys. Lett. 60, 781 (1992).
34. F. Mitsugi, T. Ikegami, K. Ebihara, J. Narayan and A.M. Grishin, Mat. Res. Soc.
Symp. 617, 3231 (2000).
35. T. Kanki, Y. Park, H. Tanaka, T. Kawai, Appl. Phys. Lett. 83, 4860 (2003).
36. J. R. Contreras, H. Kohlstedt, U. Poppe, R. Waser, C. Buchal and N. A. Pertsev,
Appl. Phys. Lett. 83, 4595 (2003).
37. S. Koo, S. Kkhartsev, C. Zetterling, A. Grishin and M. Ostling, Appl. Phys. Lett.
83, 3975 (2003).
20 Chapter 1 Introduction
38. A.M. Grishin. S.I. Khartsev and P. Johnsson, Appl. Phys. Lett. 74, 1015 (1999).
39. W. Wu, K.H. Wang, C.L. Mak, C.L. Choy and Y.H. Zhang, J. Vac. Sci. Technol.
A 18, 2412 (2000).
40. M.H. Frey and D.A. Payne, Appl. Phys. Lett. 63, 2753 (1993).
41. A. Lin, X. Hong, V. Wood, A. A. Verevkin, C. H. Ahn, R. A. McKee, F. J.
Walker, and E. D. Specht, Appl. Phys. Lett. 78, 2034 (2001).
42. F. Cracium, P. Verardi, M. Dinescu, C. Galassi and A. Costa, Sens. & Actuat. 74,
35 (1999).
43. M. Kim, J. Choi, S. Yoon, K.H. Yoon and H. Kim, Jpn. J. Appl. Phys. 41, 3817
(2002).
44. N. Ledermann, P. Muralt, J. Baborowski, S. Gentil, K. Mukati, M. Cantoni, A.
Seifert and N. Setter, Sens. & Actuat. 105, 162 (2003).
45. W. gong, J. Li, X. Chu, Z. Gui and L. Li, Jpn. J. Appl. Phys. 42, 1459 (2003).
46. L. Lian and N.R. Sottos, J. Appl. Phys. 87, 3941 (2000).
47. S. Kalpat and K. Uchino, J. Appl. Phys. 90, 2703 (2001).
48. H. Maiwa, S. Kim, N. Ichinose, Appl. Phys. Lett. 83, 4396 (2003).
49. K. Yao, Koc B., Uchino K., IEEE Trans on Ultrasonics, Ferroelectrics and
Frequency Controls. 48, 1066 (2001).
50. K. Yao, W. Zhu, Uchino E., Z. Zhang, L.C. Lim, IEEE Trans on Ultrasonics,
Ferroelectrics and Frequency Controls. 46, 1020 (1999).
51. UT. Morita, M.K. Kurosawa and T. Higuchi, Sens. & Actuat. 83, 225 (2000).
52. T.Y. Jiang, T.Y. Ng and K.Y. Lam, Sens. & Actuat. 84, 81 (2000).
53. P. Muralt, IEEE Trans. Ultrason. Ferroelectr. Frequency Control 47, 903 (2000).
54. P. Muralt, J. Micromech. Microeng. 10, 136 (2000).
21 Chapter 1 Introduction
55. J. Tsaur, Z.J. Wang, L.L. Zhang, M. Ichiki, J.W. Wan and R. Maeda, Jpn. J. Appl.
Phys. 41, 6664 (2002).
56. B. Vilquin, R. Bouregba, G. Poullain, H. Murray, E. Dogheche and D. Remiens,
J. Appl. Phys. 94, 5167 (2003).
57. C.C. Chang and C.S. Tang, Sens. & Actuat. A 65, 171 (1998).
58. W.C. Shih and M.S. Wu, IEEE Trans. Ultrason. Ferroelectr. Freq. Control. 45,
305 (1998).
59. C.Z. Rosen, B.V. Hiremath and R. Newnham, Piezoelectricity, Published by
American Institute of Physics (1992).
60. T. Sashida and T. Kenjo, An introduction to ultrasonic motors, Published by
Oxford Science Publications (1993).
61. A.M. Flynn, L.S. Tavrow, S.F. Bart, R.A. Brooks, D.J. Ehrlich, K.R.
Udayakumar and L.E. Cross, J. Microelectromech. Syst. 1, 44 (1992).
62. R.D. Klissurska, I.M. Reaney, C. Pawlaczyk, M. Kosec, K.G. Brooks and N.
Setter, J.Am.Ceram. Soc. 78, 1513 (1995).
63. K.H. Kartl and H. Rau, Sold State Commun. 7, 41 (1969).
64. C.K. Kwok and S.B. Desu, J. Mater. Res. 9, 1728 (1994).
22 Chapter 2 Apparatus and experiment procedures
Chapter 2 Apparatus and experiment procedures
2.1 Sample preparation
There are several options for fabricating ferroelectric films, such as pulsed laser deposition, sputtering, screen printing, sol-gel and metalorganic chemical vapor deposition (MOCVD).1-7 The deposition techniques used in this thesis are pulsed laser deposition (PLD) and sol-gel deposition. Both techniques have been widely used for the fabrication of metal oxide films with complex composition due to their advantages in chemical composition control.8-9
2.1.1 Pulsed laser deposition
PLD technique was first developed in 1960’s.10 The technique was not widely used during that period of time as the fabricated film had poor quality due to the lack of suitable laser source. Only after the advancement of the laser technology in the
1970’s, when powerful UV laser using excimer gases that has shorter laser pulse (<
10 ns) and higher energy (>106 W) was available, congruent vaporization and the ignition of the plasma became possible. Subsequently, the successful synthesis of
YBa2Cu3O7-δ in 1987 highlighted the advantages of PLD in fabricating thin films of multi-component complex materials.11
Compared to other deposition techniques PLD has many advantages12, such as:
1) good ability to fabricate thin film with very complex compositions from bulk
materials;
2) relatively high growth rate of about 1 – 5 Å per pulse;
3) decoupled laser energy source from deposition environment;
4) no ultrahigh vacuum requirement;
23 Chapter 2 Apparatus and experiment procedures
5) wide range of ambient reactive gas pressures (from 10-7 to 1 mbar);
6) relatively simple fabrication setup for in-situ growth of different material with
multilayer structure and
7) reduced film contamination due to the use of light for promoting ablation.
PLD also has its limitations, such as
1) relative small deposition area and
2) existence of large particulates that was ejected from the target.
However, deposition area can be increased to a certain extent by sweeping laser beam across target, while large particulates can be easily fiiltered by using a suitable spinning disc containing an aperture synchronized to the laser pulses.
The basic principle of PLD deposition process can be explained as follows. A high energy laser beam is first focused on to target surface located inside a vacuum chamber. The focused laser results in a congruent evaporation within a small target area and a consequent explosive evaporation of a thin layer before it has time to disproportionate. The evaporated ion species are then transferred to the substrate located in front of the target where the thin film is formed due to condensation.
Figure 2.1 show the diagram of a typical PLD system. We used a high energy
KrF excimer laser (pulse duration 30 ns, wavelength 248 nm, Lambda Physik Lextra
200). The laser is first focused through a focusing lens outside the vacuum chamber.
The target holder is customized such that it can hold up to 4 different targets inside the chamber. This would enable us to grow different thin film layers without breaking the vacuum by a rotating the holder to the desired target.13 This would also save us a lot of time needed to change a target. The target rotates around its axis during deposition to minimize the large particulate splashing effect and to achieve a more uniform ablation of the target. The distance between the target and the sample is 3 to
24 Chapter 2 Apparatus and experiment procedures
5 cm. The chamber can be pumped down to a vacuum of around 1×10-6 mbar by a turbo molecular pump backed by a rotary pump.
Mirror KrF Excimer Laser (248 nm)
Lens
Vacuum gauge Heater & substrate
Plume Targets Turbo Pump & Rotary Pump View Port Gas Inlet Motor
Figure 2.1 A schematic drawing of the pulsed laser deposition (PLD) system.
All substrates were first cleaned using nitric acid in an ultrasonic cleaner for 5 min to remove any natural oxide layer or oxide contaminant on the surface and subsequently cleaned with de-ionized water, acetone and ethanol. The cleaned substrates were always kept in alcohol to prevent re-oxidation or dust before being transferred to the vacuum chamber. The substrates were adhesively attached to sample holder (resistive heater) by applying a thin layer of silver paste. The temperature of the substrate was controlled by Eurotherm temperature controller. The temperature was gradually increased from room temperature to desired temperatures of 300°C to
700°C depending on the materials deposited. Ambient reactive oxygen gas was
25 Chapter 2 Apparatus and experiment procedures introduced into the chamber through a small nozzle located near the substrate. The flow rate and the pressure of the gas was controlled through a series of gas valves, needle valves and block valves placed outside of the vacuum chamber.
There are several processing parameters that affect the deposition and growth structure of the thin films. They are listed as bellow:
1) Laser power, wavelength, pulse length and repetition rate;
2) Interaction of laser with target such as laser density, spot size, angle of
incident and physical properties of the materials;
3) Interaction of the resulted plasma gas with ambient gas, target to substrate
distance, angle of incident between the plasma and the substrate;
4) Substrate lattice parameter, thermal conductivity, thermal expansion
coefficient and temperature.
However these parameters are almost independent on one another, which makes PLD very versatile in producing high quality thin films. Systematic experiments were carried out to obtain the optimum processing parameters for all the thin films in this study.
In most cases, the films were further in-situ annealed in the chamber after deposition to improve the crystallinity, and to increase oxygen content of the films.
This step is crucial especially for oxide thin films because they tend to loose oxygen during the deposition.
2.1.2 Sol-gel deposition
Sol gel deposition is another useful technique14 to fabricate some ferroelectric films. The purpose for using sol gel deposition in the present work is to make a comparison between solution deposition method and PLD method, by having thin
26 Chapter 2 Apparatus and experiment procedures films with different morphology and structure. This would help us to have a better understanding of the underlying physics between microstructure and electrical properties of ferroelectric films.
Sol-gel process starts with liquid solution of metalorganic precursors, which are hydrolyse and polymerize to become gels of different forms. The starting materials used to prepare "sol" are usually inorganic metal salts or metal organic compounds such as metal alkoxides. A “sol” prepared in this was a colloidal suspension. When the "sol" is cast into a mold, a wet "gel" would form. With further drying and heat-treatment, the "gel" is converted into dense ceramic or glass articles. If the liquid in a wet "gel" is removed under a supercritical condition, a highly porous and extremely low density material called "aerogel" can be obtained. When the viscosity of a "sol" is adjusted into a proper viscosity range, ceramic fibers can be drawn from the "sol". Thin films can be produced with such
“sols” on various substrate by spin-coating or dip-coating.
The advantages of sol-gel process are:
1) room temperature deposition
2) large area in one single process
3) no vacuum chamber requirement
4) simple processing steps
5) excellent stoichiochemistry control
These advantages have made sol-gel a very attractive deposition technique for growing large area and large quantity of ceramic films. There are also limitations for the sol-gel deposition process:
27 Chapter 2 Apparatus and experiment procedures
1) suitable metalorganic precursors might not be available for fabrication of some of the oxide films
2) substrate smoothness is stringently required for achieving crack free spin- coated films.
28 Chapter 2 Apparatus and experiment procedures
2.2 Crystal and microstructure characterizations
2.2.1 X-ray diffractions
Two x-ray diffraction characterizations were carried out to determine the crystal structure and quality of the thin film deposited. Gonio (θ-2θ) scans were used to identify the crystalline phase and determine the out-of-plane crystal orientation of the deposited film; phi (φ) scans were used to determine the in-plane crystal orientation.
In Gonio scan, an x-ray beam was incident on sample with an angle to the normal of the film surface. The x-ray was then diffracted by the crystal lattice of the samples and the diffracted x-ray intensity was measured. A schematic diagram of the gonio scan is shown in Figure 2.2.
Philips APD 1700 X-ray diffractometer with Cu Kα radiation was used for the gonio scans of the thin films. The films were first mounted onto a glass slide. The x- ray was produced by impinging a copper target with an electron beam of 30 kV with a current of 20 mA. An incident slit of 1° and a receiving slit of 2 mm were used to ensure that the received signals are the diffracted x-rays caming from only the samples. The parameters used for gonio scan were as follow:
1) 2θ range: 20° to 60°
2) Scan Step: 0.05°
3) Collection time: 1 second per step
The results were then compared to the standard powder diffraction file15 and the publications in the open journals, for determination of the phase composition, crystal orientation and lattice parameters of the films deposited.
29 Chapter 2 Apparatus and experiment procedures
Detector
X-Ray Source Receiving Slit Aperture Slit
2θ
Sample θ
Figure 2.2 X-ray diffraction (XRD) θ-2θ scan.
Phi (φ) scan was carried out using a Bruker D8-ADVANCE XRD machine.
The sample was mounted on stage that can be rotated along the axis normal to sample.
The x-ray incident angle and the detector were fixed, and the sample was rotated along the normal axis of the film surface and the diffraction pattern was collected.
The incident angle and detect angle were calculated based on the interested planes.
The X-ray used was produced by impinging an electron beam onto a copper target. The electron beam energy ranged from 20 kV to 40 kV with a current of 5 mA to 20 mA selected, depending on the sample quality and diffraction peak. The parameters of the phi scan were as follow:
1) Phi angle: 0º to 360º
2) Incident and detector angle: depending on the lattice parameters
3) Step size: 1º
4) Step time: 10 second per step
30 Chapter 2 Apparatus and experiment procedures
2.2.2 Scanning electron microscope and atomic force microscope
In order to investigate the morphology and microstructure of the ferroelectric films deposited, we employed field emission scanning electron microscope
(FESEM)16 and atomic force microscope (AFM).17
SEM was used to examine the surface morphology and the cross-section profile.1-3 In SEM characterization, electrons are extracted from a very fine tungsten tip by a high voltage of several thousands of voltage. The electrons are then focused through a series of magnetic lenses onto samples. The bombardment of the electrons will cause backscattered electrons and secondary electrons to be emitted from sample surface. These backscattered and secondary electrons are collected and analyzed.
Since the backscattered and secondary electrons collected usually come from the surface of films, the image formed would normally reflect the surface morphology of samples observed (Figure 2.3).
SEM equipment used in this study is JSM-6700F by JEOL Ltd., Tokyo, Japan.
It is equipped with a field emission electron gun of 5 kV to 20 kV and a current of 5 mA to 20 mA, which can be selected depending on condition of the samples. Since ferroelectric materials are insulators, a very thin layer of gold or platinum was coated onto surface of the sample for better conduction of electrons before. The conductive coating prevents the build-up of electric charge on the sample to ensure the quality of
SEM images.
31 Chapter 2 Apparatus and experiment procedures
Electron Gun
Condenser Lens
Objective Lens
Final Aperture Scanning Coil
Scintillator
Sample
Figure 2.3 Schematic drawing of a scanning electron microscope (SEM).
AFM technique is also widely used to characterize surface morphology of thin samples.4 In AFM characterization, a very sharp probe, in our case which is made of silicon nitride, is positioned on the surface of sample. A constant force is then maintained between the probe and the sample surface while the probe is scanned across the sample surface. By monitoring the probe, a 3-dimensional image of the sample surface can be constructed. The constant force is maintained constant by measuring the level of the reflected laser from the probe with a “light lever” sensor.
The signal is then fed into a feedback unit that control the piezoelectric driver unit of the probe. A diagram of the AFM is shown in Figure 2.4. The advantage of AFM over
SEM is that the sample does not need to be conductive. However the resolution of the
AFM highly depends on the surface roughness of the sample and the scanning areas of AFM are usually smaller compare to SEM.
32 Chapter 2 Apparatus and experiment procedures
There are two operation modes, contact and tapping mode that can be used for scanning the surface of samples. Contact mode means that the tip of probe is scanned at a near proximity of sample surface. The feedback systems of the AFM monitor the force between the tip and the sample and adjust the piezoelectric driving unit accordingly. The advantage of contact mode is that one can obtain a better contrast and closer morphology details of the sample surface. However since the tip is placed very close to the sample, there are possibilities for tip to pick up sample debris and to lose contact when the surface is too rough. In tapping mode, the tip is oscillated at a frequency near the resonance of the AFM probe. The tip is then placed at a much elevated position on top of the sample. The AFM feedback systems monitor the changes in the resonance frequency and adjust the position of the tip accordingly. The advantage of the tapping mode is that one can prevent any scratching of the sample and maintain contact with the sample even if the surface roughness is large.
We used AFM to characterize the average surface roughness and the grain size distribution of our thin films. The AFM machine used in the present study is D3000 system manufactured by Digital Instruments.
33 Chapter 2 Apparatus and experiment procedures
Feedback Loop Electronics:
Laser
Scanner
Cantilever & Tip Photo Diode
Sample
Figure 2.4 Schematic drawing of an atomic force microscope (AFM) operating in tapping mode. (Adapted from reference 17)
2.3 Electrical and ferroelectric characterizations
2.3.1 Impedance, dielectric and loss tangent measurement
The dielectric properties of the ferroelectric thin films prepared were characterized using an HP 4194A impedance analyzer.18 The measuring frequency range was from 1 kHz to 1 MHz.1,3 The thin film samples were prepared with a capacitor sandwich structure. Ferroelectric films were deposited onto bottom electrodes of either metallic conductor such as platinum or conductive oxide such as
YBCO or LSMO. Part of the ferroelectric film was chemically etched to expose the bottom electrode. Au (gold) was then sputtered on to the ferroelectric thin films as top electrodes.
34 Chapter 2 Apparatus and experiment procedures
A probe station was used for the electrical measurement of the samples. The probe station has two traveling manipulators which can be adjusted in x, y and z directions with microscrews. A tungsten tip is attached at the end of each manipulators to provide electrical contact to samples. An optical microscope is equipped above the probe station to provide a magnified view for the positioning of the tungsten probes onto samples.
2.3.2 Ferroelectric loop measurement
Ferroelectric loop, polarization versus electric field (P-E), was measured using a standardized ferroelectric test system RT6000S manufactured by Radiant
Technologies.19 The measurement was performed by applying an electric field across the ferroelectric sample. The current response was then amplified and integrated to calculate the total charge that was accumulated across the sample.1,3
Two different modes are commonly used in the P-E loop measurement of a ferroelectric sample namely Sawyer Tower and Virtual Ground modes. For the
Sawyer Tower mode (Figure 2.5), a capacitive voltage divider comprised of the test sample and a sensing capacitor (usually supplied by the user) is used. The voltage generated across the sense capacitor is proportional to the charge stored in the sample.
Due to its simplicity and low cost, the Sawyer Tower circuit has been the common method for characterizing ferroelectric devices. However, it is susceptible to significant errors from parasitic elements, particularly those associated with the sense capacitor hook up, and is limited by the accuracy to which the sense capacitor value is
35 Chapter 2 Apparatus and experiment procedures known.
Voltage Drive
Sample
Voltage out
Sawyer Tower C parasitic C sense
Figure 2.5 Equivalent circuit of Sawyer Tower mode.
The Virtual Ground mode (Figure 2.6) measures the charge stored in the ferroelectric sample by integrating the current required to maintain one terminal of the sample at zero volt – hence the term “virtual” ground. By eliminating the external sense capacitor, this circuit drastically reduces the effects of parasitic elements. The
Virtual Ground mode allows device characterization with an improved accuracy and allows results obtained from different test set up to be compared with confidence.
Therefore the Virtual Ground mode was used in all the characterization for our ferroelectric thin film samples.
36 Chapter 2 Apparatus and experiment procedures
Voltage Drive
Sample Current Integrator
Voltage out Current C parasitic Amplifier Integrating Capacitor
Figure 2.6 Equivalent circuit of Virtual Ground mode.
2.3.3 Piezoelectric constant measurement
Piezoelectric vibration was measured using a laser scanning vibrometer (OFV-
056 and OFV-3001-SF6 from PolyTec, Berlin). Characterization of piezoelectric constant using laser scanning vibrometer proposed by Yao et al.20 The experiment setup designed to measure the frequency shift of the reflected laser beam from sample surface. An alternating current (AC) signal was applied across the ferroelectric film to excite a vibration at the frequency of the AC signal. The reflected laser signal was then collected and transformed into data of frequency domain. Thus, the vibration signal can be sensitively distinguished from background noise. In our experiment, the scanning area covered both the top electrodes and the surrounding PZT films so that the movement of substrates could be observed simultaneously. A schematic illustration of the setup is shown in Figure 2.7.
37 Chapter 2 Apparatus and experiment procedures
Sample Laser Beams Laser Vibrometer
Stage
Amplifier
PC Reference Signal AC Signal Generator
Figure 2.7 Schematic drawing of the piezoelectric constant measurement setup using a laser scanning vibrometer.
38 Chapter 2 Apparatus and experiment procedures
2.4 References
1. S. Otsubo, T. Maeda, T. Minamikawa, Y. Yonezawa, A. Morimoto and T.
Shimizu, Jap. J. Appl. Phys. Part 2 29 133, (1990).
2. H. Buhay, S. Sinharoy, W.H. Kanser, M.H. Francombe, D.R. Lampe and E.
Stepkee, Appl. Phys. Lett. 58, 1470 (1991).
3. D. Roy, S.B. Krupanidhi and J.P. Dougherty, J. Appl. Phys.69, 7930 (1991).
4. K.D. Budd, S.K. Dey and D.A. Payne, Br. Ceram. Proc. 36, 107 (1985).
5. S.B. Majumder, D.C. Agrawal, Y.N. Mohapatra and V.N. Kulkarni, Int. Ferro. 9,
217 (1995).
6. M. Ichiki, J. Akedo, A. Schoroth, R. Maeda, Y. Ishikawa, Jpn. J. Appl. Phys. 36,
5815 (1997).
7. C. Lucat, F. Menil and R. VonderMuhll, Mea. Sci. & Tech. 8, 38 (1997).
8. Hubler, G. K., in Pulsed Laser Deposition of Thin Films, edited by D. B. Chrisey
and G. K. Hubler, (John Wiley & Sons Inc., New York, 1994), p. 327.
9. R. F. Haglund, Jr., Laser Ablation and Deposition, edited by J. C. Miller and R. F.
Haglund, Jr., (Academic Press, London, 1998), p.15.
10. H.M. Smith and AF Turner, Appl. Opt. 4, 147 (1965).
11. D. Dijkkamp, T. Venkatesan, X. D. Wu, S. A. Shaheen, N. Jisrawi, Y. H. Min-Lee,
W. L. McLean, and M. Croft, Appl. Phys. Lett. 51, 619 (1987)
12. S. Ghosh, Solid State Comm. 116, 585 (2000).
13. S.Y. Xu, C.W. Tan and C.K. Ong, Meas. Sci.& Tech. 9, 1912 (1998).
14. S. Yu, K. Yao, S. Shannigrahi and E.H. Tay, J. Mater. Res. 18, 737 (2003).
15. Powder diffraction file: PDF-2 database (Computer file), International Center for
Diffraction Data, 1996-2000
39 Chapter 2 Apparatus and experiment procedures
16. B.L. Gabriel, SEM: a user’s manual for materials science, (American Society for
Metals, 1985).
17. Scanning Probe Microscopy Training Notebook, Digital Instrument (1999).
18. HP 4194A Impedance Analyzer Operating Manual, Hewlett Packard (2000).
19. RT6000S Basic Operating Manual, Radian Technologies (2002).
20. K. Yao and F.E.H. Tay, IEEE Trans on Ultra. Ferro. Freq. Ctrl. 50, 113 (2003).
40 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
Chapter 3 Early growth stage of laser ablated epitaxial PZT film
3.1 Introduction
Piezoelectric materials, in particular lead-zirconate-titanate Pb(ZrxTi1-x)O3
(PZT), have shown very promising properties for various applications such as non- volatile memory elements,1,2 infrared sensors,3 surface acoustic wave device4 and actuators.5 As is well known, besides Pulsed Laser Deposition (PLD) technique, approaches, such as sol-gel, rf sputtering, metalorganic chemical vapor deposition
(MOCVD) and jet-mould system, have also been used to fabricate piezoelectric ceramic films.6-8 The merits for using PLD can be found in section 2.1.1.9-10
In the present work, perovskite conductive oxides were chosen as bottom electrode instead of conventional metal layer. Recently this issue has been extensively studied, and results of which have clearly shown the advantages in structure compatibility, chemical stability and fatigue-resisting property, by replacing metallic electrodes with conducting oxides like YBa2CuO7, (La,Sr)MnO3 (LSMO), LaNiO3
11-13 and SrRuO3. Here both PZT and LSMO have perovskite structure with similar lattice constant in their a/b planes (a=b=3.85 Å for LSMO, a=b=4.04 Å for PZT), and both require oxygen as reactive gas during PLD process. We believe that the structural compatibility and similarity between them will not only improve the device performance, but also make it possible to realize novel all-(perovskite) oxide-devices that utilize the merits of high temperature superconductor (HTS), ferroelectric and semiconducting properties.
In order to obtain high quality PZT films, it is essential to understand the correlation between microstructures and properties of the films. The final microstructures of the PZT film, such as crystalline structure, orientation, grain size,
41 Chapter 3 Early growth stage of laser ablated epitaxial PZT film grain boundaries, second phases, impurities, defects and surface roughness, are to certain level determined by the growth mode at the early stage. This chapter reports the investigations on the early growth stage of a (001)-oriented epitaxial
23-25 Pb(Zr0.52Ti0.48)O3 film on (100) SrTiO3 (STO) substrate. Previous reports had shown that PLD PZT film on LSMO film showed columnar growth structure, thus it is of interest to us to investigate if there is a difference in PLD PZT growth mechanism on atomic smooth surface. Because there are no commercially available
LSMO substrates, we chose STO as the substrate since it also has the perovskite structure with quite good lattice match (a=3.905Å) with LSMO. The second reason is that the surface roughness of the LSMO film deposited using PLD is typically 5-10 nm, which does not fulfill our requirement for studying the initial growth of PZT film.
By depositing PZT on STO substrate, which has an atomic-smooth surface, we will have a direct and clear view of the early growth stage of PZT film on lattice-matched substrate. Since STO and LSMO have very close lattice constants, it is believed that the study will offer valuable clues on the growth mechanism of epitaxial PZT films on
LSMO.
3.2 Experimental procedures
The PLD system employed in this experiment has been employed to fabricate various high quality epitaxial thin films.14-19 The details of the PLD setup have been discussed in Section 2.1.1. The target used was a sintered PZT ceramic, which had a composition at the morphotropic phase boundary (Pb(Zr0.52Ti0.48)O3). The target-to- substrate distance was 3.5 cm. The incident angle of the laser beam with respect to the target surface was 45o.
42 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
Polished (100) STO single crystal substrate have a size of 10 mm × 10 mm ×
0.5 mm. Before deposition, all STO substrates were cleaned in HF (ph=4.5) aqueous solution for 1 minute by ultrasonic vibration to remove residual oxide layer on the surface and to smoothen the substrate surface. A substrate temperature of 650°C was used for all the samples.
The deposition of PZT layer was carried out at a laser fluence of 3.0 J/cm2 and frequency of 5 Hz, under 0.11 mbar of ambient oxygen pressure. The parameters were used for all the samples. A series of samples with different deposition time, i.e.
1, 3, 6, 10, 20, 30 and 40 min, and therefore different thicknesses, were prepared. A small area at one corner of the samples was covered during the deposition such that a stage is available to measure the film thickness using a surface profiler. After the deposition, the samples were annealed under 1 bar of oxygen for 30 min at 700°C before they were naturally cooled down to room temperature.
The crystalline phases and orientation of the films were characterized using X- ray diffraction (XRD) θ-2θ scan and θ-scan (rocking curve scan). The morphology of the films was studied using an atomic force microscope (AFM, Nanoscope IIIa) operating under tapping mode. A high-resolution transmission electron microscopy
(HRTEM TEM, Philips CM 300) was used to study the microstructure and interface between PZT film and STO substrate. Cross-section specimens for HRTEM studies were prepared using standard sample preparation method including polishing, mechanical dimpling and ion-beam milling.
43 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
3.3 Results and discussion
3.3.1 Crystal Structure
Table 3.1 lists the measured results for the PZT films grown on STO substrates. The films were well crystallized and oriented, as revealed by XRD analysis. All the PZT films were oriented with c-axis perpendicular to the substrate surface. Figure 3.1 shows the XRD patterns of the PZT film deposited for 20 minutes, where only high and sharp (001) and (002) peaks of PZT and peaks of the
STO substrate are observed. The peak marked by “S” is attributed to the STO substrate, which is probably from the miscut of the substrate. Careful step scans on the (001) peak of the PZT films revealed that there was a slight increase in the lattice constant c along c-axis by 0.6% as the film thickness increased from 18 nm to 549 nm
(Figure 3.2). We did not observe a decrease of c-axis constant because the thickness of the PZT films were larger than the critical thickness where most of the strain generated from lattice mismatch of STO had been relieved. The full width half maximum (FWHM) of the rocking curve of PZT (002) peak (Table 3.1 and Figure
3.2) varied from 0.46° to 1.48° without obvious trend. This is probably attributed to the uncertain surface condition of the STO substrate, e.g. different miscut degree and different defect density. This result implies that thick PZT film is not necessarily to be associated with poor crystalline orientation. The insensitivity of the film crystalline orientation to thickness observed here is definitely an advantage for MEMS application using thick PZT films. The underlying mechanism will be discussed later.
44 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
Table 3.1 Measurement results for the PZT films grown on STO substrates
Deposition Thickness Thickness FWHM of In-plane AFM root Time by Step by PZT (002) average mean square Profiler HRTEM rocking curve grain size roughness (min) (± 5 nm ) (± 7 nm) (± 0.05°) (± 5 nm) (± 0.1 nm) 1 19 17 0.64 1.3 3 58 52 1.48 25 1.0(1) 121 5.6(2) 6 116 104 0.59 107 5.8 10 193 173 0.69 142 7.7 20 385 347 1.22 166 8.0 30 578 520 0.47 153 9.2 40 770 693 1.30 177 11.9 (1)Underlying small grains layer (2)Top larger grains layer
PZT (200)
(100) STO S Substrate (002) (001)
S Intensity (Arb unit)
15 20 25 30 35 40 45 50 55 60 2θ
Figure 3.1 XRD θ−2θ scan of the PZT film deposited on SrTiO3 (100) substrate for 20 minutes.
45 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
0.4185
1.5 0.4180
0.4175
1.0 0.4170
0.4165
0.5 0.4160 C-Axis lattice constant (nm)
0.4155 FWHM of PZT (002) Rocking Curve (degree) Curve Rocking (002) of PZT FWHM 0 100 200 300 400 500 600 700 800 Thickness (nm)
Figure 3.2 C-axis lattice constant (nm) and full width at half maximum (FWHM) of PZT (002) rocking curve against different thickness.
3.3.2 Microstructure
The morphology analysis revealed some important clues of the growth mechanism of PZT on STO. Figure 3.3 shows typical AFM micrographs of the films with different thickness as well as the cleaned STO substrate. Figure 3.3a shows the terrace steps on STO surface with height of one lattice layer (~ 0.4 nm). Figure 3.3b is the morphology of the PZT film deposited for 1 minute (~18 nm in thickness). At this stage, the film had small nuclei/grains that grow uniformly across the substrate.
However, there are voids in the film, which means the substrate surface has not been fully covered yet. When deposition time is increased to 3 min (Figure 3.3c, film thickness ~55 nm), we can see that a layer of large grains appears on top of a layer of small grain layer. To the best of our knowledge, this is the phenomenon observed in
46 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
PZT films for the first time. The lower layer has an average grain size of 25-30 nm, which has fully covered the substrate surface. The upper layer, on the other hand, has an average grain size of 100-120 nm, three times bigger than that of the former.
However, at this stage, the upper layer covers only about 30-40 % area of the film surface. The surface AFM profile (Figure 3.4) shows that the surface of this layer with large grains is ~ 15 nm above the surface of the layer with small grains. In other words, the large grains begin to form at a film thickness of about 40 nm, nearly one hundred monolayers of (001) PZT.
As the deposition time is increased to 6 min (Figure 3.3d, thickness of ~110 nm), the film structure appears uniform allover the substrate surface, and most of the grains are separated by clearly defined grain boundaries. When the deposition time is further increased to 20 min or above (Figure 3.3f. 3.3g, 3.3h), some big grains with sizes of sub-micron or micron can be observed. The big grains could be formed due to either the coalescence and merge of the smaller grains or the presence of second phase particles, precipitates and droplets - a typical feature observed in the laser ablated thin films owing to the well-known “splashing” effect. This trend resulted in the formation of irregular-shaped grains and also caused an increase in the film roughness. But surprisingly, if these big particles are not counted, the average in- plane grain size of the “main body” of the film does not significantly increase with the increase of the film thickness (Table 3.1 and Figure 3.5).
47 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
(a) (b)
1 min
(c) (d)
3 min 6 min
(e) (f)
10 min 20 min
(g) (h)
30 min 40 min
Figure 3.3 AFM surface images (1μm x 1μm) of SrTiO3 substrate and the PZT films deposited on SrTiO3 substrate with different thickness. (a) SrTiO3 substrate surface, (b)-(h) PZT films on SrTiO3 substrate with different deposition times, from 1 to 3, 6, 10, 20, 30 and 40 min respectively.
48 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
Figure 3.4 AFM surface profile view of the PZT film with a thickness of about 50 nm. The size of the bottom image is 3μm x 3μm.
From the above results, it is suggested that the epitaxial (001) PZT films grown on lattice-matched (100) STO substrate had undergone a three-dimensional
(3D) island growth mode. In the early growth stage, nuclei were uniformly distributed on the substrate surface. At thickness of about 18 nm (45 monolayers), nuclei coalesced but the substrate surface was still not fully covered. The critical nucleus size was ~ 25-30 nm and the nucleus density was ~ 1.5 × 1011/cm2 (Figure
3.3b, 3.3c). At about 40 nm (100 monolayer), the substrate surface was completely covered. As some of the grains reached the critical size of nucleation, they would start
49 Chapter 3 Early growth stage of laser ablated epitaxial PZT film to grow in size and incorporate with surrounding grains of smaller size, thus leading to a two layers structure. With increasing thickness, the film was occupied by all large grains with irregular shapes and non-uniform size distribution. When the thickness is further increased, the shapes of the grains became more irregular and the size distribution became more non-uniform. Meanwhile, the roughness of the film increased. However, the average grain size kept to be ~ 100-150 nm, showing an insignificant increase with film thickness. The PZT films have a typical column-like structure that is often observed in epitaxial thin films with complex components deposited via PLD. For example, c-axis YBCO thin film prepared at 700-750°C usually showed a column structure with an in-plane grain size of 0.2-0.4 μm20. But the grain size of the films with simple compositions usually increase with film thickness.21
It is expected that, this column-structure growth is still dominating in thicker
PZT film, as long as the deposition parameters are not changed. Similar results have been reported by other researchers.22 As mentioned above, since LSMO, PZT and
STO have very small lattice mismatch, it is reasonable to assume that the results observed in the (001)PZT/(100)STO structure may also be true observed in
(001)PZT/(001)LSMO. This is further proven in chapter 6, where epitaxial PZT was deposited on LSMO film. To maintain this structure, substrate temperature and impinging energy of the species in the laser-generated plume are considered to be the two main factors9, 10, 14 that determine the diffusivity of the adatoms on the film/substrate surface. In addition, the structure and composition of grain boundaries in the PZT film may also play an important role, because more of less they restrict the merging effect of neighboring grains, thus limiting the final in-plane grain size.
50 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
180 14
160 12
140 10 120 8 100 6 80 4 60
Average in-plane grainsize (nm) 2 40 (nm) roughness Rootmean square
20 0 0 100 200 300 400 500 600 700 800 Film thickness (nm)
Figure 3.5 Root mean square roughness and average in-plane grain size versus film thickness.
Figure 3.6 shows the schematic diagram of the film growth structure to illustrating why the orientation of PZT films is insensitive to the film thickness.
Column growth maintains the crystalline orientation and quality of thick films and enables one to obtain samples with desired electrical properties. Column structure is also advantageous in reducing the surface roughness. AFM images showed that, the difference in height points on the top surface of a grain is less than 1/3 of the in-plane
(horizontal) grain size. This means that the root mean roughness of the PZT films with grain sizes of 100 – 150 nm is limited to less than 50 nm. For a 10-μm PZT film, the root mean roughness equals to only 0.5% of the film thickness, which is acceptable for practical applications. The two merits of column growth mode, namely
51 Chapter 3 Early growth stage of laser ablated epitaxial PZT film unchanged film orientation and suppressed surface roughness, are both favourite for
MEMS applications.
Figure 3.6 A schematic diagram of the PZT film growth structure.
Now, we turn to the microstructure of the PZT films and the interface of
PZT/STO by using HRTEM. Figure 3.7 shows a cross-sectional HRTEM image of the PZT film deposited for 3 min, having a uniform thickness of 52-55 nm. The difference between the thicknesses estimated from HRTEM image and that measured with a step profiler is about 10%. The consistent lower measured thickness obtained from HRTEM may due to the process of sample preparation for TEM experiment where the samples have to be polish and grind. The values of the film thickness used in the above discussion are the average of the two sets of measured results. At high
52 Chapter 3 Early growth stage of laser ablated epitaxial PZT film magnification, the PZT/STO interface is clearly demonstrated, as shown in Figure 3.8.
Since the lattice mismatch between PZT and STO is very small, PZT film is commensurately epitaxial on the STO substrate. The crystalline correlations in Figure
3.8 are (001)PZT//(100)STO and [110]PZT//[110]STO. The interface is clear and sharp, and lattice distortions occur only in a region of 1-2 nm at the interface, marked as “A”. A few nano-meters away from the interface, there are also some local misalignment in PZT lattice, marked as “B”. Defects in zone “A” are obviously mainly caused by the defects in the STO substrate, while the defects in zone “B” may be attributed to the impurity and stacking faults of the PZT film. About 20 nm away from the interface, the lattice structure of the PZT film appears almost perfect. In general, we think that the misalignment, due to the interface stress and substrate surface defects can be buffered by a stacking of PZT with about 50 monolayers. Even on the substrate surface with serious defect, e.g. a big step, PZT film can still stack perfectly after a few nano meters of misalignment, such as that shown in Figure 3.9.
Our HRTEM analysis shows that STO substrate is very favorable for the epitaxial growth of PZT films.
53 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
GLUE
[001] [100]
PZT STO
Figure 3.7 HRTEM cross-section view of the PZT film deposited on SrTiO3 substrate for 3 min.
PZT
[001]c B
[110] B
A A [100] STO
[011]
Figure 3.8 High magnification image of the interface structure between PZT film and the SrTiO substrate. 3
54 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
B PZT
B A
A
STO
Figure 3.9 HRTEM cross-section view of the PZT film deposited on SrTiO3 substrate showing a structural defect on the SrTiO3 substrate.
3.4 Conclusions
In this chapter, the early growth stage of epitaxial PZT thin film on STO substrate was studied in terms of surface morphology, crystalline orientation, microstructure and film-substrate interface. We obtained several valuable clues of the growth mechanism. (100) STO surface was found to be very favorable for epitaxial stacking of (001) PZT lattice. The epitaxial (001) PZT film on (100) STO appeared to undergo a three-dimensional island growth mode. The nuclei of PZT film, with a size around 25-30 nm and a density of the order of 1011/cm2, were found uniformly distributed on the STO substrate surface. However, the substrate surface was not fully filled after initial 40-50 monolayers were stacked. For the first time a two-layer structure in the early stage of the film growth, at a film thickness of about 40-50 nm or about 100 monolayers, was observed, where small nuclei/grains (25-30 nm in diameter) were merged into larger grains (100-120 nm in diameter). With further
55 Chapter 3 Early growth stage of laser ablated epitaxial PZT film increase in thickness, the films showed a dominant column-like growth mode. All samples with thickness of larger than 115 nm, have an average grain size ranging from 100-150 nm. However, with increasing thickness, the shapes of the grains of the films became more irregular and the roughness increased. We can expect that, this column-structure growth will still be dominating in even thicker PZT films, provided that the deposition parameters are maintained to be unchanged. The column growth of
PZT film is advantageous for MEMS applications. This is because the PZT films grown with a column structure have a crystalline orientation unchanged with thickness. Also, the column growth can effectively suppress the overall surface roughness.
Our HRTEM analysis revealed that the interface between PZT thin film and
STO substrate is very clear and sharp, where (001) PZT film was nearly commensurately epitaxially grown on the (100) STO substrate. Distortions of the lattices occurred only in a region of 1-2 nm along the interface, even in the case that the substrate surface had serious defects like steps. Regions with local lattice misalignment were also observed in PZT film several to nano meters away from the
PZT/STO interface. The PZT film showed perfect lattice stacking as the thickness is more than 20 nm and above, indicating that the misalignment due to the interface stress and substrate defects was buffered by a stacking layer of PZT with about 50 monolayers.
56 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
3.5 References
1 S. Mathews, R. Ramesh, T. Venkatesan and J. Benedetto, Science 276, 238 (1997).
2 J.F. Scott and C.A. Paz de Araujo, Science 246, 1400 (1989).
3 Ryoichi Takayama and Yoshihiro Tomita, J. Appl. Phys. 65, 1666 (1989).
4 H. Du, D. W. Johnson, Jr., W. Zhu, J. E. Graebner, G. W. Kammlott, S. Jin, J.
Rogers, R. Willett, and R. M. Fleming, J. Appl. Phys. 86, 2220 (1999).
5 D. L. Polla, C. Ye and T. Tamagawa, Appl. Phys. Lett. 86, 3539 (1991).
6 S.B. Majumder, D.C. Agrawal, Y.N. Mohapatra and V.N. Kulkarni, Intergrated
Ferroelectrics 9, 217 (1995).
7 M. Ichiki, J. Akedo, A. Schoroth, R. Maeda, Y. Ishikawa, Jpn. J. Appl. Phys. 36,
5815 (1997).
8 C. Lucat, F. Menil and R. VonderMuhll, Measurement Sci. & Tech. 8, 38 (1997).
9 G. K. Hubler, in Pulsed Laser Deposition of Thin Films, edited by D. B. Chrisey and
G. K. Hubler, (John Wiley & Sons Inc., New York, 1994), p. 327.
10 R. F. Haglund, Jr., Laser Ablation and Deposition, edited by J. C. Miller and R. F.
Haglund, Jr., (Academic Press, London, 1998), p.15.
11 J.F. Scott, C.A. Araujo, B.M. Melnick, L.D. Macmillan and R. Zuleeg, J. Appl.
Phys. 70, 382 (1991).
12 R. Ramesh, W.K. Chan, B. Wilkens, T. Sands, J.M. Tarascon, V.G. keramidas,
D.K. Fork, J. Lee and A. Safari, Appl. Phys. Lett. 61, 1537 (1992).
13 D.J. Lichtenwalner, R. Dat, O. Auciello and A.I. Kingon, Ferroelectrics 152, 97
(1994).
14 B. L. Low, S. Y. Xu, C. K. Ong, Z. X. Shen and X. B. Wang, Supercond. Sci. Tech.
10, 41 (1997).
15 S. Y. Xu, B. L. Low, C. K. Ong and X. Zhang, J. Mater. Sci. Lett. 16, 429 (1997).
57 Chapter 3 Early growth stage of laser ablated epitaxial PZT film
16 J.-M. Liu, and C. K. Ong, Appl. Phys. Lett. 73, 1047 (1998).
17 Q. Huang, J. M. Liu, J. Li, H. C. Fang, H. P. Li and C. K. Ong, Appl. Phys. A-Mater.
Sci. & Proc. 68, 533 (1999).
18 H. C. Fang, C. K. Ong, S. Y. Xu, K. L. Tan, S. L. Lim, Y. Li and J. M. Liu, J. Appl.
Phys. 86, 2191 (1999).
19 X. Y. Zhang, C. K. Ong, S. Y. Xu and H. C. Fang, Appl. Surf. Sci. 143, 323 (1999).
20 S. Y. Xu, C. K. Ong, and X. Zhang, Solid State Commun. 107, 273 (1998).
21 K. D. Develos, M. Kusunoki and S. Ohshima, Jpn. J. Appl. Phys. 37, 6161 (1998).
22 L.X. Cao, Y Xu, B.R. Zhao, L.P. Guo, L. Li, B. Xu, Y.Z. Zhang, H. Chen, A.J. Zhu,
Z.H. Mai, J.H. Zhao, Y.F. Fu and X.J. Li, J. Phys. D: Appl. Phys. 30, 1455 (1997).
23 K.S. So and K.H. Wong, Ferroelectrics 260, 557 (2001).
24 W.B. Wu, K.H. Wong, C.L. Choy and Y.H. Zhang, Appl. Phys. Lett. 77, 3441
(2000).
25 F. Mitsugi, Y. Yamagata, T. Ikegami, K. Ebihara, J. Narayan and A.M. Grishin, Jap.
J. Appl. Phys. Part I 39, 5418 (2000).
58 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
Chapter 4 Effects of microstructure on the properties of polycrystalline lead zirconate titanate (PZT) thin films
4.1 Introduction
In this chapter, we will study on polycrystalline PZT films deposited on Pt- coated silicon substrates. The property criteria to evaluate the merits of a PZT thin film are different when targeting different practical applications. It is well known that the properties of a PZT thin film strongly depend on its microstructure.1-3 Therefore, it would be essential to find out the relationship between the microstructure and its macroscopic properties, in order to achieve samples with a specific microstructure required by a specified application. In this chapter, two different deposition techniques, pulsed laser deposition (PLD) and sol-gel technique were used to fabricate
PZT thin films with distinct microstructures. The electrical and electromechanical properties of the PZT thin films were comparatively studied.
4.2 Experimental procedures
The details of the PLD setup have been discussed in Section 2.1.1. A PZT target with the composition near the morphotropic phase boundary (MPB) was prepared following a standard ceramic process. The target to substrate distance was about 3.5 sm. Pt-coated silicon (Pt/Ti/SiO2/Si) were used as the substrates. After cleaning, the substrate was immediately moved into the deposition chamber. The chamber was pumped down to 10-6 mbar and then the substrate was heated up to
350°C in 10 min. The ablation process was carried out for 30 min at a laser fluency of
1.8 J/cm2 and a repetition-rate of 5 Hz and an oxygen atmosphere of 0.2 mbar. The film was finally annealed at 700°C at 1atm oxygen pressure.
59 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
For sol-gel deposition, the precursors were prepared with a PZT composition same to the PLD target.4 The PZT thin film was spin-coated at 5000 rpm for 30 s followed by a pyrolysis at 430°C. The spin-coating process was repeated until a desired thickness is achieved. Finally, the film was annealed at 700°C to obtain the perovskite phase.
SEM and XRD characterizations of the PZT thin films are as same as those described in Chapter 3. For electrical measurement, a corner of the samples PZT film was then chemically etched to expose the bottom Pt electrode. Gold was sputtered to serve as the top electrode. Dielectric constant and loss tangent were measured using an Hewlett Packard impedance analyzer (HP 4194A). Ferroelectric hysteresis loop was measured using a standard ferroelectric test system (Radiant Technologies
RT3000). Piezoelectric vibration was measured using a laser vibrometer (OFV-056 and OFV-3001-SF6 from PolyTech, Berlin).5
60 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
4.3 Results and discussion
4.3.1 Crystal Structure
Figure 4.1 shows the XRD results of the PZT thin films grown by the PLD and sol-gel methods, indicating that both PZT films are polycrystalline without preferential crystallographic orientation. The XRD intensities of the two films are comparable. The similar orientation character together with the comparable diffraction intensity may preliminarily exclude any significant effects of crystallographic structure on the properties of the two thin films.
PZT (110) Pt (111)
(100) (200) (111) (200)
(a) Intensity (Arb Unit) (Arb Intensity
(b)
20 25 30 35 40 45 50 2θ
Figure 4.1 XRD patterns of the PZT thin films deposited by (a) PLD method and (b) sol-gel method.
4.3.2 Microstructure
Figures 4.2a and 4.2b show the SEM images of the surface morphology of the
PZT thin films grown by PLD and sol-gel methods, respectively. The PLD PZT thin
61 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film film exhibits a compact and flat surface with unclear grain boundaries. There are only a few big particulates scattered on the film. In contrast the sol-gel PZT film has a unsmoothed surface with a looser morphology. Many open boundaries can be observed among the particles. The particles comprise of multiple grains of tens of nanometers. The open boundaries were not spread cracks as the electrical measurement of the film showed good electrical properties.
Figures 4.2c and 4.2d show the cross-sectional SEM images of the two PZT films deposited by PLD and sol-gel process, respectively. The PLD film has a densely packed structure with no apparently visible grain boundaries for the film deposited by the PLD process, while the sol-gel PZT film shows a less compact morphology with clearly observable grains and particle boundaries. The thickness of both PZT films is around 500 to 650 nm.
The SEM results showed that we can prepare PZT films with distinct morphological characteristics, but similar crystallographic characteristics. Thus, these two PZT films can be used to investigate the effects of morphology and microstructure on their electrical and electromechanical properties.
62 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
(a)
(b)
Figure 4.2 Surface SEM images of the surface of PZT thin films deposited by (a) PLD method and (b) sol-gel method.
63 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
(c)
(d)
Figure 4.2 Cross-sectional SEM images of the cross-section of PZT thin films grown by (c) PLD and (d) sol-gel.
64 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
4.3.3 Electrical property
The hysteresis loops (polarization vs electric field) of the PZT films deposited using PLD and sol-gel process are presented in Fig. 4.3. Both films exhibit good ferroelectric properties, evidence by the well-developed P-E loops. The PLD PZT
2 film has a higher remnant polarization (Pr) of 40 μC/cm and a lower coercive field of
2 35 kV/cm, while the sol-gel film has a lower Pr of 23 μC/cm and a higher coercive field of 55 kV/cm. These results are comparable to those published in the open literatures.6-10
80 )
2 60
40 C/cm
μ 20
0
-20
-40
Polarization ( Polarization -60 PLD sol-gel -80
-150 -100 -50 0 50 100 150 Electric Field (kV/cm)
Figure 4.3 P-E hysteresis loops of the PZT thin films grown by PLD and sol-gel.
Figure 4.4 shows the dielectric constant and loss tangent of the PZT films prepared with the two different deposition methods. The dielectric constant of the
PLD grown PZT film is around 1800 at 1 kHz while the loss tangent is in the range of
0.04 ~ 0.05 (Fig. 4.4a). The dielectric constant of the sol-gel PZT film is about 750 at
65 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
1 kHz and the loss tangent is in the range of 0.03 ~ 0.04 (Fig. 4.4b). All electrical and piezoelectric properties of the two PZT films are summarized in Table 4.1.
Table 4.1 The properties of the PZT films grown by PLD and sol-gel. PZT Dielectric Properties Ferroelectric Properties Piezoelectric Deposition Properties Method Dielectric Loss Remnant Coercive Constant Tangent Polarization Field d33 g33 at 1kHz μC/cm2 kV/cm pm/V mmV/N pulsed laser 1853 ± 20 0.04 ~ 0.05 40 ± 2 35 ± 3 106 ± 2 6.7 ± 0.2
sol-gel 751 ± 20 0.03 ~ 0.04 23 ± 2 55 ± 3 102 ± 2 15.4 ± 0.2
The lower dielectric constant and remnant polarization for the sol-gel derived
PZT thin films are due to the less dense morphology as compared to the PLD film, which is revealed by the SEM observation. The effective electric field that was actually applied to the ferroelectric grains in the sol-gel film could be much lower than the calculated electric field due to the existence of the open grain and particle boundaries. This may explain why the sol-gel film has a higher coercive electric field.
No substantial difference in the loss tangent was observed for the two films.
66 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
2000 0.10 (a) 1800 0.09
1600 0.08 Dielectric Constant 1400 0.07
1200 0.06
1000 0.05
800 0.04
600 0.03 Loss Tangent Loss Tangent
Dielectric Constant 400 0.02
200 0.01
0 0.00 0.1 1 10 100 Frequency (kHz)
0.10 800 (b) 0.09
700 0.08 Dielectric Constant 600 0.07
500 0.06 0.05 400 0.04 300
0.03 Loss Tangent
Dielectric Constant 200 Loss Tangent 0.02
100 0.01
0 0.00 0.1 1 10 100 Frequency (kHz)
Figure 4.4 Dielectric constant and loss tangent of the PZT films grown by (a) PLD and (b) sol-gel.
67 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
4.3.4 Piezoelectric property
To measure the piezoelectric vibration along thickness direction, a sine wave alternating current (ac) with a driving voltage of 5 V at 1.5 kHz was applied to the
PZT films. Figures 4.5a and 4.5b present the displacement measured in the frequency domain through Fast Fourier Transformation (FFT) for the films prepared by the PLD and the sol-gel process, respectively. The effective piezoelectric constant d33 of the
PZT samples was estimated using the following equation:
δ d = (4.1) 33 V where δ is the displacement of the film and V is the applied voltage across the film.
The displacement of the PLD film is 531 pm and the effective d33 is 106 pm/V, while the displacement of the sol-gel film is 513 pm and the d33 is 102 pm/V. The actual d33 of the film could be significantly larger than the calculated value as the clamping effect of the substrate is not taken into account in the calculation with equation (4.1).
It is concluded that, the microstructural difference between the two PZT films does not result in a difference in their piezoelectric constant as significant as in their dielectric constant and ferroelectric properties.
The piezoelectric voltage constant g33 can be expressed as:
d33 g33 = T (4.2) 0εε r
T where ε0, εr are the vacuum dielectric constant and the relative dielectric constant under a constant stress. Thus the values of g33 are 6.7 mmV/N for PLD film and 14.4 mmV/N for the sol-gel film, calculated basing on the measured dielectric constant and the approximated piezoelectric constant d33.
68 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
From the microscopic, as shown in Figs. 4.2b and 4.2d, it is found that the open boundaries in the sol-gel film are largely along with the thickness direction.
Therefore, the film may have a relatively higher density in the thickness direction than in the plane. Such a structural feature may be helpful in maintaining a high piezoelectric constant d33 of the sol-gel film. In addition, the loose structure of the sol- gel film can also mitigate the in-plane clamping effect, thus leading to a large displacement in the thickness direction.
The higher remnant polarization and dielectric constant and the smaller coercive field, as a result of the compact and dense microstructure make the PLD films to be more advantageous for the applications of DRAM,11-13 FRAM,14-26 capacitor27-28. However, the sol-gel PZT films, with a relatively loose microstructure, could be more preferable for the application of piezoelectric sensor, due to their substantially higher piezoelectric voltage constant.29
It should be pointed out that the microstructure of PZT films can be greatly modified by changing the processing conditions and parameters of both PLD and sol- gel. Therefore, our intention was not to compare the two deposition techniques in this study. Our purpose was to establish a better understanding on the relationship between microstructural characteristics and ferroelectric and electromechanical properties and using the two films with distinct microstructures, derived from the two different processes. This will be helpful in selecting PZT films with specific properties for specific applications.
69 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
1.0 (a)
0.8
0.6
0.4
Displacement (nm) 0.2
0.0
1.40 1.45 1.50 1.55 1.60 Frequency (kHz)
1.0 (b)
0.8
0.6
0.4
0.2 Dispalcement (nm)
0.0
1.40 1.45 1.50 1.55 1.60 Frequency (kHz)
Figure 4.5 Displacement measurement results in the frequency domain for the PZT films grown by (a) PLD and (b) sol-gel.
70 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
4.4 Conclusions
Two types of the PZT films with distinct microstructure and greatly different properties were prepared with PLD and sol-gel process, respectively. It was observed that with the PLD-derived film a denser microstructure has significantly higher dielectric constant and remnant polarization and much lower coercive electric field than the sol-gel-derived film. However, the microstructural difference of the two PZT films only resulted in a relatively small difference in their piezoelectric constant d33, and the film with looser microstructure has a substantially higher piezoelectric voltage constant g33 due to the much lower dielectric constant. PZT film with higher density in the thickness direction than in the in-plane direction may help maintain the piezoelectric constant and eliminate the in-plane clamping effect. Our results indicate that a good understanding on the relationship between microstructure and property is essential in order to prepare PZT ferroelectric thin films with specific properties required by specific applications.
71 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
4.5 References
1. W.C. Goh, S.Y. Xu, S.J. Wang and C.K. Ong, J. Appl. Phys. 89, 4497 (2001).
2. J.K. Yang, W.S. Kim and H.R. Park, Appl. Surf. Sci., 169, 554 (2001).
3. M. Alguero, M.L. Calzada, E. Snoeck and L. Pardo, J. Eur. Ceram. Soc. 19, 1501
(1999).
4. S. Yu, K. Yao, S. Shannigrahi and E.H. Tay, J. Mater. Res. 18, 737 (2003).
5. K. Yao and F.E.H. Tay, IEEE Trans on Ultra. Ferro. Freq. Ctrl. 50, 113 (2003).
6. T. Kiguchi, N. Wakiya, K. Shinozaki and N. Mizutani, Micro. Eng. 66, 708 (2003).
7. H. Funakubo, M. Aratani, T. Oikawa, K. Tokita and K. Saito, J. Appl. Phys. 92,
6768 (2003).
8. J.M. Liu, S.Y. Xu, W.Z. Zhou, X.H. Jiang, C.K. Ong and L.C. Lim, Mater. Sci. &
Eng. A 269, 67 (1999).
9. M. Tsukada, H. Yamawaki and M. Kondo, Appl. Phys. Lett. 83, 4393 (2003).
10. W.M. Gilmore, S. Chattopadhyay, A. Kvit, A.K. Sharma, C.B. Lee, W.J. Collis, J.
Sankar and J. Narayan, J. Mater. Res. 18, 111 (2003).
11. H. Yang, B. Chen, K. Tao, X. Qiu, B. Xu and B. Zhao, Appl. Phys. Lett. 83, 1611
(2003).
12. S.T. Chang and J.Y. Lee, Appl. Phys. Lett. 80, 655 (2002).
13. N. Maffei and S.B. Krupanidhi, Appl. Phys. Lett. 60, 781 (1992).
14. J. Junquera and P. Ghosez, Nature 422, 506 (2003).
15. H. Cheng, Y. Chen, C. Chou, K. Chang, C. Hou and I. Lin, J. Appl. Phys. 87, 8695
(2000).
16. J. Scott and C. Paz de Arujo, Science 246, 1400 (1989).
17. K. Saito, T. Kurosawa, T. Akai, T. Oikawa and H. Funakubo, J. Appl. Phys. 93,
545 (2003).
72 Chapter 4 Effect of microstrucuture on the properties of polycrystalline PZT film
18. S.R. Gilbert, S. Hunter, D. Ritchey, C. Chi, D.V. Taylor, J. Amano, S. Aggarwal,
T.S. Moise, T. Sakoda, S.R. Summerfelt, K.K. Singh, C. Kazemi, D. Carl and B.
Bierman, J. Appl. Phys. 93, 1713 (2003).
19. M. Tsukada, H. Yamawaki and M. Kondo, Appl. Phys. Lett. 83, 4393 (2003).
20. J. F. Scott, Ferroelectric Memories (Springer, Berlin, 2000).
21. Z. G. Zhang, D. P. Chu, B. M. McGregor, P. Migliorato, K. Ohashi, K. Hasegawa,
and T. Shimoda, Appl. Phys. Lett. 83, 2892 (2003).
22. F. Yan, Y. Wang, H.L.W. Chen, C.L. Choy, Appl Phys. Lett. 82, 4325 (2003).
23. K. Tokita, M. Aratani and H. Funakubo, Appl. Phys. Lett. 82, 4122 (2003).
24. Auciello, J.F. Scott and R. Ramesh, Phys. Today 51, 22 (1998).
25. F. Mcnally, J.H. Kim and F.F. Lange, J. Mater. Res. 15, 1546 (2000).
26. Y. Zhu, M. Shen, R. Xu and W. Cao, Surf. Coat. Tech. 160, 277 (2002).
27. M.H. Frey and D.A. Payne, Appl. Phys. Lett. 63, 2753 (1993).
28. A. Lin, X. Hong, V. Wood, A.A. Verevkin, C.H. Ahn, R.A. McKee, F.J. Walker
and E.D. Specht, Appl. Phys. Lett. 78, 2034 (2001).
29. E. Dimitriu, R. Ramer, C. Midea and C. Tanasoiu, Ferroelectrics 241, 1851 (2000).
73 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT
Chapter 5 Epitaxial La0.7Sr0.3MnO3 conductive film as bottom electrode for PZT film
5.1 Introduction
In order to fabricate high quality epitaxial ferroelectric film on silicon substrate, perovskite oxide has to be used as bottom electrode. In this study, the perovskite oxide that we selected is La0.7Sr0.3MnO3 (LSMO). The high electrical conductivity and good lattice matching with perovskite ferroelectric oxides, such as
PZT and barium strontium titanate (BST), make LSMO an attractive candidate as bottom electrode material for ferroelectric thin films integrated with silicon substrate.
Although epitaxial LSMO thin films grown on lattice matched single crystal
1 oxide substrates such as SrTiO3 (STO), and LaAlO3 (LAO), have been widely reported, available information on LSMO thin films grown on silicon substrates is very limited in the literatures.2-4 Due to the fact that there is a big lattice mismatching between silicon and LSMO, and chemical reactions between LSMO and silicon easily take place at high temperatures to form interfacial phases, it is necessary to use appropriate buffer layers to achieve high quality epitaxial LSMO film growth on silicon substrates. It is also of interest to investigate the epitaxial mechanism of
LSMO film on buffered silicon substrates.
In this chapter, the deposition of LSMO films on yttria-stabilized zirconia
(YSZ) buffer layer on silicon will be described. YSZ layer not only has a good lattice matching with LSMO but also prevents the formation of an interfacial layer between
LSMO and silicon. Introducing an additional buffer layer of YBa2Cu3O7-δ (YBCO) on top of the YSZ layer can promote the epitaxial growth of (001)-oriented LSMO thin films on silicon. The epitaxial growth mechanism, crystal structures, surface
74 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT morphology, resistivity, magnetoresistance, and their interrelationships will also be discussed in this chapter.
5.2 Experimental Procedure
(100)-oriented silicon substrates were cleaned for 5 min using nitric acid in an ultrasonic cleaner to remove any natural oxide layer or oxide contaminant on surface.
The substrates were subsequently cleaned with de-ionized water, acetone and ethanol.
All the YSZ, YBCO, and LSMO thin films were deposited by PLD. Detailed processing parameters the three films are summarized in Table 5.1. Two groups of multilayer samples, namely LSMO/YSZ/Si and LSMO/YBCO/YSZ/Si, were prepared. In addition, an LSMO film on lattice-matched single crystal LAO substrate was fabricated for comparison.
Both θ-2θ and φ-scan XRD studies were conducted to determine the crystalline phase and the crystal orientation of the films. SEM examination was carried out to observe the morphology of the films. A four-probe resistance measurement system (Keithley 220 current source, Keithley 182 voltmeter and
Lakeshore 330 temperature controller) was used to analyze the electrical properties of the samples from room temperature to 77 K. Magnetoresistance (MR) experiments with magnetic field up to 3000 Gauss (home made electromagnet) were also carried out at 77 K to evaluate the magnetic properties of LSMO films.
75 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT
Table 5.1 Summary of PLD processing parameters for the YSZ, YBCO, and LSMO thin films.
Layer Deposition Oxygen Laser Repetition Deposition Thickness Temperature Ambience fluence Rate Time (0C) (mbar) (J/cm2) (Hz) (min) (x10-9 m) YSZ 670 5x10-4 5 5 13 10 YBCO 650 5x10-1 3.5 3 8 200 LSMO 670 2x10-1 5 5 45 1000
5.3 Results and discussion
5.3.1 Crystalline structure
XRD patterns of the LSMO/YSZ/Si and LSMO/YBCO/YSZ/Si samples are shown in Figure 5.1. Slight (001)-orientation preference and several unidentified peaks are observed in the LSMO film grown on YSZ/Si, whereas the LSMO film grown on YBCO/YSZ/Si exhibits a perfect (001) orientation and no unidentified peak is present.
XRD φ scans for the LSMO/YBCO/YSZ/Si and LSMO/YSZ/Si samples are illustrated in Figures 5.2a and 5.2b, respectively. The YSZ layers exhibit a cube-on- cube epitaxial relationship to the (100)-oriented single crystal silicon substrates for both samples, as revealed by the matched 4-fold symmetry for both YSZ and silicon shown in Figures 5.2a and 5.2b. The LSMO film on YSZ, which shows a slight (001) orientation, exhibits weak 4-fold in-plane symmetry with broad diffraction peaks of plane {113}, as shown in Figure 5.2b. From these results it can be concluded that the
LSMO film on YSZ consists mostly of randomly oriented grains with a small number of epitaxial grains.
76 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT
LSMO Unidentified Phase YBCO * YSZ (110) (002)
(001) (002) * * * * LSMO/YSZ/Si Intensity (Abr Unit) Intensity
(005) LSMO/YBCO/YSZ/SI
20 25 30 35 40 45 50 55 60 2θ
Figure 5.1 XRD patterns (θ-2θ) for the LSMO thin films deposited on YBCO/YSZ/Si and YSZ/Si substrates.
The φ scan pattern of the LSMO{103} plane of LSMO/YBCO/YSZ/Si sample exhibits strong 8-fold diffraction peaks, as shown in Figure 5.2a. This confirms that there are two sets of in-plane orientations in the LSMO layer with an in-plane shift by
45° from each other. The two sets of in-plane orientation are found to result from the corresponding two sets of the in-plane orientation of YBCO and LSMO, as evidenced by the φ scan results of the {103} planes for both LSMO and YBCO layers in the inset in Figure 5.2a. As Figure 5.1 shows a perfect (001) orientation, the LSMO film is epitaxial on YBCO/YSZ/Si but has two sets of in-plane orientation, with a shift by
45° from each other.
77 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT
(a)
LSMO {103}
Intensity (Arb unit) YBCO {103} 0 50 100 150 200 250 300 350 Phi Angle (III)
(II) Intensity (Arb Unit)
(I)
0 50 100 150 200 250 300 350 Phi Angle
(b) (III)
(II)
Intensity (Arb Unit) Intensity (I)
0 50 100 150 200 250 300 350 Phi Angle
Figure 5.2 (a) XRD patterns (φ scan) for (I) Si {113}, (II) YSZ {113} and (III) LSMO {113} peaks for the LSMO/YBCO/YSZ/Si multilayer. The inset figure shows the relative φ angle relationship between YBCO {103} and LSMO {103}. (b) X-ray diffraction results (φ scan) for (I) Si {113}, (II) YSZ {113} and (III) LSMO {113} peaks for the LSMO/YSZ/Si multilayer.
78 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT
5.3.2 Microstructure
Cross-sectional SEM images of LSMO/YBCO/YSZ/Si and LSMO/YSZ/Si multilayer structures are shown in Figures 5.3a and 5.3b, respectively. The LSMO film on YBCO exhibits a columnar structure, while the LSMO film on YSZ/Si has a randomly stacking structure. The columnar grains in the LSMO film on
YBCO/YSZ/Si are 100-175 nm in width across the film. The LSMO film on YSZ/Si has smaller grain size (more grain boundaries) than the LSMO film on YBCO/YSZ/Si.
This is also supported by the surface SEM images of the films (not shown). The SEM images indicate that the LSMO film on YBCO/YSZ/Si has a better epitaxial quality than that on YSZ/Si, which is in a good agreement with XRD results. Both LSMO films show a root-mean-square surface roughness of about 10 nm AFM.
79 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT
(a)
LSMO
YBCO YSZ
Si
(b)
LSMO
YSZ
Si
Figure 5.3 Cross-sectional SEM images of cross-sections through (a) LSMO/YBCO/YSZ/Si and (b) LSMO/YSZ/Si.
80 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT
5.3.3 Resistivity and magnetoresistance
Figure 5.4 shows the temperature-resistivity curve of the LSMO films deposited on LAO, YBCO/YSZ/Si and YSZ/Si. The resistivity of the LSMO film on
YBCO/YSZ/Si is much lower than that of the LSMO film on YSZ/Si, but is still higher than the resistivity of the epitaxial LSMO film grown on lattice-matched LAO.
3.2 (c) 3.0 2.8
cm) 2.6 Ω 2.4 2.2
1.2 1.0 (b) 0.8 0.6 Resistivity (m 0.4 (a) 0.2 0.0 100 150 200 250 300 Temperature (K)
Figure 5.4 Temperature dependence of resisitivity for the LSMO thin films deposited on (a) LAO, (b) YBCO/YSZ/Si, and (c) YSZ/Si substrates.
The magnetoresistance (MR) of the LSMO films deposited on YBCO/YSZ/Si and YSZ/Si are given in Figure 5.5. A magnetic field of up to 3000 Gauss was applied perpendicular to the surface of the samples at 77 K. It is found that the LSMO film on
YSZ/Si has much larger MR (~16%) than the LSMO film on YBCO/YSZ/Si (~2%).
Similar results were reported by J.Y. Gu et al,5 who found that LSMO films with a higher density of grain boundaries showed a larger MR.
81 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT
2 0 (a) -2 -4 -6
-8
MR(%) -10 -12
-14 (b) -16 -18 -3000 -2000 -1000 0 1000 2000 3000 Magnetic Field (Gauss)
Figure 5.5 Magnetoresistance (MR) vs. magnetic field for the LSMO thin films deposited on (a) YBCO/YSZ/Si and (b) YSZ/Si substrates.
5.3.4 Discussion
The growth of YSZ on (100)-silicon is governed by a [100]//[100] cube-on- cube mechanism, as reported previously by our group.6 When deposited onto YSZ film, YBCO film exhibited a strong c-axis orientation along the thickness direction and two in-plane orientations with a 45° shift from each other, i.e.,
YBCO[100](001)//YSZ[100](001) and YBCO[110](001)//YSZ[100](001), as shown in Figure 5.2a. A similar dual in-plane orientation for YBCO thin films has been reported in the literature.7-9 The orientation of LSMO[100](001) is more preferred when grown on YBCO/YSZ/Si than YSZ/Si, as evidenced by the XRD patterns shown in Figure 5.2. Without introducing the YBCO layer, the LSMO film did not exhibit the same in-plane orientation as YSZ when deposited on YSZ/Si. The
82 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT epitaxial growth of YBCO[110](001)//YSZ[100](001) is obviously due to the small mismatching of 6.24% between (110) lattice of YBCO and (100) lattice of YSZ.
Schematic diagrams to illustrate both the in-plane orientations of YBCO on YSZ are shown in Figure 5.6. The Cu-O plane layers are proposed to be in contact with YSZ.
In Figure 5.6a, the Cu-O plane of YBCO is grown onto the O-O plane of YSZ (at 3/4 of a YSZ unit cell) to form the YBCO[100]//YSZ[100] orientation. In Figure 5.6b, the
Cu-O plane of the YBCO is turned by 45º in-plane on the Zr-Zr plane of YSZ to realize the YBCO[110]//YSZ[100] epitaxial growth.
(a) (b)
(001)
(010) (001) (010) (100) Oxygen (100) Zirconium
Copper
Figure 5.6 A schematic diagram of the proposed epitaxial growth mechanisms of the YBCO thin film for the two observed in-plane orientations on the YSZ buffer layer: (a) YBCO [100](001)//YSZ [100](001), in which the Cu-O plane of the YBCO is grown on the O-O plane of the YSZ (3/4 height of the unit cell of YSZ); (b) YBCO [110](001)//YSZ [100](001), the Cu-O plane of the YBCO is turned 45º in-plane on the Zr-Zr plane of YSZ.
83 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT
It is necessary to analyze the structures of YBCO and LSMO to understand the reason why YBCO can be epitaxially grown on YSZ/Si by forming a
YBCO[100]//YSZ[100] in-plane orientation while LSMO, which has similar lattice parameters, does not grow on YSZ/Si with the same in-plane orientation. The mismatching for YBCO[100]//YSZ[100] and LSMO[100]//YSZ[100] are 24.90% and
24.71%, respectively. In orthorhombic YBCO crystal structure, as shown in Figure
5.7a, the Cu-O bonds at the second and third planes are slightly bent, whereas the Mn-
O plane in the perovskite LSMO structure is flat, as shown in Figure 5.7b. The bent
Cu-O plane in YBCO is more suited for the epitaxial growth of LSMO as depicted in
Figure 5.6a, because it tolerates the lattice mismatching and provides more angular flexibility for forming chemical bonds between copper and oxygen at the interface. In addition, the electron configuration 3d104s1 of copper could also form Cu+ that electrically neutralizes oxygen vacancies at the interface between YSZ and YBCO.
All these characteristics can not be found in LSMO/YSZ. It is therefore essential to have an additional YBCO template layer in order to grow completely c-axis oriented
LSMO films on YSZ buffered silicon substrate.
84 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT
(001) (001)
(010) (010) (100) (100) Yttrium Lanthanum or Strontium Barium Oxygen Oxygen Manganese Copper
Figure 5.7 Structure of (a) orthorhombic YBCO crystal and (b) perovskite LSMO crystal.
The difference in resistivity between LSMO/YBCO/YSZ/Si and
LSMO/YSZ/Si films can be understood in terms of electron polarization and scattering mechanisms. Inside grains, the conduction electrons are spin polarized, i.e. having the same magnetic dipole direction, and are easily transferred between pairs of
Mn3+ and Mn4+ ions. However, when these electrons travel across grain boundaries, they experience strong spin-dependent scattering. A large density of grain boundaries would lead to an increase in scattering and hence result in an increase in resistivity.
The lower resistivity of the LSMO film grown on YBCO/YSZ/Si as compared to the
LSMO film on YSZ/Si is due to the fact that the former has a complete c-axis orientation with a columnar grain structure. It is also understandable that the LSMO
85 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT film on YBCO/YSZ/Si has higher resistivity compared to the epitaxial LSMO film on
LAO substrate.
With the application of a magnetic field, the randomly oriented magnetic domains in a polycrystalline sample are re-oriented, which leads to a significant decrease in the resistivity. Therefore, LSMO film on YSZ/Si has the largest MR as compared to the films on the YBCO/YSZ/Si and LAO.5,10
5.4 Conclusions
(001)-oriented epitaxial LSMO thin films with two in-plane orientations were fabricated on silicon substrates with YSZ and YBCO as buffer layers, using PLD. The epitaxial LSMO films on the YBCO layers had a cross-sectional columnar structure.
The dual in-plane orientations of the LSMO film were formed by epitaxial growth from the dual orientations in the YBCO template layer. The crystal orientation relationship among the LSMO, YBCO, and YSZ films can be expressed as,
LSMO/YBCO[110](001)//YSZ[100](001) and
LSMO/YBCO[100](001)//YSZ[100](001). In contrast, when LSMO films were deposited directly on YSZ/Si without using the YBCO template layer, they demonstrated a mixture of randomly oriented polycrystalline grains and (001)- oriented grains. Therefore, introduction of a YBCO layer is essential to obtain c-axis- oriented epitaxial LSMO films due to the dual-in-plane orientation growth mechanism of YBCO on YSZ/Si. The difference between the in-plane orientations of the LSMO and YBCO films on YSZ/Si is mainly attributed to the disparity in their crystallographical structure. The LSMO film with the dual-in-plane orientations on
YBCO/YSZ/Si exhibited a much lower resistivity as compared to the LSMO film on
YSZ/Si, because the later had more randomly oriented grains and hence more electron
86 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT scattering at grain boundaries. The low resistivity of the LSMO film on
YBCO/YSZ/Si means that it is a good candidate as bottom electrode for the subsequent growth of PZT film. However, due to high density of grain boundaries, the
LSMO film on YSZ/I showed higher MR effect.
87 Chapter 5 Epitaxial LSMO conductive film as bottom electrode for PZT
5.5 References
1 J. Li, C.K. Ong, Q. Zhan and D.X. Li, J. Phys.: Condens. Matter 14, 6341 (2002).
2 Z. Tranjanovic, C. Kwon, M.C. Robson, K.-C. Kim, M. Rajeswari, R. Ramesh, T.
Venkatesan, S.E. Lorland, S.M. Bhagat and D. Fork, Appl. Phys. Lett. 69, 1005
(1996).
3 J.Y. Gu, C. Kwon, M.C. Robson, Z. Tranjanovic, K. Ghosh, R.P. Sharma, R.
Shreekala, M. Rajeswari, T. Venkatesan, R. Ramesh and T.W. Noh, Appl. Phys. Lett.
70, 1763 (1997).
4 J.-H Kim, S.I. Khartsev and A.M. Grishin, Appl. Phys. Lett. 82, 4295 (2003).
5 J.Y. Gu, S.B. Ogale, M. Rajeswari, T. Venkatesan, R. Ramesh, V. Radmilovic, U.
Dahmen, G. Thomas and T.W. Noh, Appl. Phys. Lett. 72, 1113 (1998).
6 S.J. Wang, S.Y. Xu, L.P. You, S.L. Lim and C.K. Ong, Supercond. Sci. Tech 13,
362 (2000).
7 S.M. Garrison, N. Newman, B.F. Cole, K. Char, and R.W. Barton, Appl. Phys. Lett.
58, 2168 (1991).
8 D.M. Hwang, T.S. Ravl, R. Ramesh, S.W. Chan, C.Y. Chan, L. Nazor, X.D.Wu, A.
Inam and T. Venkatesen, Appl. Phys. Lett. 54, 1690 (1990).
9 L.A. Tietz, C.B. Carter, D.K. Lathrop, S.E. Russek, R.A. Buhrman and J.R. Michael,
J. Mater. Res. 4, 1072 (1989).
10 M. F. Hundley, M. Hawley, R. H. Heffner, Q. X. Jia, J. J. Neumeier, J. Tesmer, J. D.
Thompson, and X. D. Wu, Appl. Phys. Lett. 67, 860 (1995).
88 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate
Chapter 6 Pseudo-epitaxial lead zirconate titanate (PZT) thin film on silicon substrate
6.1 Introduction
Epitaxial perovskite ferroelectric Pb(ZrxTi1-x)O3 (PZT) thin films deposited on lattice matching substrates, such as LAO or STO, have exhibited superior ferroelectric properties in comparison with polycrystalline PZT films.1-3 In order to exploit the superior ferroelectric properties for silicon-based device applications, intensive efforts have been made to grow epitaxial PZT thin films on silicon substrate. It has been reported that epitaxial PZT thin films can be developed on multilayer buffered silicon
4 2 substrates. They include YBa2Cu3O7/YSZ/Si, LaSr0.5Ca0.5O3/SrTiO3/Si,
5 1 Ir/MgAl2O4/SiO2/Si, LaSr0.5Ca0.5O3/CeO2/YSZ/Si, and
6 SrRuO3/SrTiO3/MgO/TiN/Si. Although PZT films exhibiting good epitaxial quality have been achieved on silicon substrate, their ferroelectric properties are significantly inferior to those grown on lattice matching perovskite oxide layers. In this chapter, a pseudo-epitaxial PZT thin film, due to broad full width half maximum peak in the
XRD φ scan and granular film morphology, was prepared on silicon substrate with a
LSMO/ YBCO/YSZ7 buffer structure by a PLD. Despite the imperfect epitaxial quality, the pseudo-epitaxial PZT thin film exhibited a substantially larger ferroelectric polarization than those “ideal” epitaxial films on silicon substrates.
6.2 Experimental procedures
The deposition and characterization of LSMO/YBCO/YSZ/Si and PZT films can be referred to chapter 5 and chapter 4, respectively.
89 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate
6.3 Results and discussion
6.3.1 Crystal structure
XRD of PZT/LSMO/YBCO/YSZ/Si is shown in Figure 6.1. The PZT films has a strong (001)-orientation preference of PZT is observed, without any pyrochlore phase. XRD rocking curve full width at half maximum (FWHM) of PZT (001) is
1.71°. XRD φ scan results for the PZT/LSMO/YBCO/YSZ/Si are presented in Figure
6.2. The LSMO layer exhibits a cube-on-cube epitaxial relationship to silicon substrate as discussed in Chapter 5. The PZT film on the LSMO layer exhibits 4-fold symmetry with broad diffraction peaks of plane {113}, as shown in Figure 6.2. This confirms that there is only one set of in-plane orientation in the PZT layer. The in- plane and out-of-plane orientations of our PZT film are similar to an epitaxial film with a cube-on-cube crystal orientation with respect to the silicon substrate. However, the XRD φ scan peaks of the PZT layer were very broad, with a FWHM of 12.18°, indicating that the epitaxial quality is far from perfect.
90 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate
PZT LSMO YBCO
(001) YSZ (002)
(002) (001) Intensity (Arb Unit) Intensity (006) (002) (004)
20 25 30 35 40 45 50 55 60 2θ
Figure 6.1 XRD patterns (θ-2θ) for the PZT thin films deposited on LSMO/YBCO/YSZ/Si.
PZT
LSMO
Intensity (Arb unit) Intensity
Si
0 50 100 150 200 250 300 350 Phi Angle
Figure 6.2 XRD patterns (φ scan) for (I) Si {113}, (II) LSMO {113} and (III) PZT {113} peaks for the PZT/LSMO/YBCO/YSZ/Si multilayer.
91 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate
6.3.2 Microstructure
Surface SEM image of the PZT film and the cross-sectional image of the
PZT/LSMO/YBCO/YSZ/Si multilayer structure are presented in Figure 6.3a and
Figure 6.3b, respectively. The thickness of the PZT film is 714 nm. Both the surface and cross-section images of the PZT film show a granular morphology with pores of about 20 nm uniformly distributed across the film. We attribute the existence of these nanometer-sized pores to the volatilization of the excess 10 wt% PbO during the final annealing process. The extra PbO is unlikely to evaporate during the deposition process at the low deposition temperature of 350°C, and is incorporated inside the
PZT film. However, these excess PbO, which are most likely located at the grain boundaries or interstitial sites, will evaporate during the final annealing process at
700°C for 1 hour and thus leaving these nanometer-sized pores in the film. The pores produced by evaporation of PbO have also been observed in polycrystalline PZT films under different processing conditions.13-14 Furthermore when we looked at the SEM surface image of the PZT film before the annealing process, it did not show similar nanometer-sized pores which further supports that the nanometer-sized pores could be due to the evaporation of the excess PbO. Nevertheless, the excess PbO might not be the only reason to the formation of the pores. The granular morphology is a result of the low deposition temperature followed by the high temperature annealing. When
PZT first nucleates on surface of the LSMO layer during low temperature deposition, it would not have much mobility energy to relocate itself to the lowest energy position.
When subsequently annealed at a much higher temperature of 700°C, the PZT film started to undergo an internal nucleation process, which is a solid-solid transportation process.15
92 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate
(a)
(b)
Figure 6.3 SEM images of (a) surface and (b) cross-section of PZT/LSMO/YBCO/YSZ/Si heteromultilayer structure.
93 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate
Under such internal nucleation process, the mobility of the adatoms is very limited, thus leading to the formation of granular morphology of the PZT film. The microscopic granular morphology of the PZT film is consistent with the relatively large FWHM in both θ-2θ and φ scan XRD patterns. Therefore, we call the PZT thin film in the present study a pseudo-epitaxial film.
6.3.3 Electrical properties
The P-E hysteresis loop for the PZT thin film on LSMO/YBCO/YSZ/Si is
2 presented in Figure 6.4, showing a remnant polarization of (Pr) about 53 μC/cm and a coercive field of about 60 kV/cm. The room temperature dielectric constant of the epitaxial PZT film is about 1678 while the loss tangent is in the range of 0.04 ~ 0.05 at 1 kHz (Figure 6.5). The remnant polarization is significantly higher than whose
(10~30 μC/cm2) of the epitaxial PZT films on silicon substratereported in the literatures.1,3-6
94 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate
100 80 ) 2 60 40
C/cm 20 μ 0
-20 -40 -60 -80 Polarization ( -100 -120 -400 -300 -200 -100 0 100 200 300 400 Electric Field (kV/cm)
Figure 6.4 P-E hysteresis loop of a PZT thin film sample deposited on a LSMO/YBCO/YSZ/Si substrate.
1800 0.5
1700 0.4
1600 0.3
1500
0.2 1400 Loss Tangent Dielectric Constant 0.1 1300
1200 0.0 0.1 1 10 100 Frequency (kHz)
Figure 6.5 Dielectric constant and loss tangent of a PZT film deposited on LSMO/YBCO/YSZ/Si substrate.
95 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate
Why is the remnant polarization of the pseudo-epitaxial PZT thin film with nanometer-sized pores and granular morphology larger than whose of the high quality epitaxial PZT thin films on a silicon substrate? It has been well-recognized that the remnant polarization of a ferroelectric thin film is closely related to stress, and tensile stress leads to a lower remnant polarization.16 According to the θ-2θ scan XRD results, the thin film PZT (001) XRD peak slightly shifts towards a higher angle as compared to that of the target, as shown in Figure 6.6. This observation indicates that an in- plane tensile stress exists in the PZT thin film. Polycrystalline PZT thin films deposited on silicon substrate usually also show a tensile residual stress.17 Due to the thermal expansion coefficient (TEC) mismatch between Si and PZT, which are 3.5 ×
10-6/°C and 6 – 7 × 10-6/°C respectively, a large tensile residual stress could be generated during the cooling process after annealing. If only taking into account the thermal mismatch effect, the residual stress σ in the PZT thin film can be estimated using the following equation:
σ Y ()α α 211111 Δ−= T (6.1) where α1, α2 are the TECs of PZT and Si, respectively, Y11 is the Young’s modulus of the PZT thin film, and ΔT is the difference between the annealing temperature and room temperature. The estimated strain and tensile stress are 0.2% and 264 MPa, respectively. Such a large tensile stress could exist in a defect-free epitaxial PZT thin film deposited on Si, which would substantially lower the remnant polarization.5,18
However, due to the granular morphology with relatively open grain boundaries and nanometer-sized pores, the residual stress in our pseudo-epitaxial PZT thin film could be relieved to certain a extent although not completely. This may be the major reason why our pseudo-epitaxial PZT thin film exhibits a higher remnant ferroelectric polarization than the “ideal” epitaxial film on silicon. It is also understandable that the
96 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate remnant polarization of the pseudo-epitaxial PZT thin film is certainly not as good as the epitaxial PZT thin films grown on lattice matching perovskite substrate, in which tensile stress is not substantial.
Film PZT (001)
Target PZT (001)
Intensity (Arb Unit)
21.0 21.5 22.0 22.5 23.0 2θ
Figure 6.6 XRD pattern (θ-2θ) of PZT (001) peaks for the target and film.
6.3.4 Piezoelectric properties
To measure the piezoelectric vibration along thickness direction, a sine wave alternation current (ac) with a peak-to-peak voltage of 9 V at 5 kHz was applied to the
PZT film. Figure 6.7 presents the instantaneous vibration data when the dilatation of the PZT films reaches the maximum magnitude. The effective piezoelectric constant d33 of the PZT samples can be approximated for comparison purpose using equation
(4.1). The average dilatation of the film is calculated by averaging the displacement at all the points on top and subtracting the average displacements of the points at bottom.
97 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate
The result showed that the PZT film had an effective piezoelectric constant of 80.70 pm/V.
700 )
600 m
p 500
(
t 400 n
e 300 m
e 200
c
a 100 l
p 0 s
i -100 D -200 0.9 0.1 0.8 0.2 0.7 0.3 0.6 0.4 0.5 ) X 0.5 0.4 m A 0.6 x 0.7 0.3 (m is 0.2 (m 0.8 is m 0.9 0.1 Ax ) Y
Figure 6.7 Displacement measurement results in the time domain for the PZT films on LSMO/YBCO/YSZ/Si driven by an AC peak-to-peak 9 V at 1500 Hz.
6.4 Conclusions
In summary, a pseudo-epitaxial PZT thin film was fabricated on silicon substrate buffered with a LSMO/YBCO/YSZ heterostructure template by PLD. The pseudo-epitaxial PZT thin film was characterized with broad XRD peaks and granular morphology with nanometer-sized pores distributed across the film. Despite the imperfect epitaxial quality, the pseudo-epitaxial PZT thin film exhibited a substantially larger ferroelectric polarization than an “ideal” epitaxial PZT thin film deposited on Si. The improvement in ferroelectric polarization of our pseudo-epitaxial
PZT films has been attributed to the partial relief of the tensile stress by virtue of the granular pseudo-epitaxial structure with nanometer-sized pores. The results indicate
98 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate that only improving the epitaxial quality without considering the tensile stress effects may not be sufficient in achieving optimal ferroelectric polarization of a ferroelectric film on Si.
99 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate
6.5 References
1 T. Kiguchi, N. Wakiya, K. Shinozaki and N. Mizutani, Micro. Eng. 66, 708 (2003).
2 Y. Wang, C. Ganpule, B.T. Liu, H. Li, K. Mori, B. Hill, M. Wuttig, R. Ramesh, J.
Finder, Z. Yu, R. Droopad and K. Eeisenbeiser, Appl. Phys. Lett. 80, 97 (2002).
3 H. Funakubo, M. Aratani, T. Oikawa, K. Tokita and K. Saito, J. Appl. Phys. 92,
6768 (2003).
4 J.M. Liu, S.Y. Xu, W.Z. Zhou, X.H. Jiang, C.K. Ong and L.C. Lim, Mater. Sci. &
Eng. A 269, 67 (1999).
5 M. Tsukada, H. Yamawaki and M. Kondo, Appl. Phys. Lett. 83, 4393 (2003).
6 W.M. Gilmore, S. Chattopadhyay, A. Kvit, A.K. Sharma, C.B. Lee, W.J. Collis, J.
Sankar and J. Narayan, J. Mater. Res. 18, 111 (2003).
7 W.C. Goh, K. Yao and C.K. Ong, J. Appl. Phys. 97, 073905 (2005).
8 H. Hu, C. Zhu, Y.F. Lu, Y.H. Wu, T. Liew, M.F. Li, B.J. Cho, W.K. Choi and N.
Yakovlev, J. Appl. Phys. 94, 551 (2003).
9 S. Zhang and R. Xiao, J. Appl. Phys. 83, 3842 (1998).
10 J.J. Araiza, M. Cordenas, C. Falcony, V.H. Mendez-Garcia, M. Lopez, G.
Contreras-puente, J. Vac. Sci. Technol. A. 16, 3305 (1998).
11 J.M. Liu, Q. Huang, J. Li, L.P. You, S.Y. Xu, C.K. Ong, Z.G. Liu and Y.W. Du, J.
Appl. Phys. 88, 2791 (2000).
12 A.A. Voevodin, J.G. Jones and J.S. Zabinski, J. Vac. Sci. Technol. A. 19, 1320
(2001).
13 Y. Yao, S.G. Lu, H. Chen and K.H. Wong, J. Appl. Phys. 96, 5830 (2004).
14 S. Chen and I. Chen, J. Am. Ceram. Soc. 81, 97 (1998).
15 X.Y. Chen, K.H. Wong, C.L. Mak, X.B. Yin, M. Wang, J.M. Liu and Z.G. Liu, J.
Appl. Phys. 91, 5728 (1991).
100 Chapter 6 Pseudo-epitaxial PZT thin film on silicon substrate
16 T. Kumazawa, Y. Kumagai, H. Miura, M. Kitano and K. Kushida, Appl. Phys. Lett.
72, 608 (1998).
17 K. Yao, S. Yu and F. EH. Tay, Appl. Phys. Lett. 82, 4540 (2003).
18 J. S. Speck, A. Seifert, W. Pompe, and R. Ramesh, J. Appl. Phys. 76, 477 (1994).
19 K. Yao and F.E.H. Tay, IEEE Trans on Ultra. Ferro. Freq. Ctrl. 50, 113 (2003).
101 Chapter 7 Overall Conclusions
Chapter 7 Overall conclusions and future work
7.1 Conclusions
The following major conclusions have been drawn from the results on the ferroelectric Lead Zirconate Titanate (PZT) thin films in this thesis:
1. The experimental results showed that SrTiO3 (STO) is favorable for growing
epitaxial PZT film by Pulsed Laser Deposition (PLD). The growth of PZT
film deposited on STO is governed by a three dimensional island mode. We
observed a two-layer growth structure at a PZT film thickness of 40 – 50 nm.
At the early growth stage, the nuclei of the PZT film had a size of 20 – 30 nm
and a density of 1011/cm2. As the thickness was increased to 40 nm, small
grains started to merge into larger grains with a size of 100 – 120 nm. With
further increase in thickness, a dominant column-like growth mode was
observed, with an average grain size of 100 – 150 nm. This growth mode is
found to be favorable to grow thick PZT film since the film roughness is still
maintained at an acceptable level.
2. High resolution transmission electron microscope images (HRTEM) showed a
sharp interface between the PZT film and the STO substrate. Some
dislocations defects were found within 1 – 2 nm at the interface. A local lattice
misalignment was also observed. The PZT film showed perfect stacking lattice
at thickness of 20 nm and above, indicating that the misalignment in the PZT
film due to the interface stress and substrate defects was healed after stacking
for 50 monolayers.
3. We also observed that the difference in microstructure of the PZT film would
lead to significantly difference in electrical properties. The PZT film with
102 Chapter 7 Overall Conclusions
dense microstructure had a significantly high dielectric constant and remnant
polarization with low coercive field. However, morphology difference only
resulted in relatively smaller difference in loss tangent and piezoelectric
properties. The PZT film with loose microstructure could have a substantially
high piezoelectric voltage constant g33 due to the much low dielectric constant.
4. Yttria-Stabilized Zirconia (YSZ) and Yttrium Barium Copper Oxide (YBCO)
were found to be excellent buffer layer for growing perovskite conductive
oxide on silicon substrate. The introduction of a YBCO layer was essential to
obtain c-axis-oriented epitaxial La0.7Sr0.3MnO3 (LSMO) films through the
dual-in-plane orientation growth mechanism of YBCO on YSZ/Si.
5. LSMO was found to be an excellent bottom electrode for PZT film on silicon
substrate due to excellent electrical conductivity and good lattice matching
with PZT film.
6. A pseudo epitaxial PZT film was deposited on LSMO/YBCO/YSZ template
layer on silicon substrate by PLD. Although this pseudo epitaxial PZT film
had good crystallography orientation, it exhibited a granular morphology with
nano-pores distributed evenly across the whole PZT film. Despite of the
imperfect epitaxial quality, the pseudo epitaxial PZT film had a ferroelectric
polarization that was substantially larger than that of an “ideal” epitaxial film
deposited on silicon substrate. The improvement in ferroelectric properties of
the pseudo-epitaxial was attributed to the partial relief of the tensile stress by
virtue of the granular pseudo-epitaxial nature with nano-sized pores.
7. Beside high epitaxial quality, the effect of tensile stress must also be taken into
account, when achieving PZT fin films with high performances.
103 Chapter 7 Overall Conclusions
7.1 Future Work
We have developed a deep understanding on the relationship between microstructure and the properties of PZT films deposited by PLD and sol-gel methods.
We have further clarified the connections between microstructure properties and electrical and piezoelectric properties of the PZT films. There are still many questions that we can ask, including:
1. What is the relationship between the in-plane orientation of the PZT film
and its electrical and piezoelectric properties as well?
2. How can we realize a stress free but quality epitaxial film on silicon
substrate?
3. Is tensile residual stress the main cause for the low remnant polarization
of epitaxial PZT films on silicon?
4. What is the next step to integrate this pseudo epitaxial PZT film in MEMS
or microelectronics devices?
5. How to further improve the piezoelectric properties of the PZT films
deposited by PLD?
104