has continued for many decades with acute interest following a roller-coaster pattern. For example ... materials: its high energy radiation stability and amenability to wet chemical etching. [21]. ...... Status Solidi B 235, 260 (2003). [23] J. D. Ye ...
UNIVERSITY OF PUERTO RICO
SYNTHESIS AND CHARACTERIZATION OF 3d-TRANSITION METAL IONS DOPED ZnO BASED DILUTE MAGNETIC SEMICONDUCTOR THIN FILMS
By KOUSIK SAMANTA
A Dissertation in Partial Fulfillment of the Requirements for the Degree of DOCTOR OF PHILOSOPHY
Program in Chemical Physics Department of Physics Faculty of Natural Sciences
Supervised by Dr. Ram S. Katiyar, Professor
San Juan, Puerto Rico APRIL, 2009
©Copyright 2009 by Kousik Samanta
SYNTHESIS AND CHARACTERIZATION OF 3d-TRANSITION METAL IONS DOPED ZnO BASED DILUTE MAGNETIC SEMICONDUCTOR THIN FILMS
ACCEPTED BY FACULTY OF THE CHEMICAL PHYSICS PROGRAM OF THE UNIVERSITY OF PUERTO RICO IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY
03%), the mixed state of Cu2+ and Cu1+ were detected.
II
ACKNOWLEDGEMENTS First and foremost, my deepest and sincerest appreciation goes to my advisor Prof. Ram S. Katiyar for suggesting the problem, supervising the work and being a potential source of inspiration at each stage of this thesis research work. The completion of this thesis has been possible, only due to his intellectual support. I hope that that his work and his example would continue to inspire me during my entire carrier in future. I would like to express my deepest gratitude to Dr. Pijush Bhattacharya for his scientific insight, continuous support, and guidance throughout my work. I benefited tremendously from his depth of knowledge in many diverse fields. His confidence in me has always been a motivating factor to keep me striving to reach my goals. I am thankful to him for supervising this thesis and constant discussions. My sincere thanks to Prof. Luis F. Fonseca, Prof. Gerardo Morell and Prof. Carlos R. Cabrera as the potential thesis committee members, they have spent their valuable time to review my thesis, and providing necessary suggestions. I would like to extend my thanks to Prof. Carlos Rettori and Prof. K. V. Rao for the collaboration and invaluable discussion about magnetism. I like to thanks material Characterization Center (MCC) for routine structural and XPS characterizations of my samples. Thanks to the members of CCMR, Cornell University,
for
facilitating
high
resolution
transmission
electron
microscopy
measurements. I wish to thank all the non-teaching stuff of the Physics Department and EPSCoR for their cooperation during the research period I gratefully acknowledge two years financial support from the NSF fellowship and one year DOE fellowship.
III
TABLE OF CONTENTS ABSTRACT……….……………………………………………………………………...I ACKNOWLEDGEMENTS……………………………………………………………III TABLE OF CONTENTS….............................................................................................V LIST OF FIGURES….....................................................................................................IX LIST OF TABLES...................................................................................................…..XII LIST OF ABBREVIATIONS……….……………………………………………….XIII PUBLICATIONS AND PRESENTATIONS……………………………………….XIV
CHAPTER 1 INTRODUCTION………………………………………………………….…………....1 1.1 Brief Description of Spintronics……………………………………………………..1 1.2 Importance of ZnO in Optoelectronics………………………………………………2 1.3 Significance of ZnO based DMSs for Spintronics…………………………………...3 1.4 Proposed Spintronics Devices……………………………………………………....12 1.4.1
Spin-Light emitting diode (SLED)……………...………………………….12
1.4.2
Spin-Transistor……………………………………………………………...13
1.4.3
Spin-Field Effect Transistor (SFET)…………………………….……….....13
1.5 Statement of Research Problems…………………………………………………...15 1.6 References…………………………………………………………………..............16
CHAPTER 2 EXPERIMENTAL PROCEDURES…….………………………………………….....20 2.1 Materials Synthesis………………………………………………………..………20 2.2 Pulse Laser Deposition (PLD) Technique…………………….………………….20
V
2.3 Characterization Techniques…………………………………………………......22 2.3.1
X-ray Diffraction……………………………………………………….…22
2.3.2
Optical Techniques…………………………………………………….….23 2.3.2.1
Optical Transmission Spectroscopy…………………………...........23
2.3.2.2
Photoluminescence Spectroscopy………………………………......24
2.3.3
Raman Spectroscopy…………………………………………………..…..26
2.3.4
High Resolution Transmission Spectroscopy (HRTEM)………………....28
2.3.5
Magnetometry………………………………………………….………….28 2.3.5.1
SQUID Magnetometer...…………………………..………………..28
2.3.5.2
Vibrating Sample Magnetometer (VSM)……………………...…....29
CHAPTER 3 SYNTHESIS AND CHARACTERIZATION OF Zn1-xCoxO CERAMIC TARGETS AND THIN FILMS……………………………..............................................................30 3.1 Structural and Lattice Dynamical Properties of Zn1-xCoxO Ceramic Targets And Thin Films………………...………………………………………………….32 3.1.1
Introduction………………………………………………………………..32
3.1.2
Experimental Procedure…………………………………………………...33
3.1.3
Results and Discussions…………………………………………………...34 3.1.3.1
XRD Analysis………………...….………………………………...34
3.1.3.2
Raman Scattering Analysis……………………………...…............37
3.1.3.3
HRTEM Micro-Graph Analysis……………………......………….44
3.1.4
Summary…………………………………………………………………..46
3.1.5
References…………………………………………………………………47
3.2 Temperature Dependent Raman Active Optical Modes in Zn1-xCoxO Ternary Alloy…………………………………………………………………..….49 3.2.1
Introduction………………………………………………………………..49
3.2.2
Results and Discussions…………………………………………………...50
3.2.3
Summary…………………………………………………………………..58
VI
3.2.4
References…………………………………………………………………60
3.3 Optical Properties of Zn1-xCoxO Thin Films…………………………..………...61 3.3.1
Introduction………………………………………………………………..61
3.3.2
Results and Discussions…………………………………………………...64 3.3.2.1
UV-VIS Transmission Analysis………………………………..…..64
3.3.2.2 Photoluminescence (PL) of ZnO Thin Films at 77K…………..........66 3.3.2.3
Temperature Dependent PL Analysis…………..…………...............66
3.3.2.4
Photoluminescence of Zn1-xCoxO Thin Films at 77K……….......….69
3.3.3
Summary…………………………………………………………………..70
3.3.4
References…………………………………………………………............71
3.4 Magnetic Properties of Co Doped ZnO thin Films……………………………...73 3.4.1
Introduction………………………………………………………………..73
3.4.2
Results and Discussions…………………………………………………...74 3.4.2.1
Ferromagnetism in Zn1-xCoxO Thin Films………………………….74
3.4.2.2
Effect of Carrier Concentration on Ferromagnetic Properties of Co dope ZnO Thin Films…………………………………………...76
3.4.2.3
Optical Bandgap Analysis of ZCO:Al Thin Films…………………79
3.4.2.4
Oxidation State of Co in ZCO:Al Thin Films……………................81
3.4.3
Summary…………………………………………………………………..83
3.4.4
References…………………………………………………………………84
CHAPTER 4 STRUCTURAL, OPTICAL, AND MAGNETIC PROPERTIES OF Cu AND Mn DOPED ZnO THIS FILMS……………………………………………………..…......86 4.1 Microstructural and Ferromagnetic Properties of Cu doped ZnO Thin Films………………………………………………………………….....................87 4.1.1
Introduction………………………………………………………………..87
4.1.2
Experiments……………………………………………………………….88
4.1.3
Results and Discussions…………………………………………………...89 VII
4.1.3.1 X-Ray Diffraction Analysis……………………………………...…..89 4.1.3.2 Raman Scattering Analysis…………………………………………..89 4.1.3.3 High Resolution TEM Analysis……………………………...............91 4.1.3.4 X-Ray Photoelectron Spectroscopy Analysis…………………...…...92 4.1.3.5 Room temperature PL analysis of Zn1-xCuxO thin films……...……..95 4.1.3.6 Magnetic Properties of Zn1-xCuxO Thin Films………………...…….97 4.1.4
Summary…………………………………………………………………..99
4.1.5
References………………………………………………………………..101
4.2 Optical Properties of Zn1-xCuxO Thin Films…………………………..……….103 4.2.1
Introduction………………………………………………………………..103
4.2.2
Results and Discussions……………………………………………...........104 4.2.2.1 Optical Transmission Analysis……………………………………..104 4.2.2.2 Photoluminescence Analysis………………………...……………..106
4.2.3
Summary……………………………………………..…………..............108
4.2.4
References…………………………………………………………..……109
4.3 Multi-Phonon Raman Scattering in Mn Doped ZnO………………..………...110 4.3.1 Results and Discussions……………………………………………………110 4.3.1.1 XRD Analysis…………………………...………………………….110 4.3.1.2 Raman Scattering Analysis…………………………………............112 4.3.1.3 Optical Transmission Spectra Analysis…………………………….115 4.3.2
Summary………………………………………………..………………..117
4.3.3
References…………………………………..……………………………118
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LIST OF FIGURES 1.1 Schematic diagram of the ferromagnetic metal and semiconductor hetero-structure; the problem of resistance scattering at the interface and rapid decay of spin polarization away from the interface [26]………..……………………………..........5 1.2 The schematic diagram of an ideal dilute magnetic semiconductor……….………...5 1.3 Electronic configuration of 3d and 4s states of transition metal elements [29]……...8 1.4 The schematic representation of the density of electronic states available to electrons in a normal metal and in a ferromagnetic metal whose majority spin states are completely filled. E is the electron energy; EF is the Fermi level; and N(E) is density of states [30]……..…………………………………………………………………...8 1.5 Schematic diagram of SLED (p-n junction) [58]…..……………………………….11 1.6 Schematic diagram of spin transistor (n-p-n) [58]……….……………………........12 1.7 A schematic of a Spin Field Effect Transistor (Datta-Das transistor). In this device, the gate voltage is used to control the precession of spins from a ferromagnetic emitted to ferromagnetic collector [59]……….…………………………………....14 2.1. Schematic diagram of the pulsed laser deposition (PLD) set-up …………………. 21 2.2. Diagram of the experimental geometry of X-ray diffraction ……………………... 23 2.3. Schematic diagram of the possible optical transitions in the semiconductor………25 2.4. Schematic diagram of the photoluminescence (PL) set-up.………………………...25 2.5. Diagram of the optical phonon modes and their vibrational directions in the wurtzite (C6v) ZnO crystal structure …………………………………………………………27 3.1 Schematic representation of a wurtzite ZnO structure having lattice constants a in the basal plane and c in the basal direction; u parameter is expressed as the bond length or the nearest-neighbor distance b divided by c (0.375 in ideal crystal); α and β (109.47° in ideal crystal) are the bond angles [1]……..………………………….33 3.2 X-ray diffraction pattern of Zn1-xCoxO bulk targets ……………………………… 35 3.3 X-ray diffraction pattern of Zn1-xCoxO thin films…………......…………………....35 3.4 Room temperature Raman spectra of Zn1-xCoxO (x = 0-0.2) targets ……………... 38
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3.5 Room temperature Raman spectra of Zn1-xCoxO (x = 0-0.15) thin films on Al2O3 Substrates …………………………………………………………………………..38 3.6 Correlation length (L) Vs Raman peak frequency (ω) for E2low modes of Zn1-xCoxO targets for different Co concentrations …………………………………………….41 3.7 Change in full width and half maxima (FWHM) of E2low modes of Zn1-xCoxO targets for different Co concentration ……………………………………………………...41 3.8 Optical transmission spectra of Zn1-xCoxO (x=0.1, 0.15) and ZnCo2O4 thin films on Al2O3 substrates recorded at 300 K.………………………………………………..41 3.9 Identification of secondary phase of ZnCo2O4 in 10 & 15% Co doped ZnO………41 3.10 (a) and (b) Cross-sectional HRTEM micrograph of 5 and 10% Co doped ZnO thin film; insets shows the FFT of the selected area of the films…….………………….45 3.11 Room temperature Raman scattering of ZnO and Zn1-xCoxO ternary alloy…….....51 3.12 Temperature dependent Raman spectra of ZnO in the range of 80 to 800K………51 3.13 (a) and (b) are the temperature dependent frequency shift and line width change of ZnO and 3% Co doped ZnO respectively……………………………………........55 3.14 (a) and (b) show the temperature dependent linewidth change and frequency shift of E2low mode of ZnO respectively……………………………………………………57 3.15 Comparison of the temperature dependent of E2 (high) phonon lifetime in ZnO and Zn1-xCoxO alloys………………………………………………………...................57 3.16 Band structure and symmetries of hexagonal ZnO. The splitting into three valence bands (A, B, C) is caused by crystal field splitting and spin-orbit coupling [5]…..62 3.17 Optical transmission spectra of Zn1-xCoxO thin films; inset shows the variation of bandgap with Co concentrations…………………………………………………..65 3.18 Splitting of the low energy level of Co2+ in ZnO lattice, under crystal field and spin-orbit interaction [24]…………………………………………………………65 3.19 Photoluminescence spectra for ZnO thin film on Al2O3 substrate at 77 K ……… 67 3.20 Temperature dependent PL spectra for ZnO thin film; the eA0 appears due to increase of temperature …………………………………………………………...67 3.21 Photoluminescence spectra for Zn1-xCoxO (x = 0 to 0.15) thin film at 77 K ….... 69
X
3.22 Hysteresis loop (M-H) of 10% Co doped ZnO thin film at 2 K…………………..75 3.23 Hysteresis loop (M-H) of 3, 5 and 15% Co doped ZnO thin films at 2 K.………..75 3.24 Room temperature ferromagnetism in 10% Co doped ZnO thin film; upper left inset shows the TR FM of 5% Co doped sample; lower right inset shows the temperature variation of ZFC/FC measurement…………………………………..75 3.25 Room temperature in-plane magnetization, M (H), of ZCO and Al-doped ZnO thin films. Inset shows the field cool in-plane M (T) at H = 1 T for the same four films……………………………………………………………………………….77 3.26 The Curie-Weiss linear fit of the inverse paramagnetic susceptibility, χp-1, of the ZCO and Al: ZCO films for T ≤ 120 K…………………………………………...80 3.27 The optical transmission spectra of Zn0.9-xCo0.1O:Alx thin films..………………...80 3.28 The XPS spectra of Zn0.9-xCo0.1O: Alx thin films, no signature of Co0 were detected……………………………………………………………………………80 4.1 XRD spectrum of Zn1-xCuxO thin films, CuO and Cu2O related peaks in 5% Cu doped sample at 38.38 and 61.32 respectively……………………………………..90 4.2 Raman spectra of Zn1-xCuxO thin films on Al2O3 substrate, the CuO related Ag mode at 296.8 cm-1 observed in Zn0.95Cu0.05O thin film…...……………………….90 4.3 (a) High resolution TEM image showing the interface between Zn0.99Cu0.01O and Al2O3 substrate; (b) HRTEM micrograph of Zn0.99Cu0.01O film; (c) HRTEM image of polycrystalline Zn0.95Cu0.05O thin film; and (d) Lattice dislocation [dotted circle in (c)] exhibit stacking fault in the atomic row along (0001) planes. All insets show the FFT of the corresponding micrograph…………………….......................................93 4.4 XPS spectra of Cu2p peaks of Zn1-xCuxO thin films, inset shows the de-convolution of Cu 2p3/2 of Zn0.95Cu0.05O thin film……………………………………………….94 4.5 PL spectra of Zn1-xCuxO thin film at 300 K; NBE transition red shifted due to Cu doping……...……………………………………………………………………….96 4.6 M-H loop of the Zn1-xCuxO thin films at room temperature, maximum saturation magnetization (Ms) of 0.76 μB/Cu was observed in 3% Cu doped ZnO sample…...98 4.7 Schematic diagram and 3d orbitals splitting of Cu2+ (d9) in Zn2+ side of ZnO lattice under tetrahedral crystal field and spin-orbit interaction………..………………….98 4.8 Optical transmission spectrum of Zn1-xCuxO thin films, the reduction of band edge transition with Cu content (inset)..………………………………………………...105
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4.9 Photoluminescence spectrum of Zn1-xCuxO thin films at 77 K; inset shows the deconvolution of FXA and FXB transition…...………………………………………107 4.10 Temperature dependent PL spectrum of Zn0.95Cu0.05O thin film, the appearance of eA0 at the higher energy side of DAP band with elevated temperature (inset)..…107 4.11 XRD spectra of Zn1-xMnxO ceramic targets....…………………………………..111 4.12 XRD spectra of Zn1-xMnxO thin films…………………………………………...111 4.13 Raman spectra of Zn1-xMnxO ceramic targets at room temperature, inset shows the shift of E2low mode towards lower frequency side..……………………….....113 4.14 Identification of secondary spinal phase ZnMn2O4 precipitated in 10% Mn doped ZnO………………………………………………………………..……………..113 4.15 Optical transmission spectra of Zn1-xMnxO and ZnMn2O4 thin films; insert shows the increase of optical bandgap energy due to Mn doping……………................116
LIST OF TABLES 3.1 The fitting values of the anharmonic constants for the E2 (high) mode in ZnO and Zn1-xCoxO alloys…….……………………………………………………………...55 3.2 Band Structure related properties of Wurtzite ZnO………………………………...63
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LIST OF ABBREVIATIONS USED IN THIS THESIS APF
Alloy Potential Fluctuation
BMP
Bound Magnetic polaron
CT
Charge Transfer
CCD
Charged Coupled Device
DAP
Donor Acceptor Pair
DBE
Donor Bound Exciton
DMS
Dilute Magnetic Semiconductor
DHO
Damped Harmonic Oscillator
FC
Field Cooled
FFT
Fast Fourier Transformation
FM
Ferromagnetism
FX
Free Exciton
HRTEM
High Resolution Transmission Electron Microscopy
LSDA
Local Spin Density Approximation
MCD
Magnetic Circular Dichroism
NBE
Near Band Edge
PL
Photoluminescence
PLD
Pulse Laser Deposition
RTFM
Room Temperature Ferromagnetism
SQUID
Superconducting Quantum Interference Device
SFET
Spin Field Effect Transistor
SLED
Spin Light Emitting Diode
SC
Spatial Correlation
TM
Transition Metal
UV-VIS
Ultraviolet-Visible
VSM
Vibrating Sample Magnetometer
XPS
X-Ray Photoelectron Spectroscopy
XRD
X-Ray Diffraction
ZFC
Zero Field Cooled
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PUBLICATIONS AND PRESENTATIONS Thesis Related Publications: 1. Microstructural and ferromagnetic properties of Zn1-xCuxO thin films, K. Samanta, P. Bhattacharya, and R. S. Katiyar, J. Appl. Phys. 105, 113929 (2009) 2. Optical and magnetic properties of Zn0.9-xCo0.1O:Alx thin films, K. Samanta, P. Bhattacharya, J. G. S. Duque, W. Iwamoto, C. Rettori, P. G. Pagliuso, and R. S. Katiyar, Solid State Commun. 147, 305 (2008) 3. Structural and optical properties of nano-crystalline Zn1-xMnxO, K. Samanta, S. Dussan, P. Bhattacharya, and R. S. Katiyar, Appl. Phys. Lett. 90, 261903 (2007) 4. Temperature dependent E2 Raman modes in the ZnCoO ternary alloy, K. Samanta, P. Bhattacharya, and R. S. Katiyar, Phys. Rev. B 75, 035208 (2007) 5. Local and global magnetic properties of Zn1-xCoxO and Mn doped GaAs thin films, W. Iwamoto, R. R. Urbano, P. G. Pagliuso, C. Rettori, K. Samanta, P. Bhattacharya, R. S. Katiyar, J. H. D. de Silva, A. Pereira, G. de M. Azevedo, and S. B. Oseroff, IEEE Transaction on Magnetics 42, 2700 (2006) 6. Raman scattering studies in dilute magnetic semiconductor Zn1-xCoxO, K. Samanta, P. Bhattacharya, R. S. Katiyar, W. Iwamoto, P. G. Pagliuso, and C. Rettori, Phys. Rev. B 73, 245213 (2006) 7. Structural and optical properties of Zn1-xCoxO and ZnCo2O4 thin films, K. Samanta, P. Bhattacharya, R. S. Katiyar, and C. Rettori, Proc. Mat. Res. Soc. Symp. 891, EE10.09.1 (2006) 8. Optical properties of Zn1-xCoxO thin films on Al2O3 (0001) substrate, K. Samanta, P. Bhattacharya, and R. S. Katiyar, Appl. Phys. Lett. 87, 101903 92005) 9. Raman spectroscopy of V and Co doped ZnO ceramics and thin films, K. Samanta, N. Awasthi, B. Sundarakannan, P. Bhattacharya, and R. S. Katiyar, Proc. Mat. Res. Soc. Symp. 829, B2.28.1 (2005) (Ribbon Award Winner)
Collaborative Publications: 1. Self-assembled ZnO nano structure for field emission device, P. Bhattacharya, D. Varshney, K. Samanta, and R. S. Katiyar, J. Nano Res. 4, 19 (2008)
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2. The structural and magnetic properties of (In1-xFex)2O3 (0.00 ≤ x ≤ 0.25) system: Prepared by gel combustion method, O. D. Jayakumar, I. K. Gopalakrishnan, S. K. Kulshreshtha, A. Gupta, K.V. Rao, D.V. Louzguine-Luzgin, A. Inoue, K. Samanta, M. K. Singh, and R. S. Katiyar, Appl. Phys. Lett. 91, 052504 (2007)
Thesis Related Presentations: 1. Blue luminescence in Ferromagnetic Zn1-xCuxO thin films, K. Samanta, P. Bhattacharya, and R. S. Katiyar, Materials Research Society Fall Meeting, Boston, MA, Dec 1-5, 2008 (Poster Presentation) 2. Structural and Magnetic properties of Zn1-xCoxO Thin films, K. Samanta, P. Bhattacharya, C. Rettori, and R. S. Katiyar, International Conference on Magnetic Materials, Kolkata, WB, India, Dec 11-16, 2007 (Oral Presentation) 3. Structural and Luminescence properties of Cu doped ZnO thin films, K. Samanta, P. Bhattacharya, and R. S. Katiyar, Materials Research Society Fall Meeting, Boston, MA, Nov 24-28, 2007 (Oral Presentation) 4. Room temperature Ferromagnetism in dilute magnetic semiconductor Zn0.9xCo0.1O:Alx thin films, K. Samanta, P. Bhattacharya, and R. S. Katiyar, Material Science &Technology Conference and Exhibition, Detroit, MI, Sept 16-20, 2007 (Invited Talk) 5. Effect of free carriers on Ferromagnetic properties of Co doped ZnO thin films, K. Samanta, P. Bhattacharya, R. S. Katiyar, and C. Rettori, Materials Research Society Spring Meeting, San Francisco, CA, April 9-14, 2007 (Oral Presentation) 6. Structural, Magnetic, and Luminescence properties of Co substituted ZnO, K. Samanta, P. Bhattacharya, R. S. Katiyar, and C. Rettori, Materials Research Society Fall Meeting, Boston, MA, Nov 27th - Dec 1st, 2006 (Poster Presentation) 7. Optical and Magnetic properties of Zn1-xCoxO and ZnCo2O4 thin films, K. Samanta, P. Bhattacharya, R. S. Katiyar, W. Iwamoto, R. R. Urbano, P.G. Pagliuso, and C. Rettori, American Physical Society March Meeting, Baltimore, Maryland, March 1316, 2006 (Oral Presentation) 8. Structural and optical properties of Zn1-xCoxO and ZnCo2O4 thin films, K. Samanta, P. Bhattacharya, R.S. Katiyar, W. Iwamoto, R. R. Urbano, P.G. Pagliuso, and C. Rettori, Materials Research Society Fall Meeting, Boston, MA, Nov 28th - Dec 2nd, 2005 (Poster Presentation)
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9. Self assembled ZnO nanostructure using polystyrene sphere, K. Samanta, S. Bhattacharyya, P.Bhattacharya, and R.S. Katiyar, American Ceramic Society’s 107th Annual Meeting, Baltimore, Maryland, April 10-13 2005 (Oral Presentation) 10. Raman spectroscopy of V and Co doped ZnO ceramics and thin films, K. Samanta, N. Awasthi, B. Sundarakannan, and P. Bhattacharya, and R. S. Katiyar, Material Research Society Fall meeting, Boston, MA, Nov 29th - Dec 3rd 2004 (Poster Presentation)
XVI
SYNTHESIS AND CHARACTERIZATION OF 3d-TRANSITION METAL IONS DOPED ZnO BASED DILUTE MAGNETIC SEMICONDUCTOR THIN FILMS
Dedicated to My Parents and my Wife Soma For their Love, Support, and Encouragement
CHAPTER 1 INTRODUCTION 1.1
Brief Description of Spintronics Over the last 50 years there has been a revolution in microelectronics technology
from the earliest transistor to supercomputing microprocessor based on the digital logic of the electron. The foundation of the digital logic (bit) is based on the charge state of the electron. However, the continuous scaling down and more logic in microelectronics devices, would be in such a level that the each logic would be in atomic scale, and that’s the end of silicon based microelectronics roadmap [1]. The rapid miniaturization of the device is approaching in such a limit that the generated heat cannot be dissipated fast enough and the unwanted quantum-mechanical effects would prevent them from functioning properly. The utilization and control of the spin of the electron in the devices can overcome the above problem. The field of spintronics is multidisciplinary in nature; the core concept of this field is to manipulate the polarized spin degrees of freedom in solid-state systems. It combines standard microelectronics with spin-dependent effects that arise from the interaction between electrons and magnetic field. Thus, the combination of bandgap engineering and the integration of magnetic degrees of freedom offer remarkable opportunities for a new generation of devices with completely different functionalities. Moreover, the spin of an electron can be switched from one state to another much faster than charge can be moved around a circuit, spintronic devices are expected to operate faster and produce less heat than conventional microelectronic components. For the spintronics revolution to happen, however, we need to find a way to
1
manipulate, inject, and detect the spin of electrons in semiconductors, since these materials are likely to remain central in device physics for the foreseeable future. The prediction of room-temperature ferromagnetism in transition metal doped semiconductors like GaN and ZnO has sparked a race towards the development of materials that can serve as a source of spin-polarized electrons and can be actively controlled by the application of external fields [2].
The goal of spintronics is to understand the interaction between the particle spin and its solid-state environments and to make the energy efficient multifunctional devices, such as, spin-based transistor that would replace conventional transistors in integrated logic circuits and memory devices, thus allowing the miniaturization trend to continue; spin light-emitting diode (SLED) that generates left or right circularly polarized light for use in encrypted communication. Looking further into the future, spintronic devices could even be used as quantum bits, the units of information processed by quantum computers.
1.2
Importance of ZnO in Optoelectronics There has been a great deal of interest in zinc oxide (ZnO) semiconductor
material for its prospects in optoelectronics applications owing to its direct wide band gap (Eg) 3.3 eV and a large excitonic binding energy (60 meV) at 300 K. Research on ZnO has continued for many decades with acute interest following a roller-coaster pattern. For example, lattice parameters of ZnO were investigated for many decades [3-6]. Similarly, optical properties and processes in ZnO as well as its refractive index were extensively studied many decades ago [7-14]. Vibrational properties by techniques such as Raman scattering were also determined early on [15-20]. The large exciton binding energy of (60 meV) and small excitonic Bohr radius ( rB ~ 1.8 nm), which makes the excitons stable even at room temperature paves the way for an intense near-band-edge excitonic
2
emission at room temperature, and sharp transitions facilitating very low threshold semiconductor laser. It should be noted that besides the above-mentioned properties of ZnO, there are additional properties which make it preferable over other wide-band-gap materials: its high energy radiation stability and amenability to wet chemical etching [21]. Several experiments confirmed that ZnO is very resistive to high-energy radiation [22-24] making it a very suitable candidate for space applications. ZnO is easily etched in all acids and alkalis, and this provides an opportunity for fabrication of small-size devices. ZnO has recently found other applications as well, such as fabrication of transparent thin-film transistors, where the protective covering preventing light exposure is eliminated since ZnO-based transistors are insensitive to visible light. The carrier concentrations can be increased up to 2 × 1021 cm−3 by heavy substitutional doping into ZnO [25]. By controlling the doping level, electrical properties can be changed from insulator, through n-type semiconductor to metal, while maintaining optical transparency that makes it useful for transparent electrodes in flat-panel displays and solar cells. This field is also interesting by theoretical prediction [2] and perhaps experimental confirmation that the 3d transition metals doped ZnO could be the potential dilute magnetic semiconductor (DMS) with ferromagnetic Curie temperature well above 300 K for spintronics applications.
1.3
Significance of ZnO based DMSs for Spintronics In order to develop and realize spintronics devices it is essential to resolve several
issues related to the stable source of sufficient spin polarized carriers and reliable transportation of these spin polarized carriers within the devices. These impose two requirements on the system; the system must have a ferromagnetic element which 3
supports a sufficient storage of the spins polarized electrons and secondly the system should have a semiconducting portion through which conventional device operation can be performed. At a glance, the ferromagnetic metal and semiconductor hetero-structure may serve the purpose. However, the conductivity mismatch between the ferromagnetic metal and the semiconductor heterostructures, causes resistance scattering at the interface, and as a result the degree of spin polarization should be very low as shown in Fig. 1.1 [26]. Moreover, in the metal/semiconductor interface, some of the magnetic atoms diffuse to the semiconductor part. Each magnetic atom contains magnetic moment and they are randomly oriented. When the spin-polarized electrons pass through the interface, it experiences a scatter between two different spin channels, and the electrons will loose their spin-polarization shortly. In order to achieve high spin polarization with a ferromagnetic metal source, either the conductivities of the two materials must be closely matched, or the degree of spin polarization in the metal must be 100%. Neither of these solutions is easily attainable in metal-semiconductor devices [27]. The use of dilute magnetic semiconductor (DMS) as the source of spin-polarized carriers represents another approach to spin injection. The DMSs are semiconducting materials in which a fraction of the host cations can be substituted by proper magnetic ions or rare earths. In an appropriate DMS, the net electronic angular momentum of individual magnetic dopants is ferromagnetically coupled with the charge carriers. Such materials are expected to have spin-polarized states either in the valance or conduction band. The use of DMS as spin injector can eliminate the problem of conductivity matching at the interface. However, many magnetic semiconductors have Curie temperatures (Tc) that are well below room temperature, making them unsuitable for device applications.
4
Ferromagnet
Semiconductor
• Lattice mismatch • Defects • Fermi level mismatch Spin scattering ∝ resistivity mismatch Figure 1.1: Schematic diagram of the ferromagnetic metal and semiconductor heterostructure; the problem of resistance scattering at the interface and rapid decay of spin polarization away from the interface [26].
Electronics
Optics
Magnetism
Magnetic impurity Semiconductor host
Figure 1.2: The schematic diagram of an ideal dilute magnetic semiconductor
5
The enormous interest on DMS materials is due to its potential application in “Spintronics” devices, which exploit spin in magnetic materials along with charge of electrons in semiconductors. Transition metal ions with partially filled d orbitals (Mn, Fe, Co, Ni, and Cu etc.) are used as magnetic atoms in DMS. The unpaired electrons in partially filled d orbitals are responsible to exhibit magnetic behavior in DMS materials; the delocalized conduction band electrons and valence band holes interact with the localized magnetic moments associated with the magnetic atoms. The electronic structure of the host lattice is influenced by the strong hybridization of 3d orbitals of the transition metal ions with the s or p orbitals of the neighboring anions. This hybridization gives rise to the strong magnetic interaction between the localized 3d spins and the carriers in the host valence band [28]. The 3d-TM doped ZnO has attracted intense attention in the search for high TC ferromagnetic DMS materials, since Dietl et al. [2] predicted that it could exhibit ferromagnetism above room temperature upon doping with transition elements such as Mn, Co, Cu, etc. This in simple terms is due to the strong sp-d hybridization, which involves the valence and conduction band in the host, owing to small distance from its nearest neighbor and small spin dephasing spin-orbit interaction. Zinc oxide is of wurtzite structure which is formed by tetrahedral (s-p3) bonding and the TM elements have valence electrons corresponding to the 4s orbital, and have partially filled 3d shells. Generally, 3d transition-metal ions substitute for the cations of the host semiconductors, i.e., Zn sites in ZnO; the particular TM element, for example, Mn, Co, Cu etc, contributes its 4s electrons to the s-p3 bonding, and can therefore substitutionally replace the Zn in the tetrahedral bonding to form a TM2+ charge state. The 3d orbital of the Mn2+ ion is exactly half-filled with 5 electrons among the 10 available states, with an
6
energy gap between the up-spin (↑) occupied states and empty down-spin (↓) states. For other transition metals, such as Co and Cu one of the bands is usually partially filled (up or down), as shown in Fig 1.3 [29]. The TM-d band of the transition metal, hybridize with the host valence bands (O-p orbital in ZnO) to form the tetrahedral bonding. This hybridization gives rise to the exchange interaction between the localized 3d spins and the carriers in the host valence band. In this simple picture, the s band of the conduction band does not mix with the TM-d bands, but it is still influenced by the magnetic ion. The important characteristic of a ferromagnetic material is the spontaneous magnetization below the Curie temperature, also referred to as the critical temperature. As shown in Fig 1.4 in ferromagnetic materials [30], the d band is divided into spin-up and spin-down subbands, and the up and down states are displaced in energy with respect to one another, so that the spin-up band is filled first, and the spin-down states contain the remaining, if any, electrons. The difference in the number of spin-up and spin-down electrons gives rise to the observed spontaneous magnetic moment. Above TC, the ferromagnetic material looses its permanent magnetism due to thermal agitations. In order to have practical applications in functional devices, it would be desirable, to have a Curie temperature well above room temperature. Further for some device applications, it is also desirable to have carrier-induced ferromagnetism, so that the magnetic properties of the DMS can be manipulated by external means, such as through manipulation of the electron or hole concentration. The prediction of room temperature ferromagnetism could be possible in p-type TM- doped ZnO based DMS materials [2].
7
Figure 1.3: Electronic configuration of 3d and 4s states of transition metal elements [29]
Figure 1.4: The schematic representation of the density of electronic states available to electrons in a normal metal and in a ferromagnetic metal whose majority spin states are completely filled. E is the electron energy; EF is the Fermi level; and N(E) is density of states [30]
8
This in simple terms is in part due to the strong p-d hybridization, which involves the valence band in the host, owing to small distance with nearest neighbor and small spin dephasing spin-orbit interaction. Even though the common wisdom indicates hole mediated ferromagnetic exchange interaction to be dominant, Sato et al. [31] predicted that the ferromagnetic state Co2+(d7) in Co-doped ZnO could be stabilized by s-d hybridization, pointing to the possibility that high-Curie-temperature ferromagnetic materials could be realized in n-type ZnO as well. A number of approaches have been explored to synthesize the single phase transition metal doped ZnO based DMS and to realize their magnetic properties. Despite the weight of the predictions support only p-type Zn1-xMnxO leading to ferromagnetism; experimental observations of ferromagnetism for insulating Zn1-xMnxO [32, 33] and ntype Zn1-xMnxO thin films [34, 35]. In case of Co doped ZnO samples, high TC (>300 K) ferromagnetism was observed for insulating Co-doped ZnO films [36, 37]. Ando et al. [38] reported a large magneto-optical effect in Zn1-xCoxO thin films as measured by magnetic circular dichroism (MCD) spectra, suggesting Zn1-xCoxO to be suitable as a DMS material, although the p-d exchange interaction is antiferromagnetic in the samples studied. The MCD relies on the optical transitions allowed under various optical polarizations involving split off bands due to Zeeman Effect which is enhanced by sp-d exchange interaction. Ferromagnetism with T > 300K was also observed in Zn1C
x(Co0.5Fe0.5)x
O thin films prepared by magnetron co-sputtering and post annealing in
vacuum [39]. However, bulk Zn1-xCoxO has been found to be antiferromagnetic in polycrystalline powder samples prepared by both solid-state and liquid-phase reactions [40]. This antiferromagnetic behavior is likely to be associated with the formation of Co
9
clusters observed in Zn1-xCoxO powder films, together with the existence of interstitial Co atoms instead of substitutional Co on Zn sites. Sati et al. [41] has reported recently the paramagnetic behavior of plasma-assisted MBE grown single crystalline Co doped ZnO thin films below helium temperature while it’s antiferromagnetic above it. Recently, Akdogan et al. has reported two phase ferromagnetism in Co doped ZnO thin films grown on Al2O3 substrate by rf-sputtering [42]. One is due to the substitution of Co in Zn lattice side and the second magnetic phase originates from the Co cluster on the substrate. The complete paramagnetic behavior was found in Al, Co co-doped ZnO polycrystalline powder [43]; where as, Lui et al. [44] observed the paramagnetic behavior in sol-gel prepared Co doped ZnO powder below 5 K, but the co-precipitation of Al with Co doped ZnO samples were ferromagnetic even at 360 K. We have observed room temperature ferromagnetism in PLD grown Co, Al co-doped ZnO thin films [45]. There are also suggestions based on experimental data which show that homogeneous films of Zn1xCoxO
which are preferred for device applications tend to exhibit spin-glass behavior,
whereas inhomogeneous Zn1-xCoxO films are more likely to demonstrate roomtemperature ferromagnetism, suggesting perhaps that Co clusters might be the source of observed high TC ferromagnetism in Zn1-xCoxO thin films [46]. There are several contradicting reports on Cu doped ZnO samples available in the literature; where some authors have confirmed [47, 48] the occurrence of FM in this system, while others have ruled it out [49]. Even in studies, where room temperature ferromagnetism is reported, the effect of carrier type on the ferromagnetic properties is unclear [50]. Buchholz et al. found that p-type carriers are essential for realizing ferromagnetism in the ZnO:Cu
10
system but non-ferromagnetic in n-type system [50]. In sharp contrast to this, Hou et al. reported ferromagnetism in n-type ZnCuO films [51]. The mechanisms behind the observed ferromagnetism in DMS materials, particularly early on, are still not fully understood and appreciated. The theory dealing with ferromagnetism driven by the exchange interaction between carriers and localized magnetic ions was first proposed by Zener [52]. The theory indicates that direct superexchange between the magnetic ions is not ferromagnetic but the indirect super-exchange involving carrier mediation is. Therefore, the features of DMS are induced by the exchange interaction between localized d shell electrons of the magnetic ions and the delocalized band carrier states (s or p origin). Only in recent years, a large number of theoretical efforts have been undertaken to explain the detailed exchange mechanisms. Models based on the mean-field theory, first principle calculations [53, 54], and bound magnetic polaron (BMP) [55-57], etc., all have been developed to account for the magnetic properties observed experimentally, although each has its own limitations. Due to the complexity of the DMS systems based on ZnO and especially the possible presence of secondary phase precipitates, it is still difficult to find a universal theory, other than possibly full blown numerical ab initio calculations, to explain all the above phenomena observed in Zn1-xTMxO system. It is fair to state that, the state of transition metal-doped ZnO systems is still in its infancy, and it is too early to give a definitive description of the exact mechanism(s) governing the experimental observations regarding the origin of reported magnetization behavior. A better understanding of the above mention mechanisms will certainly provide the much needed guidance for material design.
11
1.4
Proposed Spintronics Devices
The goal of spintronics is to exploit the charge as well as the spin property of the electrons in the conventional semiconductor devices.
1.4.1 Spin-Light Emitting Diode (SLED) The schematic diagram of spin light emitting diode (SLED) and its band diagram (magnetic p-n junction) are shown in Fig. 1.5 [58]. The SLED in which spin-polarized electrons are injected from a ferromagnetic layer (pale blue) into a semiconductor structure (orange) recombine with holes in the active region (yellow) to produce circularly polarized light (pink, where the arrow indicates the direction of polarization), and it could be useful for encrypted communication.
Figure 1.5: Schematic diagram of SLED (p-n junction) [58].
12
1.4.2 Spin-Transistor A design for a spintronic transistor is the magnetic tunnel transistor, shown in Fig. 1.6 [58], in which the injected electrons are filtered depending on their spin as they tunnel through a thin insulating layer (red), as happens in a magnetic tunnel junction, before passing through a Schottky barrier. The output current in the "collector" semiconductor can therefore be controlled by changing the spin alignment of the "emitter" and "base" ferromagnetic layers.
Figure 1.6: Schematic diagram of spin transistor (n-p-n) [58].
1.4.3 Spin-Field Effect Transistor (SFET) The first proposed semiconductor spintronic device was the Spin Field Effect Transistor, by Datta and Das in 1990 [59] (Figure 1.7). The spin-polarized current is injected from source side of the device. The gate voltage is used to control the precession
13
of spins via the Rashba-spin orbit interaction from a ferromagnetic source to ferromagnetic drain. Since the degree of spin precession is dependent on the voltage applies, the transport through the device will be affected by the end of channel spin alignment of the current relative to that of the ferromagnetic collector at the drain end of the device. This control of the drain current through application of a gate voltage is analogous to what is seen in a charge based field effect transistor. The potential benefits of this device are the low operational currents and higher speeds than traditional FETs. This could have a great impact on the overall development of spin-based devices. Implementation of these structures has been slow due to difficulties in the fabrication and operation of these devices.
Figure 1.7: A schematic of a Spin Field Effect Transistor (Datta-Das transistor). In this device, the gate voltage is used to control the precession of spins from a ferromagnetic source to ferromagnetic drain [59].
14
1.5
Statement of Research Problem
The above mentioned review elaborates the present status of ZnO based dilute magnetic semiconductor for Spintronics applications. A large number of groups have explored the synthesis and magnetic properties of these materials in more detail using a variety of growth methods. One lesson that has been learned in these investigations is that thorough materials characterization, with a number of spectroscopic, diffraction, and imaging techniques, must be carried out in order to accurately determine what phases are present. Without such characterization, it is all too easy to interpret a ferromagnetic response as arising from the formation of a true DMS, when in reality a nonmagnetic oxide semiconductor phase containing secondary magnetic impurity phases has formed. It is often desirable from a magnetic point of view to incorporate several atomic percent of the magnetic dopant. However, doing so can result in exceeding the solid solubility limit, thereby promoting secondary phase formation. After a thorough literature survey on ZnO based DMS materials, we have identified some of the critical issues of DMS which need investigation for the realistic device applications. The main objective of this investigation is to synthesis the single phase Zn1xTMxO
based DMS thin films and to optimized the process parameters. The prime focus
was to investigate the structural and lattice dynamical properties, doping effects, electronic band structure and excitonic behavior, and to identify the origin of observed ferromagnetism in DMS materials. The total investigation in this thesis is broadly classified into these four categories.
15
1.6
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19
CHAPTER 2 EXPERIMENTAL PROCEDURES 2.1
Material Synthesis The material synthesis is one of the key features for the development and
realization of the semiconductor based spintronics applications. The ability to produce high quality single phase dilute magnetic semiconductors (DMS) is the driving parameter to investigate this materials system (DMS) for spintronics applications. The ceramic powders of Co, Mn, and Cu doped ZnO were prepared by solid solution reaction mechanism. The stoichiometric mixing of Co3O4, MnO2, and CuO with ZnO were carried out by 20 hours normal ball milling for the Co, Mn, and Cu doping, respectively. All the chemicals were 99.999% pure and parched from Alfa Aesar. The dry powders were calcined at 1000ºC for 4 hour for the Zn1-xTMxO (TM = Co, Mn, and Cu) phase formation. The calcined powders were well mixed with PVA as a binder, and the targets for the PLD were fabricated by using 1 inch dye and hydraulic press. The targets were sintered in two steps, at 650ºC for 4 h at 3º C/min to remove the organics (PVA) and then at 1200ºC for 4 h at 6 ºC/min.
2.2
Pulse Laser Deposition (PLD) Technique In the pulsed laser deposition (PLD) technique, high power laser pulses were used
to evaporate materials from the target surface such that the stoichiometry of the material is preserved in the interaction.
20
Figure 2.1: Schematic diagram of the PLD set-up
As a result, a supersonic jet of particles (plume) was directed normal to the target surface. The plume expands away from the target with a strong forward directed velocity distribution of different particles. The ablated species condense on the substrate, placed opposite to the target. A schematic diagram of the typical PLD system is shown in Fig. 2.1. An Excimer laser (KrF, 248nm, 10Hz) with laser energy of 2.5 J/cm2 was used to deposit the films. The deposition chamber was evacuated to 1 x 10-6 Torr background pressure. The substrate temperature was maintained at 650°C. Oxygen was introduced by keeping the pressure at 2-5 mTorr. The target to substrate distance was maintained at 5
21
cm. The deposition period was 17 min, and the thickness of the films grown was ~800 nm. The single crystal (0001) oriented Al2O3 was used as a substrate throughout all the thin films deposition.
2.3
Characterization Techniques
2.3.1 X-ray Diffraction Technique The X-ray diffraction technique was used to analyze crystalline structures, phase formation, presence of secondary phase, and lattice parameters of the bulk materials and thin films. This non-destructive technique can provide appropriate informations on structural properties and thin films quality. For diffraction applications, short wave lengths X-ray in the range of few angstroms to 0.1 Å are used. The web length of X-ray is comparable to the dimension of atoms and ideally suited for probing the structural arrangements of atoms in the crystal lattice. This high energy X-ray can penetrate deep in to the material and provide useful informations about the crystal structure. When X-ray photons collide with the atom, some of them deflected away from the direction where they originally travel. In the case of elastic scattering (no change in wavelength as well as energy) only momentum is transfer and the measured X-ray carries information about the atomic distribution in the lattice. Constructive interference of diffracted waves from different atoms gives the resultant intensity distribution, which is strongly modulated by the atomic distribution. The peaks in an X-ray diffraction pattern are directly related to the inter-planer spacing; the schematic diagram of X-ray diffraction of a periodic lattice is shown in Fig.2.2.
22
2d sin θ =
θ d
Figure 2.2: Diagram of the experimental geometry of X-ray diffraction
For a given set of lattice plane with an inter-planer spacing d, the condition for diffraction (peak) to occur can be written as (Bragg’s Low): 2d Sinθ = nλ. In this work, the XRD measurements were carried out by Siemens D5000 X-ray difractometer using Cu Kα radiation of 1.543 Å.
2.3.2 Optical Techniques These approaches provide informations related to electronic band structure of the material and the interaction with the electromagnetic radiation. It’s also providing the informations about band-edge and defects related transitions.
2.3.2.1
Optical Transmission Spectroscopy
Optical transmission consists of shining a UV-VIS light source through a sample and observing the spectrum of transmitted light on the other side. This technique is useful
23
to study the absorption related features that may or may not show up during luminescence measurements but it can provide an indication of the optical band-edge within a semiconductor material. In the case of dilute magnetic semiconductors thin films, it can be used to observe transitions related to internal d-shell transmissions, which very often end up in the ultraviolet or the infrared radiation. The UV-Visible spectrometer (Perkin Elmer model Lambda 2S) was used to perform the transmission on ZnO and Zn1-xTMxO thin films samples in the range of 300 to 800 nm of wave length using the combination of Tungsten and Deuterium lamp.
2.3.2.2
Photoluminescence Spectroscopy
One of the most important and versatile techniques to determine the information about electronic band structure, excitonic character, and more importantly the defects within the sample, is photoluminescence. In this process, electron-hole pairs (exciton) are generated by the application of an incident laser beam on the surface of a sample. These excitons recombine often through the radiative transition back to the ground state. Some of the basic observed recombination pathways are shown in Fig.2.3. By measuring the wavelength of the emitted photon from the observed recombination, information can be derived about the electronic band structure, donor and acceptor levels, defect types, impurities, crystalline quality, and defect densities within the materials system. The schematic diagram of the experimental setup used in a photoluminescence study is shown in Fig.2.4. The photoluminescence (PL) was excited using the ultraviolet (UV) 351 and 325 nm line of Ar+ and HeCd lasers, respectively.
24
Figure 2.3: Schematic diagram of the possible optical transitions
Figure 2.4: Schematic diagram of the photoluminescence set-up
25
The emitted light was collected by a T64000 Raman spectrometer from Horiba, Inc. utilizing a 0.64m monochromator with a 2400 g/mm grating with an accuracy of 2 cm-1 and detected with an LN2-cooled charge-coupled devices (CCD) camera. The sample temperature was varied from 77 to 300 K using a continuous flow liquid N2 optical cryostat.
2.3.3 Raman Spectroscopy Raman scattering spectroscopy is another technique to investigate light wavematter interaction to probe certain materials properties. Specifically, Raman scattering measures the interaction of light via inelastic scattering from an incident laser beam off of a material. One method for inelastic scattering is to transfer this energy into lattice vibrations or phonons. The energy of these lattice vibrations is quantized and a function of the local bonding and atoms involved in the structure. Thus, by measuring the energy transferred to or from phonons to photons, which is manifested as a Stokes or Anti-stokes shift in the in-elastically scattered light source; valuable information regarding the crystalline quality and lattice dynamics of the material can be gained. Fig.2.5 shows the optical phonon modes and their vibrational direction in the wurtzite ZnO lattice. The modes that can be observed are highly sensitive to the polarization of the incident light and the orientation of the crystal upon which the laser light is incident. Raman spectroscopy measurements in this work were performed by using a Jobin-Yvon T64000 Triple-mate instrument. The radiation of 514.5 nm from Ar+ laser was used for excitation to perform Raman measurement.
26
E1
E2L
E2H
Zn O
c-axis
A1
B1L
B1H
Figure 2.5: Diagram of the optical phonon modes and their vibrational directions in the wurtzite (C6v) ZnO structure
27
2.3.4 High-resolution Transmission Electron Microscopy (HRTEM) The cross sectional TEM samples were prepared by mechanical process. The samples were attached (with glue) with a smooth surface, then mechanically thinned both side by using diamond papers (30μm to 0.25μm) and finally ion milled. The HRTEM measurements we carried out by using FEI Tecnai F20 system operated at 200 keV.
2.3.5 Magnetometry In order to determine the macroscopic magnetic properties of the transition metal (TM) doped ZnO thin films, we did the SQUID and Vibrating Sample Magnetometer (VSM) measurements. The SQUID measurements were carried out in Brazil and the VSM measurements in our laboratory. SQUID Measurements were taken between 2 K and 350 K and VSM measurements at 300K.
2.3.5.1
SQUID Magnetometer
Superconducting interference device (SQUID) magnetometer is the standard measurement technique for highly sensitive magnetization studies. In this instrument, the DMS samples are drawn through a coil of superconducting wire in the presence of a magnetic field. The moving magnetic field from the sample induces a current in the wire, which through signal processing can be analyzed and converted to a signal proportional to the magnetization of the sample. The noise floor on this technique can be as low a 10-8 emu, which would make it ideal for studies of dilute magnetic systems. The superconducting magnet and coil must be cooled to cryogenic temperatures, as the critical temperature for the superconducting wire coil is 20 K. In these experiments,
28
SQUID measurements were performed using Quantum Design MPMS-5 magnetic property measurement system. In a majority of experiments, the magnetic field of the sample was applied in parallel to the plane of the sample.
2.3.5.2
Vibrating Sample Magnetometer (VSM)
Another form of magnetometer that was used in these studies was vibrating sample magnetometer (VSM). In this setup, the sample is mounted on a sample tail and placed between the coils of an electromagnet. The magnet supplies a field, which the sample is oscillated at a known, low frequency. The moving magnetic field from the sample induces a current in the pick up coils, which is sent, amplified, and converted to a known magnetic signal. This technique is nice because it is inexpensive and fast for routine measurements. However, it is much more difficult to perform low temperature studies, as it needs an additional cryostat. In addition, with these dilute samples, the magnetization of the sample is very often much less than that of the sample holder, such that the data must be background corrected in order to pull out the data from the sample from the measurement noise. In addition, the noise floor on this measurement technique is two orders of magnitude greater than what it is for the SQUID measurement technique, so it is difficult to measure very weakly magnetic samples. In this work, measurements were performed using a Lakeshore 7404 Vibrating Sample Magnetometer, with the magnetic field applied perpendicular to the plane of the thin film.
29
CHAPTER 3 SYNTHESIS AND CHARACTERIZATION OF Zn1-xCoxO CERAMIC TARGETS AND THIN FILMS This chapter is devoted to synthesis of Co doped ZnO ceramic targets, optimization of the parameters of the PLD grown thin films, structural and the study of lattice dynamical, optical, and magnetic properties of Zn1-xCoxO thin films. We have investigated the effect of calcination temperature to synthesize single phase Zn1-xCoxO ceramic powder. We have calcined the stoichiometric mixture of ZnO and Co3O4 in the temperature range of 500°C to 1100°C and obtained complete phase formation at 900°C to 1000°C. At 1100°C we observed that it starts melting, depending on the concentration of Co in ZnO. We have also optimized the process parameters (temperature, oxygen pressure, and laser energy) for the thin film deposition by PLD technique. The optimized parameters for the highly c-axis oriented thin films on single crystalline (0001) oriented Al2O3 substrates were 650°C temperature, 2-5 mTorr oxygen pressure, and 350 mJ laser power. The PLD grown Zn1-xCoxO thin films were polycrystalline and highly oriented. The lattice dynamical studies by Raman scattering showed the structural stability for up to 800K and no secondary phase was detected with Co doping up to the 10%. The low temperature photoluminescence showed the basic ZnO excitonic structure and the free exciton was stable even at room temperature. The SQUID measurements (2–300K) of Zn1-xCoxO thin films showed room temperature ferromagnetism with maximum saturation magnetization of 1.1 μB/Co in 10% Co doped ZnO thin film. Further increase of Co concentration reduced the saturation magnetization abruptly. 30
The effect of carrier concentration on ferromagnetic property was investigated by controlled co-doping with Al to increase the carrier concentrations. The ZFC/FC measurements confirm the ferromagnetic nature of Zn1-xCoxO thin films for up to 350 K.
31
3.1
Structural and Lattice dynamical Properties of Zn1-xCoxO Ceramics and Thin Films
3.1.1 Introduction Zinc Oxide is a II-VI compound semiconductor whose iconicity resides at the borderline between covalent and ionic semiconductor. At ambient conditions, the thermodynamically stable phase of ZnO is wurtzite with a hexagonal unit cell and it belongs to the space group of C6v4. A schematic representation of the wurtzite ZnO structure is shown in Fig. 3.1 [1]. The structure is composed of two interpenetrating hexagonal-close-packed (hcp) sub-lattices, each of which consists of one type of atom displaced with respect to each other along the threefold c-axis by the amount of u = 3/8 = 0.375 (in an ideal wurtzite structure) in fractional coordinates. Each sub-lattice include four atoms per unit cell and every atom of one kind (group-II atom) is surrounded by four atoms of the other kind (group VI), or vice versa, which are coordinated at the edges of a tetrahedron. Recently, ZnO alloying with 3d transition metals (TM) have attracted much attention as a dilute magnetic semiconductor (DMS), with room temperature ferromagnetism for spintronic applications [2-6]. In the host wurtzite ZnO, the isovalent transition metals (Co2+, Mn2+, Cu2+ etc.) are the substitutes at the Zn cationic sites. The coupling of the localized d electrons of the TM with the host semiconducting band gap leads to a number of exciting properties, such as magneto-optical and magneto-electrical effects [7-9]. There are several theoretical arguments as well as experimental reports that predict the Zn1-xCoxO is the most promising DMS material for room temperature ferromagnetism [10-11]. However, the major drawbacks with the experimental reports are often the parameter windows that are very narrow with a poor reproducibility.
32
Figure 3.1: Schematic representation of a wurtzite ZnO structure having lattice constants a in the basal plane and c in the basal direction; u parameter is expressed as the bond length or the nearest-neighbor distance b divided by c (0.375 in ideal crystal); α and β (109.47° in ideal crystal) are the bond angles [1].
For instance, recent controversial reports [12] reveal that the ferromagnetism in Zn1xCoxO
is not an inherent property, but because of the segregated secondary phases or
clustering related to Co, Co3O4 and the isomorphic ZnxCo3-xO4 phase. Non-destructive characterization technique ‘Raman spectroscopy’ is extensively used to study the substitution of Co in the ZnO host lattice and to identify the impurity phases.
3.1.2 Experimental Procedure The ceramic targets of Zn1-xCoxO (x=0-0.15) and ZnCo2O4 for PLD were prepared by solid solution reaction mechanism by using ZnO and Co3O4 powders. The
33
powders of ZnO and Co3O4 were mixed with stoichiometric amounts (x = 0.01, 0.05, 0.10, 0.15, 0.20) and ball milled for 24 h. The dry powders were calcined at 1000ºC for 4 h and the pellets were sintered at 800ºC for 4 h at 3º C/min and then at 1200ºC for 4 h at 6ºC/min. No intentional carrier doping was used. High quality Zn1-xCoxO thin films were grown on single crystalline (0001) oriented Al2O3 substrate by pulsed laser deposition (PLD) technique using an Excimer laser (KrF, 248nm, 10Hz) with laser energy of 2.5 J/cm2. The deposition chamber was initially evacuated to 1 x 10-6 Torr base pressure. The substrate temperature was maintained at 650°C. Deposition was done under oxygen, at a pressure of 2 mTorr, for 17 min yielding films of about 0.8 μm thickness. The crystal structure and the phase formation of ZnO and Zn1-xCoxO ceramics and thin films were characterized with the X-ray diffraction technique using Siemens D5000 X-ray difractometer (XRD) with CuKα radiation. The Raman scattering studies were performed in the backscattering geometry using Jobin-Yvon T64000 Triple-mate instrument. The radiation of 514.5nm from a Coherent Argon ion laser was focused to ~ 2 μm in diameter on the samples. An LN2-cooled charge-coupled device (CCD) system was used to collect and process the scattered data.
3.1.3 Result and Discussions 3.1.3.1
Structural Analysis
The X-ray diffraction (XRD) patterns of Zn1-xCoxO (x = 0-0.15) ceramic targets and thin films, grown on Al2O3 (0001) substrates by PLD technique are shown in Fig. 3.2 and Fig. 3.3, respectively.
34
(201) (112) (200)
(103)
(110)
(102)
(101)
(002)
(100)
Intensity (a. u)
15%Co 10%Co 5%Co 3%Co 1%Co ZnO 30
40
50
2 θ (degree)
60
70
Al2O3
(0002)
Figure 3.2: X-ray diffraction pattern of Zn1-xCoxO bulk targets
Intensity (a. u)
15% Co 10% Co
5% Co 3% Co 1% Co ZnO 30
35
40
45
50
55
2θ (degree) Figure 3.3: X-ray diffraction pattern of Zn1-xCoxO thin films
35
All of the diffraction peaks (Fig. 3.2) from the targets were corresponding to hexagonal ZnO structure. There was a small increase in a-axis value corresponding to the shift of (1000) peak position towards higher diffraction angle and a decrease in c-axis value corresponding to the leftward shift of (0002) peak position in the Zn1-xCoxO targets with the increase of Co concentration, which is consistent with the reported value by Risbud et al. [13]. The intense (0002) peak from the Zn1-xCoxO thin films clearly shows the same wurtzite structure of the film as in pure ZnO, with highly c axis orientation (Fig. 3.3). No additional peaks were observed, indicated that there was no structural change and/or additional phase formation due to incorporation of Co in ZnO in the limit of XRD detection. The (0002) peak position gradually shifted towards higher diffracting angle with increase of Co concentration up to 10%; further increase of Co concentration (~ 15%) the (0002) peak shows a shift in reverse direction. This feature indicates that Co is going to the Zn lattice side up to 10% and after that it’s going to either intersatial position or forming Co nano-clusters or undetected (in XRD) secondary phases in the sample. The c-axis value in the thin films decreased from 0.5195 to 0.5167 nm with increase of Co concentrations (up to 10%) as calculated from (0002) peak position. The decrease of caxis is understandable because of the difference in Co2+ ionic radius (0.058 nm) and that of tetrahedrally coordinated Zn2+ (0.06 nm) [14]. The evolution of the a and c cell parameters, and the unit cell volume of the Zn1-xCoxO samples as a function of Co concentration were also obtained. The substitution of divalent, high-spin Co in tetrahedral coordination of Zn results the changes in the cell parameters and the cell volume of the ZnO host. Interestingly, while the substitution results in a decrease in the c-axis (in keeping with the smaller radius of Co2+); the a-axis actually increases. If the Co2+ ions
36
were in an octahedral environment in the wurtzite structure, it would be signaled by a significant increase in the cell parameters since octahedral Co2+ has a radius between 0.65 Å (low spin) and 0.745 Å (high spin) [13].
3.1.3.2
Raman Scattering Analysis
The wurtzite structure of ZnO has the space group C46v with two formula units per primitive cell with all atoms occupying C3v sites. Each Zn atom is tetrahedrally coordinated to four O atoms and vice versa. The numbers of optical modes for the ZnO structure are given by [15]
Γ = A1 (z, z2, x2 + y2) + 2B1 + E1 (x, y, xz, yz) + 2E2(x2 – y2, xy).
Where, B1 modes are silent in Raman scattering, A1 and E1 modes are polar and hence, exhibit different frequencies for the transverse-optical (TO) and longitudinal-optical (LO) phonons, because of the macroscopic electric field associated with the LO phonons. The non-polar E2 modes have two frequencies, namely E2high and E2low, associated with the motion of oxygen (O) atoms along with zinc (Zn) sub-lattice vibrations [15]. The first order Raman modes for ZnO and Co doped ZnO ceramic targets (used for thin film deposition) are shown in Fig 3.4. We observed five normal modes at 99.8, 379, 408.7, 437.7, and 584 cm-1, corresponding to E2low, A1(TO), E1(TO), E2high, and E1(LO), respectively. The Raman scattering studies of Zn1-xCoxO bulk and thin films at room temperature revealed that there were several additional features due to the Co substitution (Fig. 3.4 and Fig. 3.5) in ZnO.
37
low
high
E2
E2
Intensity (abr. units)
M
20%Co 15%Co
A1(TO)
10%Co
E1(LO)
5%Co
E1(TO)
3%Co
1%Co
ZnO
100
200
300
400
500
600
700
-1
Raman shift (cm ) Figure 3.4: Room temperature Raman spectra of Zn1-xCoxO (x = 0-0.2) targets
S
low
E2
high
Intensity (abr. unit)
A1-TO
E2
AM
S
S S AM
Co15%
Co10% Co5%
Co1% ZnO
100
200
300
400
500
600
700
-1
Raman Shift (cm ) Figure 3.5: Room temperature Raman spectra of Zn1-xCoxO (x = 0-0.15) thin films on Al2O3 substrates
38
The E2low mode of ZnO was shifted towards the lower frequency (up to 10%Co) and there was an increase in FWHM for up to 10% of Co substituted ZnO; the FWHM nearly saturated on the further increase (15 and 20%) of Co concentration.
The broad band
centered at 540 cm-1 and an additional mode at 470 cm-1 in ceramic targets were observed in Co doped samples (Fig. 3.4). The shift of E2low mode was detected as 1.50 and 1.30 cm-1 corresponding to the bulk and thin films samples. Moreover, the FWHM was found to increase gradually up to 10% and saturated in higher Co (15 and 20%) concentration in ZnO. The backscattering Raman spectra for highly c-axis oriented ZCO thin films clearly show (Fig. 3.5) E2high and E2low modes besides strong optical modes of Al2O3 substrates. The atomic substitution of Co in ZnO host lattice induces structural disorder. This disorder breaks the translational symmetry of the allowed phonons of the host lattice and leading to the contribution of q≠0 phonons to the Raman line shape, corresponding to the finite size effect. The disorder-induced effects (lower frequency shift and broadening) in Zn1-xCoxO thin films were explained by alloy potential fluctuations (APF) using a spatial correlation (SC) model [16-20]. In an ideal crystal, the region over which the spatial correlation function of the phonon extends is infinity. When the crystal is alloying, the spatial correlation region of the phonon becomes finite owing to the potential fluctuation of the alloying disorder, which gives rise to the relaxation of q=0 selection rule in Raman scattering. The assumption of a Gaussian attenuation factor exp(-2r2/L2), where L is the diameter of the correlation region, leads to an average over q with a similar weighting factor exp(-q2L2/4) upon Furrier transformation. It successfully used to account for the q vector relaxation related to the finite size effect [16] and the structural disorder [17]. We
39
assumed a finite spatial correlation region in the alloying material and then the Raman intensity at a frequency ω can be written as [21],
4πq 2 exp( − q 2 L2 / 4)dq , 0 [ω − ω ( q )]2 + [ Γ / 2]2 0 1
I (ω ) ≅ ∫
(1)
Where, q has the unit of 2π/a (q = q' ± 2nπ/a), a is the lattice constant, n is the reflective index of the material, and Γ0 (=3.66 cm-1) is the FWHM of E2low mode of undoped ZnO Raman line. Assuming one-dimensional linear chain model, the dispersion relation for wurtzite ZnO structure can be written as follows by assuming the analytical mode relationship,
ω(q) = A + B Cos (πq),
(2)
Where, A = 73.8 cm-1 and B = 26 cm-1 for the E2low phonon dispersion according to the ab initio phonon dispersion relation calculated for ZnO [22]. Considering the correlation length, L, as an adjustable parameter, we obtained the value of L by fitting the Raman line shape of E2low band. The estimated L values corresponding to 1, 3, 5 and 10 % Co doped ZnO are 18.5, 16.4, 15.1, and 12.8 nm, respectively. Figure 3.6 shows that the E2low phonon peak shifts to the lower frequency side as the correlation length decreases; the agreement in both experimental and theoretical cases is quite good. The FWHM of E2low phonon of Zn1-xCoxO is shown in Fig. 3.7. The FWHM increases with increase in Co concentration up to 10%, from 3.66 to 5.85 cm-1 and nearly saturated for further increases of Co concentration.
40
99.6
Th
-1
Frequency shift (cm )
99.3
3%Co
99.0 −ο− Ex
1%Co 5%Co
98.7
Co
10%Co
98.4 98.1 97.8 97.5 97.2 96.9 96.6 96.3
6
8
10
12
14
16
18
20
22
24
Correlation length (L) in nm Figure 3.6: Plot of Correlation length (L) Vs Raman peak frequency for E2low modes of Zn1-xCoxO targets for different Co concentrations. The solid line represents the theoretical values using equation (1) and (2).
6.0
-1
FWHM (cm )
5.5 5.0 4.5 4.0 3.5 0
5
10
15
20
Co Concentration (%) Figure 3.7: Full width and half maxima (FWHM) of E2low modes of Zn1-xCoxO targets for different Co concentration
41
The additional Raman modes at 333 cm-1 and 542 cm-1 were detected in ZCO samples. These additional modes belong to the disorder induced multiphonon scattering. It was interesting to note that with the increase of Co concentrations up to 20%, the intensity of the multiphonon mode at 540 cm-1 and E1 (LO) at 584 cm-1 was increased substantially. In thin film spectra a broad shoulder around 548 cm-1 was clearly evident on Co substitutions (Fig. 3.4). Similar increase in the intensity of multiphonon mode in the range of 500-600 cm-1 was also reported for P+ implanted ZnO, which was attributed to the defect induced band [23]. They showed that after annealing the intensity of this multiphonon mode was reduced. Manjon et al. [24] claimed that the Raman mode around 580 cm-1 corresponding to B1 (high) silent mode of wurtzite ZnO using ab initio calculations. This mode could be observed in disorder activated Raman scattering due to relaxation of the Raman selection rules produced by the breakdown of the transnational symmetry. We tried to explain the increase of E1 (LO), besides defects, as due to resonant Raman effect at sub-bandgap absorption related to d-d transition of Co in ZCO samples. The optical absorption spectra ZCO clearly showed strong absorption band in the subbandgap (1.8-2.3 eV) region related to d-d crystal field splitting and the charge transfer absorptions as shown in Fig. 3.8. The Raman excitation source, Ar+ laser, has energy of 2.4 eV (514.5 nm), which is above the sub-bandgap states. This increase of E1 (LO) mode was due to the electron-phonon interaction. Recently, Sahoo et al. observed the same phenomena in ZnO nano-particles excited by various excitations in the range 2.41 eV to 3.815 eV [25]. Several recent results have been reported on anomalous Raman modes in doped and alloyed ZnO bulk and thin films grown by different techniques [26, 27].
42
Transmission (%)
80 ZnCo2O4
10%Co
60
15%Co
40
20
d-d transitions
0 1.75
2.00
2.25
2.50
2.75
3.00
3.25
Photon energy (eV) Figure 3.8: Optical transmission spectra of Zn1-xCoxO (x=0.1, 0.15) and ZnCo2O4 thin films on Al2O3 substrates.
L
Intensity (a. u)
E2
H
E2
15% Co 10% Co ZnCo2O4 100
200
300
400
500
600
700
800
900
-1
Raman shift (cm ) Figure 3.9: Identification of secondary phase ZnCo2O4 in 10 & 15% Co doped ZnO
43
To understand the additional modes (AM) due to Co substitution we have also studied ZnCo2O4 bulk and thin films on Al2O3 substrates. The possible secondary phases in Co substituted ZnO are Co clusters, Co3O4 and/or its isomeric compound ZnxCo3-xO4 as mentioned earlier. A comparison of Raman spectra for 10 and 15% Co doped ZnO with ZnCo2O4 ceramic targets are presented in Fig. 3.9. In the case of ceramic targets, for Zn1xCoxO
we were able to identify the additional modes mainly due to ZnCo2O4, however
for thin films; we were unable to identify any mode related to these secondary phases.
3.1.3.3
HRTEM Micro-graph Analysis
The micro-structure, interface quality, and defects or dislocations were investigated by HRTEM micrograph analysis. Fig. 3.10 (a) and (b) shows the crosssectional TEM micrograph of 5 and 10% Co doped ZnO thin films on Al2O3 substrates. The interface in both cases showed the interface-reaction between sample and substrate. The 5% Co doped sample showed that the film is highly ordered and nearly single crystalline. No defects or dislocations were observed; the lattice fringes are around 0.28 nm and selected area FFT performed on the micrograph [inset of Fig. 3.10 (a)] showed spots corresponding to the film only. The 10% Co doped sample also showed the good crystalline quality [Fig. 3.10 (b)]. However, some dislocations are observed at the lower side of the micrograph.
44
5% Co
0.28 nm
¯
5 nm
Al2O3
Figure 3.10: (a) Cross-sectional HRTEM micrograph of 5% Co doped ZnO thin film; inset shows the FFT of the selected area of the film
10% Co
0.28
5 nm Figure 3.10: (b) Cross-sectional HRTEM micrograph of 10% Co doped ZnO thin film; inset shows the FFT of the selected area of the film
45
3.1.4 Summary We prepared single phased Co doped ZnO ceramic targets (for PLD) by solid state reaction mechanism. Thin films of Co doped ZnO were grown on Al2O3 (0001) substrate by PLD technique. The films were highly c-axis oriented and free from any secondary phase as observed in XRD measurements. Extensive Raman scattering studies were carried out for structural analysis and to identify the phase separation. The bulk Raman spectrum showed single phased material up to 10% of Co doping in ZnO. In higher Co concentration, the secondary phase of ZnCo2O4 was detected; where as XRD could not detect the trace amount of this phase separation. In thin films, only non-polar E2 modes were detected besides the strong substrate signal. According to Raman selection rule, only E2 and E1(LO) modes were allowed in the highly c-axis oriented ZnO samples. The E2(low) mode was found to shift towards lower frequency side and FWHM brooded gradually up to 10% of Co doping. The shift and broadening of E2(low) modes towards the lower frequencies were considered due to the alloy potential fluctuations, which were analyzed using a spatial correlation model and the results agreed with the experimental data. It also clearly showed that the Co ions were occupying Zn substitutional sites. The HRTEM micrograph showed that the samples were nearly single crystalline, defects free. Although, small interface reaction between sample/substrate and lattice dislocation were detected in 10% Co doped ZnO sample. Our experimental studies give a unique and novel approach to establish the upper limit of uniform and homogeneous Co substitution in the ZnO lattice as compared to other techniques, which generally give a statistical average of such information.
46
3.1.5 References [1]
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48
3.2
Temperature Dependent Raman Active Optical Modes in Zn1-xCoxO Ternary Alloy
3.2.1 Introduction The substitution of magnetic ions in ZnO host lattice for fabrication of ternary alloys, substantially affects its lattice dynamical properties. The inelastic scattering of light, such as Raman Scattering from the material has provided a great deal of information concerning the optical modes of vibrations at the center of Brillouin zone [1]. In pure materials both the frequencies and the linewidth of a phonon change with temperature, which could be attributed to the vibrational potential energy [2-3]. The Raman spectrum of a solid is highly sensitive to the lattice temperature; therefore Raman microprobe could be used as a temperature probe [4]. This capability creates exciting opportunities and provides a powerful tool for the study of the fundamental interactions. The parameters of Raman mode, such as frequency, linewidth and hence lifetime provide the basic information of lattice dynamics. The evolution of these parameters with temperature for the material under study is also important for its practical applications. Lin et al. have investigated the disorder induced first order Raman active LO mode and their anharmonic properties of Zn1-xBexSe thin films on GaAs substrates [5]. The anharmonic properties of low frequency E2 mode of ZnO was investigated by Aku-Leh et al. [6]; but not many reports are available on lattice dynamics and anharmonic properties of II-VI based ternary alloy semiconductors. To the best of our knowledge, there are no reports on temperature dependent anharmonicity in high frequency E2 mode of ZnO and Co substituted ZnO, covering temperature range up to 800K.
49
The details investigations on the behavior of Raman active non-polar E2 modes of technologically important ZnO and Co doped ZnO samples were carried out in the temperature range from 80 to 800K. The strong temperature dependence of its frequency and linewidth were explained in terms of the anharmonic decay mechanism dominated by three and four phonon coupling.
3.2.2 Result and Discussion The Raman active optical modes of ZnO and Zn1-xCoxO at room temperature are shown in Fig.3.11. The presence of non-polar E2 modes in all the samples indicates that the Co doping does not change the Wurtzite structure of ZnO. It was found that the increase of Co concentration in pure ZnO gradually decreases the intensity of E2high mode and the peak position shifted towards the lower frequency. The intensity of the multiphonon peak at 538 cm-1 increased with the increase of Co concentration and it merged with the E1 (LO) phonon. The translational symmetry of wurtzite ZnO breaks by the insertion of Co2+ in the Zn2+ cationic site in ZnO host lattice. The disorder induced frequency shift was explained by alloy potential fluctuation (APF) using spatial correlation model [7]. Our main focus in this work was to investigate the effect of temperature on the frequency, linewidth, and damping of the zone centered E2high and E2low modes in Zn1-xCoxO ternary alloys. The frequency of E2 modes decreased gradually and the linewidth becomes broadened with increase of temperature (Figure 3.12). The analysis for higher Co concentrations (>3%) in ZnO was excluded due to rapid decay of E2high peak intensity with temperature.
50
low
E2
Intensity (arb. units)
E1(LO)
high
E2
M
10%Co
E1(TO) A1(TO)
3%Co
1%Co
ZnO
100
200
300
400
500
600
700
-1
Raman shift (cm ) Figure 3.11: Room temperature Raman scattering of ZnO and Zn1-xCoxO ternary alloy
low
ZnO
E2
high
Intensity (arb.units)
E2
800K 750K 650K 550K 450K 350K 250K 80K
100
200
300
400
500
600
700
-1
Raman shift (cm ) Figure 3.12: Temperature dependent Raman spectra of ZnO in the range of 80 to 800K
51
In order to account the accurate value of the phonon frequency (not the peak center) and linewidth of E2 modes at different temperature, we use damped harmonic oscillator (DHO) model to fit Raman profile [inset of Fig. 3.13 (a), (b)] as [8],
I (ω ) =
χ 0 Γ0ωω o2 ( n + 1) (ω02 − ω 2 ) 2 + ω 2 Γ02
(1)
Where, n = exp(−hω / k BT ) − 1 , is the phonon occupation number, ω 0 is the peak frequency, Γ0 is the linewidth, and χ0 is related with the peak intensity. The peak frequency of the E2high mode with temperature is depicted in Fig. 3.13 (a) and (b). The frequency shift with temperature decreased linearly above 150 K. The decrease of the phonon frequency with increase in temperature can be explained by perturbation model according to which the frequency shift is mainly due to the effect of thermal expansion and the anharmonic coupling to other phonon. We can write the temperature dependent Raman frequency shift as fallows [9],
ωT = ω0 + Δωe (T) + Δωd (T),
(2)
Where, ω0 is the harmonic frequency of the optical mode, Δωe (T) is the contribution from thermal expansion; it depends on the alloy composition and mode polarization. The third term Δωd (T) is due to the anharmonic coupling of phonons of other branches. The term
Δωe (T) can be written as, T
Δωe (T) = − ω 0 γ ∫ [α c (T ) + 2α a (T )]dT , 0
52
(3)
Where, ω0 is the harmonic frequency of high frequency E2 mode, the Gruneisen parameter of the corresponding mode in ZnO γ = 1.66 [10], αc = 2.49×10-6 K-1 and αa = 4.31×10-6 K-1 are the parallel and perpendicular thermal expansion coefficient of ZnO, respectively [11]. For the anharmonic coupling term Δωd (T), we can model it by taking into account cubic anharmonic phonon coupling (three phonon process) and quartic (four phonon process) terms in the anharmonic Hamiltonian:
⎡ ⎤ 2 Δωd(T) = C ⎢1 + ⎥+ h ω k T exp( / 2 ) − 1 0 B ⎣ ⎦ ⎡ ⎤ 3 3 D ⎢1 + + + Higher order terms 2 ⎥ ⎣ exp(hω 0 / 3k B T ) − 1 {exp(hω 0 / 3k B T ) − 1} ⎦
(4)
The anharmonicity parameter C and D are taken as an adjustable parameter to fit the experimental data. The E2 (high) phonon cannot decay in to either two LA (TA) phonons, because the energy gap between the acoustic and optical phonon branches is more than one half of the E2high phonon energy, or one TO and one LA (TA) phonons, because E2high is at the lower edge of the optical phonon branch. Therefore, only the three-phonon (fourphonon process) has been taken into account in the anharmonic coupling term Δωd (T) for E2high, where we have considered the symmetric decay to the frequency ω0/3. The solid lines in Fig. 3.13 are the fitting results of the experimental Raman frequency shift with temperature using Equation 2. The current model describes well the downshift of phonon
53
frequency with increasing temperature. The fitting parameters D and ω0 have been listed in Table 3.1, along with the frequency temperature coefficient (dω/dT) of different composition. The temperature coefficients, dω/dT, lie in a narrow margin for all the composition. In order to identify the different temperature dependent behavior of volume expansion and four-phonon process, we have also displayed the net contribution of these effects to the Raman shift. The linewidth analysis is more difficult to achieve since different contributions have to be considered. Among these, the finite resolution of the spectrometer [12], the disorder induced effects in the alloy [13] and the broadening due to the anharmonic decay of the phonons are the major factors. The alloy potential fluctuation also induced asymmetric Raman line broadening [14], and such a broadening is independent of temperature. The thermal expansion is a manifestation of the lattice anharmonicity that has its own shifts in the frequency, but not on the linewidth. The pure anharmonic broadening is due to the temperature dependent phonon damping, which mainly arises from the decay into phonons with lower energies. Similar to the phonon frequency shift, the damping can be calculated by assuming the asymmetric decay into two phonons and symmetric decay in to three phonons. Due to the lack of asymmetric two-phonon (three-phonon process) decay in E2high, the temperature dependent phonon linewidth can be written as [1],
Γ(T) = Γ0 + Γd,
⎡ ⎤ 3 3 Γd = B ⎢1 + + 2 ⎥ ⎣ exp(hω 0 / 3k B T ) − 1 {exp(hω 0 / 3k B T ) − 1} ⎦
54
(5)
(6)
442
20 18
-1
440
16
438
14 436 12 434
10
432
8
430 428
6 420
430
440
450
460
100 200 300 400 500 600 700 800
(b)
440
3% Co
20 18
438
16
436
14
434
12
432
10 8
430
6
428
4
Linewidth (cm-1)
ZnO
(a)
Linewidth (cm ) Frequency shift (cm-1)
Frequency shift (cm-1)
442
410
420
430
440
450
460
470
100 200 300 400 500 600 700 800
Temperature (K)
4
Temperature (K)
Figure 3.13: (a) and (b) are the temperature dependent frequency shift and line width change of ZnO and 3% Co doped ZnO respectively. The experimental data points are extracted by fitting the Raman line shape with the equation (1).
TABLE 3.1: The best-fit values of anharmonic constants for the E2 (high) mode in ZnO and Zn1-xCoxO alloys
Sample
ω0 (cm-1)
B (cm-1)
D (cm-1)
Γ0 (cm-1)
dω/dT (cm-1K-1)
ZnO
441.5
0.28
-0.135
5.4
-0.014
1%Co
441.0
0.3
-0.142
6.2
-0.015
3%Co
440.2
0.35
-0.302
6.6
-0.0174
55
Where, Γ0 is the harmonic linewidth independent of temperature and it is caused by the defects, including isotopic mixture. Γd is the damping part induced due to the fourphonon process (quartic anharmonicity). The parameter ‘B’ is the constant related to this decay process; for ZnO, Γ0 = 5.4 cm-1 and B = 0.28 cm-1. Figure 3.13 (a) and (b) show that the agreement between the calculated curve, given by the solid line, and the experimental points for ZnO and ZCO alloy are reasonably good. The fitting parameter
Γ0 and B are also listed in Table 3.1. The anharmonic constant B and D are numerically greater for ZCO alloys than for ZnO, which indicates that the degree of anharmonicity experienced by E2high mode in ZCO is greater than ZnO. From the figures we can see the faster temperature dependent change (slightly greater slope) in frequency for ZCO than for ZnO. In ZnO the anharmonicity depends only on temperature, where as in the case of ZCO alloys, there are two components of anharmonicity, one corresponding to the temperature and the other due to the compositional disorder. The disorder-induced anharmonicity is also sensitive to the temperature; the increase of temperature produced a large number of phonons, causing in to the increased probability of inelastic scattering between the phonons and the substitutive atoms [5]. If the substituting atoms increase, the probability of inelastic (anharmonic) decay of phonons becomes larger. Thus increasing temperature make the compositional-disorder-induced anharmonicity more obvious. In the case of E2low, the situation is quite different, the linewidth and frequency of this mode behaves practically harmonic with respect to the temperature for the case of both ZnO and ZCO samples. The fluctuations observed in Fig. 3.14 (a) and (b) show no particular trend and most of them are within the experimental uncertainty. The E2low average frequency is 98.98 ± 0.5 cm-1, and its average linewidth is 3.5 ± 0.3 cm-1.
56
8
103.5 -1
3.5 ± 0.3 cm
(a)
-1
Frequency (cm )
-1
FWHM (cm )
6 5 4 3
-1
98.98 ± 0.5 cm
(b)
7
102.0 100.5 99.0 97.5
2 96.0
1
100
200
300
400
500
600
700
100
800
200
300
400
500
600
700
800
Temperature (K)
Temperature (K)
Figure 3.14: (a) and (b) shows the temperature dependent linewidth change and frequency shift of E2low mode of ZnO respectively. The fluctuations observed in figure show no particular trend and most of them are within the experimental uncertainty, i.e. the E2low mode is harmonic or nearly so.
2.0
ZnO 1Co 3Co
Lifetime (ps)
1.8 1.6 1.4 1.2 1.0 0.8 0.6 0.4 100
200
300
400
500
600
700
800
Temperature (K) Figure 3.15: Comparison of the temperature dependent of E2 (high) phonon lifetime in bared ZnO and Zn1-xCoxO alloys
57
These experimental results indicate that the E2low mode is harmonic or nearly so, i.e., its potential energy is quadratic in the interatomic displacements. This harmonic behavior is unchanged with Co substitution, because of the very similar mass and ionic radius of Zn2+ and Co2+ ions. It is well known that the damping is usually related to the lifetime of the decay process of involved phonons. The decay time of the phonon with temperature in ZnO and ZCO alloys for E2high modes is shown in Fig. 3.15. The estimated lifetime can be calculated by τ =1/(πcΓ), where c is the velocity of light and Γ is the FWHM of the mode. This gives τ = 1.82 ps at 80K and 0.65 ps at 800K for pure ZnO, where for the case of 1% and 3% Co doped ZnO, the lifetime goes down to 1.74 ps and 1.53 ps at 80K, respectively. Where as, the lifetime of 1% and 3% Co doped ZCO sample merge above 300K. The shorter lifetime of high frequency E2 mode of ZCO alloys clearly demonstrates the more possibility of decay than the ZnO, consistent with the above arguments for faster decay in ZCO alloys.
3.2.3 Summary We used Raman microprobe to investigate the anharmonic behavior of non-polar E2 modes of ZCO ternary alloys in the temperature range 80 to 800K. We analyzed the observed temperature dependent frequency and linewidth characteristics in this alloy by using a statistical model, involving thermal expansion and symmetric three-phonon coupling. In case of E2high mode the frequency shift towards the lower energy side was analyzed in light of the theory of anharmonic phonon-phonon interaction and thermal expansion of the lattice; whereas, the linewidth behavior was analyzed in terms of anharmonic effect of three-phonon decay mechanism. It was found that the higher
58
phonon anharmonicity for ZnCoO alloys compare to pure ZnO and it increased with the compositional disorder. The low temperature lifetime of E2 phonon in ZnO, 1% and 3% Co doped ZnO samples were found to be 1.82, 1.74 and 1.54 ps, respectively. But, in the case of E2low mode, the line width and frequency behaved practically harmonic with respect to temperature and independent of Co substitutions.
59
3.2.4 References [1]
M. Balkanski, R. F. Wallis, and E. Haro, Phys. Rev. B 28, 1928 (1983)
[2]
P. G. Klemens, Phys. Rev. 148, 845 (1966)
[3]
T. R. Hart, R. L. Aggarwal, and B. Lux, Phys. Rev. B 1, 638 (1970)
[4]
H. Fujimori, H. Komatsu, K. Ioku, S. Goto, and M. Yoshimura, Phys. Rev. B 66, 064306 (2002)
[5]
L. Y. Lin, C. W. Chang, W. H. Chen, Y. F. Chen, S. P. Guo, and M. C. Tamargo, Phys. Rev. B 69, 075204 (2004)
[6]
C. Aku-Leh, J. Zhao, R. Merlin, J. Menendez, and M. Cardona, Phys. Rev. B 71, 205211 (2005)
[7]
K. Samanta, P. Bhattacharya, R. S. Katiyar, W. Iwamoto, P.G. Pagliuso, and C. Rettori, Phys. Rev. B 73, 245213 (2006)
[8]
J. Menendez and M. Cardona, Phys. Rev. B 29, 2051 (1984)
[9]
H. Tang and I. P. Herman, Phys. Rev. B 43, 2299 (1991)
[10]
D. G. Mead and G. R. Wilkinson, J. Raman. Spectroscopy 6, 123 (1977)
[11]
H. Iwanaga, a. Kunishige, and S. Takeuchi, J. Mater. Sci. 35, 2451 (2000)
[12]
P. Verma, S. C. Abbi, and K. P. Jain, Phys. Rev. B 51, 16660 (1995)
[13]
H. H. Burke and I. P. Herman, Phys. Rev. B 48, 15016 (1993)
[14]
P. Parayanthal and F. H. Pollak, Phys. Rev. Lett. 52, 1822 (1984)
60
3.3 Optical Properties of Zn1-xCoxO Thin Films 3.3.1
Introduction Optical studies are an essential tool for examining the crystalline quality, defect
states, and formation of impurity bands due to the incorporation of transition metals into the lattice. This section is devoted to the substitutional effects on the optical properties of ZnO. The wide band gap (direct 3.37 eV) semiconductor ZnO sees a vivid research interest because of its possible applications in blue/UV light emitting diodes, lasers devices, and UV sensors, which are currently based on GaN [1-7]. The advantage of ZnO over GaN is the much higher excitonic binding energy (60 meV) compared to GaN (25 meV) [8], which makes the excitons stable even above the room temperature. Due to this excitonic nature of ZnO, many research groups attributed stimulated emission at room temperature (RT) with very low thresholds [9-18]. The high exciton binding energy and the small excitonic Bohr radius (aB ~ 1.8 nm) may give RT excitonic lasing under considerably low thresholds [13, 15-17]. It is possible to increase excitonic binding energy in ZnMgO/ZnO/ZnMgO quantum wells [19]; and the band gap energy can be tuned from 3.3 eV (at room temperature) to 4 eV by alloying with Mg [20]. Zinc oxide is an excellent emitter which crystallizes in wurtzite symmetry and there exist numerous photoluminescence investigations of near band gap recombination. In high quality bulk crystals or thin films the luminescence line width of excitonic recombination is as narrow as 40 μeV and many fine spectroscopic details can be observed. The conduction band of ZnO is constructed from s-like symmetric state ( Γ7c ). The valance band is a p-like state. It splitted into three bands by crystal field and spin orbit interaction, which dominate the near bandgap intrinsic absorption and emission spectra [5].
61
The symmetry of the upper valence sub-band called A-band is (Γ7) corresponding to the heavy hole, the middle one is B-band (Γ9) corresponding to light hole, and the lower most symmetry Γ7 (C-band) originated due to the crystal field splitting [21]. The optical properties of the intrinsic excitonic transition process in ZnO, effect of TM doping on excitonic behavior, formation of defects level, and inter-band transitions were investigated in details by the optical transmission spectroscopy (UV-VIS) and low temperature photoluminescence spectroscopy.
Figure 3.16: Band structure and symmetries of hexagonal ZnO. The splitting into three valence bands (A, B, C) is caused by crystal field splitting and spin-orbit coupling [5].
62
Table 3.2: Band Structure related properties of Wurtzite ZnO [5]
63
3.3.2
Result and Discussion
3.3.2.1
UV-VIS Transmission Spectra Analysis
The UV-visible transmission spectra of ZnO and Zn1-xCoxO thin films are shown in Figure 3.17. Optical transmission measurements of the Zn1-xCoxO samples confirm the interatomic transitions within the divalent Co atoms. The undoped ZnO thin film shows a sharp absorption edge at 373.9 nm (3.31 eV), which is lower than the expected room temperature band edge of ZnO at ~360 nm (3.4 eV); this reduction is due to the thermal broadening of the band edge and by the large concentration of intrinsic donor impurity states during the film growth process. The absorption edge shifts towards the lower energy side with increasing Co concentration (Fig. 3.17). In order to calculate the band gap energy of the thin films, we assumed the absorption coefficient α ∝ -ln T corresponding to the direct band gap of the wurtzite structure [22]. We draw a plot of [α (hυ )]2 against the photon energy hν . The decrease of optical band gap in Zn1-xCoxO thin films at room temperature is due to the sp-d exchange interaction between the band electrons and the localized d electrons of the Co2+ ions. The s-d and p-d exchange give rise to the negative and positive corrections to the conduction and valance band edges, respectively, leading to the band gap narrowing [23]. Besides band edge transition, additional absorptions at 1.88, 2.03, and 2.19 eV are observed in all Co doped ZnO thin films. These bands are due to the inter-atomic d-d transitions in the high-spin d7 state of the Co2+ ions. These transitions are corresponding to the 4A2 (F) to 4T2 (F), 4T1 (F), and 2E (G), respectively [24].
64
10%Co
75
ZnO
60 45 30 15 0 1.6
Band edge transition (eV)
Transmission (%)
90
1%Co 5%Co
3.4 3.2 3.0 2.8 2.6 0
2.0
2 4 6 8 Co Conct. (%)
2.4
10 12
2.8
3.2
3.6
Photon energy (eV) Figure 3.17: Optical transmission spectra of Zn1-xCoxO thin films; inset shows the variation of bandgap with Co concentrations.
Figure 3.18: Splitting of the low energy level of Co2+ in ZnO lattice, under crystal field and spin-orbit interaction [24]
65
3.3.2.2
Photoluminescence Properties of ZnO Thin Film at 77 K
Photoluminescence (PL) is one of the most powerful techniques to investigate the exciton structure of the semiconductor. There exist several reported data on excitonic transitions in ZnO at wide range of temperature [25-28]. Figure 3.19 shows the PL spectrum of ZnO thin film on Al2O3 substrate in the fundamental excitonic region at 77K. The emission line at 3.368eV is considered as the free excitons (FX); the radiative recombination of an electron from conduction band to the valance band gives this transition. The optical transition at 3.354 eV is due to the donor bound-exciton (D0X). The discrete electronic energy levels in ZnO band gap are generated by the dopant or defects in semiconductor material. The type of defects and band structure of the semiconductor material influences the electronic states of the bound exciton. Neutral or charged donors and acceptors can form bounded exciton. The neutral shallow donorbound exciton (DBE) normally dominates in the low-temperature (4 K) PL spectrum of high-quality ZnO films. The most intense peak at 3.307 eV is attributed as donor acceptor pair (DAP) recombination with the binding energy of 124 meV. The first and second LO phonon replica of the DAP band were at 3.235 and 3.164 eV, respectively.
3.3.2.3
Temperature dependent Photoluminescence Analysis
Temperature dependent PL measurements are used to study the temperature evolution of the excitons behavior in ZnO PL spectra. Temperature variation of PL spectrum of ZnO thin film is shown in Fig. 3.20.
66
PL Intensity (abr. units)
λ = 325 nm
DAP
T = 77 K
0
DX
FX
1LODAP 2LODAP LO
LO
2.95 3.00 3.05 3.10 3.15 3.20 3.25 3.30 3.35 3.40 3.45
Photon energy (eV) Figure 3.19: Photoluminescence spectra for ZnO thin film at 77 K
ZnO λ = 325 nm
0
PL Intensity (a. u)
eA
300 K 273 K 200 K 175 K 150 K 125 K 100 K 77 K
3.05 3.10 3.15 3.20 3.25 3.30 3.35 3.40 3.45
Photon energy (eV) Figure 3.20: Temperature dependent PL spectra for ZnO thin film; the eA0 appears due to increase of temperature
67
Due to the increase of temperature the intensity of free exciton (FX) at 3.368 eV increases and dominates above 200 K, because of the ionization of impurities that used to form bound exciton at low temperature [29]. The PL intensity of D0X transition decreases with increasing temperature, and the transition energy shows a redshift. This redshift of the D0X transition is due to the decrease of ZnO band-gap energy with temperature. The DAP transition shows a different behavior with temperature. The transition energy is almost unchanged or slightly redshifted up to 150 K, after that it becomes the part of the broad emission band as the temperature increases up to 300 K. For the DAP emission, a new peak emerges at the higher energy side with increase of temperature. This feature is typical for DAP transition, and the new peak at higher energy side is caused by band-toimpurity transition [30, 31]. The impurity involve in this transition must have high binding energy; keeping in mind, the higher binding energy of the acceptor in ZnO, this band-to-impurity transition is attributed to the conduction band-to-acceptor transition (eA0), in which the free electrons in the conduction band recombine with acceptors. The transition energy ( E eA ) of eA0 can be express as [31], E eA = E g − E A + k B T / 2 , where, Eg
is the band gap energy, EA is the binding of the acceptor. The band gap energy at a particular temperature can be estimated from the FX transition energy and the free excitonic binding energy (60 meV). The acceptor binding energy at 125 K has been calculated as 109 meV using EFX = 3.358 eV, EeA = 3.314 eV. This acceptor binding energy from the PL measurement date is in good agreement with the value of 107 meV, which was estimated by Zhang et al. [29].
68
3.3.2.4
Photoluminescence of Zn1-xCoxO Thin Films at 77 K
The PL spectra of Zn1-xCoxO thin films at 77 K are shown in Fig 3.21. No appreciable change in the peak position of free exciton was detected in the near bandedge structure. There is a shift of 3 meV was observed in DAP peak position towards the lower energy with increase of Co concentrations. The optical absorption spectra showed that the band-edge decreased with increase in Co concentration in Zn1-xCoxO thin films. The band-edge shifted to the lower energy side with increase in Co concentration is mainly due to the sp–d exchange interactions between the band electrons and the localized d electrons of the Co+2 ions substituting Zn ions [23]. Similar bandgap decrease was also reported for other II-VI solid solutions with 3d materials, such as Zn1-xMnxSe [32]. In this case, the low temperature (10-100K) PL spectra were red shifted with respect to ZnSe for low Mn concentration. No such shift was observed in case Zn1-xCoxO thin films at 77 K.
DAP
0
(D ,X)
PL Intensity(a.u)
77K λ = 325 nm
1LODAP 15%Co
FX
2LODAP
10%Co 5%Co 1%Co ZnO
3.10
3.15
3.20
3.25
3.30
3.35
3.40
Photon energy (eV) Figure 3.21: Photoluminescence spectra for Zn1-xCoxO (x = 0 to 0.15) thin film at 77 K
69
3.3.3 Summary Optical properties of Zn1-xCoxO thin films were investigated by optical transmission and photoluminescence spectroscopy. The optical transmission spectra showed the decrease of near band edge (NBE) transition with increase of Co concentration. The interband d-d transitions were observed in Co doped ZnO samples, which confirmed the substitution of Co2+ at Zn2+ lattice site. However, no appreciable red shift of near band-edge (NBE) transition was observed in PL spectrum of Zn1-xCoxO thin films at 77 K. The reduction of band-edge transition in optical transmission spectra is due to the sp-d exchange interaction. We observed well-resolved PL spectrum of ZnO and Co doped ZnO thin films. With the increase of temperature the appearance of eA0 emission at the higher energy side of the peak at 3.309 eV confirmed the characteristics of DAP transition rather than accepter bound exciton.
70
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B. P. Zhang, N. T. Binh, Y. Segawa, K. wakatsuki, and N. Usami, Appl. Phys. Lett. 83, 1635 (2003)
[30]
D. J. As, F. Schmilgus, C. Wang, B. Schottker, D. Schikora, and K. Lischka, Appl. Phys. Lett. 70, 1311 (1997)
[31]
J. F. Wang, D. Masugata, C. B. Oh, A. Omino, S. Seto, and M. Isshiki, Phys. Status Solidi A 193, 251 (2002)
[32]
R. B. Bylsma, W. M. Becker, J. Kossut, U. Debska, and D. Yoder-Short, Phys. Rev. B 33, 8207 (1986)
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3.4
Magnetic Properties of Co Doped ZnO Thin Films
3.4.1 Introduction The dilute magnetic semiconductor (DMS) becomes a new area of interest in condensed matter. This is due to its rich physics and potential applications in the field of spintronics [1]. Particularly, these materials are interesting for the development of spin based information storage and data processing devices [2, 3]. The basic requirement for practical applications is to achieve Curie temperatures (Tc) well above the room temperatures [4]. The recent discovery of high-Tc ferromagnetism (FM) with large magnetic moments per transition metal (TM) ion in oxide semiconductors such as ZnO [5-7] and TiO2 [8], and the controversial results [9-12] challenges our understanding of magnetism in DMS. Recently, Sati et al. [10] reported the paramagnetic behavior of plasma-assisted MBE grown single crystalline Co doped ZnO thin films bellow helium temperature and above it is antiferromagnetic in nature. The complete paramagnetic behavior was found in Al, Co co-doped ZnO polycrystalline powder [13]; where as, Lui et al. [14] observed the paramagnetic behavior in sol-gel prepared Co doped ZnO powder bellow 5 K, but the co-precipitation of Al with Co doped ZnO samples were ferromagnetic even at 360 K. In order to develop magneto-optic and spin-electronic devices that could operate in ambient conditions, it is essential to understand the origin of high-Tc ferromagnetism in these DMS materials. The Co doped ZnO (ZCO) is one the most prominent DMS for spintronics applications. This is due to strong FM coupling that is expected for this system. In this section we studied the effect of additional free carrier (donor) by Al co-doping with Co on the ferromagnetic properties of ZCO thin films using pulsed laser deposition technique.
73
3.4.2 Result and discussion 3.4.2.1
Ferromagnetism in Zn1-xCoxO Thin Films
The magnetic properties of Co doped ZnO thin films are investigated by SQUID magnetometer measurement. The magnetic hysteresis loop (M-H) of Zn1-xCoxO thin films at 2K are shown in Fig.3.22 and Fig.3.23. The saturation magnetization at 2K for 3, 5, and 10% Co substituted ZnO was observed as 0.11, 0.24, and 1.2 μB/Co, respectively. It is evident from these data that the 10% of Co substitution produces the maximum saturation magnetic moment (1.2 μB/Co), and the Ms value is found to decrease for 15% Co substitution. This decrease of Ms can be correlated to the inverse of frequency shift and the saturation of FWHM of E2low Raman mode for 15% Co substitution. Therefore, the additional Co concentrations (>10%) are not substituting at the Zn site in the ZnO host lattice. Harima et al. [15] have also reported phase separation at 15% Co substitution in ZnO from the Raman spectra. The substitution of Co at Zn site was found to be proportional to Co concentration till 10% and with further increase of Co it preferably formed ZnCo2O4. The inset in Fig. 3.23 shows the hysteresis loop for ZnCo2O4 film, which is completely different than those found for the Co doped ZnO films. The ZnCo2O4 film presented a very small Ms value (~ 4 x10-4μB/Co) and a large coercive filed of Hc ≈ 5 kOe. The magnetization results of Fig. 3.22 and Fig.3.23 make us to conclude that the ferromagnetic loops in our Zn1-xCoxO thin films (with Co concentration up to 10%) are not due to the precipitation of any secondary phase formation of ZnCo2O4 or Co cluster in our films. The room temperature ferromagnetism was achieved in 5 and 10% Co doped ZnO thin films as shown in Fig.3.24.
74
1.5
0.5
ZnO 3% Co 5% Co 15% Co
20 15
T=2K
10
M (10−2μB/Co)
M (μB/Co)
1.0
Zn1-xCoxO/Al2O3 x = 0.10 Thickness 800 nm
0.0 -0.5
5
T=2K 0 0.08
-5
-10
0.00 -0.04
-15
-1.0
ZnCo2O4
0.04
-0.08
-20
-15 -10 -5
-1.5 -15
-10
-5
0
5
10
-15
15
-10
-5
0
5
0
10
5
10
15
15
Field (kOe)
H (kOe) Figure 3.22: Hysteresis loop (M-H) of 10% Co doped ZnO thin film at 2 K
Figure 3.23: Hysteresis loop (M-H) of 3, 5 and 15% Co doped ZnO thin films at 2 K
1.5 8 6
-5
M (10 emu)
1.0
4
ZnO 5%Co 2K 300 K
2
Zn0.9Co0.1O 300 K 350 K
0 -2
0.5
-6 -8 -16 -14-12 -10 -8 -6 -4 -2 0 2 4 6 8 10 12 14 16
H (kOe)
0.0 1.4
H=1T
Zn1-xCoxO (x=0.1)
-0.5
1.2
1.1
-1.0
1.0
-1.5
ZFC FC
1.3
M (μB/Co)
M (μB/Co)
-4
-15
-10
-5
0
0
50
5
100
150 200 T (K)
250
10
300
350
15
H (kOe) Figure 3.24: Room temperature ferromagnetism in 10% Co doped ZnO thin film; upper left inset shows the TR FM of 5% Co doped sample; lower right inset shows the temperature variation of ZFC/FC measurement.
75
The lower right inset shows the ZFC/FC measurement of 10% Co doped ZnO thin film. The temperature verses magnetization shows almost temperature independent saturation magnetization above 120 K. This phenomenon confirms that the observed ferromagnetism is not mediated by the exchange interaction through free carriers. Moreover, our magnetization results are in general agreement with the current interpretation that the residual impurities (Zni and O2v) act as shallow n-type donors, allowing a long-range magnetic coupling between the Co2+ localized magnetic moments via the conduction band [16].
3.4.2.2
Effect of Carrier Concentration on Ferromagnetic properties of Co doped ZnO Thin Films
The origin of room temperature ferromagnetism in DMS still remains controversial. The theoretical approach suggests that there is double exchange interaction leading to ferromagnetism in ZnCoO [17]. Park et al. claimed that the residual hydrogen impurities in ZnO could mediate a strong short-range ferromagnetic interaction [18]. Large concentrations of mobile carriers were not observed in high Curie Temperature ZnCoO thin films report by Yan et al. [19]. In our case 10% Co substituted ZnO also did not show high mobile carrier concentration besides the residual impurities. The high-Tc ferromagnetic property in Co-doped ZnO (ZCO), mediated by donor impurity band was tested by controlled introduction of shallow donors (Al) in the Zn0.9xCo0.1O:Alx
(x = 0.005 and 0.01) thin films. We have obtained the maximum saturation
magnetic moment Ms (4 emu/cc) at ~ 300 K for the high resistive (ρ >103 Ω-cm) ZCO thin films. In the case of Al doped ZCO sample, we have also obtained room temperature FM, however, the saturation decreased significantly to Ms ≤ 1 emu/cc (Fig. 3.25). 76
Magnetization (emu/cc)
4
Zn0.9Co0.1 O:Alx T = 300 K x=0 x = 0.005 x = 0.01
2
0 H = 1 T (FC) H // Film plane
20
15
-2 10
5
-4
0
0
50
100
150
200
250
300
T (K)
-15
-10
-5
0
5
10
15
H (kOe) Figure 3.25: Room temperature in-plane magnetization, M (H), of ZCO and Al-doped ZnO thin films. Inset shows the field cool in-plane M (T) at H = 1 T for the same four films.
77
The high-Tc FM in high resistive Zn0.9Co0.1O film (no Al doping) may be explained by the phenomenon of the hybridization of Co ion states and the charge carrier introduced by shallow donors at the Fermi level [20]. The isovalent Co2+ in ZnO itself doesn’t introduced carriers; the interstitial Zn (Zni) is the most common donor in ZnO. The charge transfer (CT) transition of an electron from the ZnO donor (Zni) level to Co2+ state creates a hole in donor level (acceptor level) and a shallow ionized (Co+) donor state close to the conduction band, these two oppositely charge particles generated through the CT transition can form bound state due to long range Coulomb interaction. This transition + ( Co 2+ ⇒ Co + + hVB ) is refer to as ligand valance band to transition metal CT transition
(LVB-TM-CT), and a considerable amount of charge is shifted between the dopant and semiconductor in this process [21]. Considering Co+ as a shallow donor in the ZCO system the exchange should work in the same way as the donor impurity band exchange (BMP model) proposed by Coey et al. [22]. The introduction of 0.5 and 1.0% Al in ZCO sample generates the free carriers and the resistivity drops to ~ 0.033 and 0.02 Ω-cm respectively. These free carriers degenerate the Fermi level to the conduction band, which is supported by our optical band gap analysis. These free carriers in Al doped ZCO samples may destabilize the magnetic polarons, and reduced the saturation magnetization. The temperature dependence of the field cool (H = 1 T) magnetization, M (T) is shown in the inset of Fig. 3.25. The temperature independent saturation magnetization above ~ 120 K, suggested that delocalized free carriers and metallic conductivity are not responsible for the room temperature ferromagnetism in ZCO and Al doped ZCO samples. The low temperature tail (T < 120 K) of M (T) was fitted to a paramagnetic Curie-Weiss type of magnetic susceptibility. The obtained Co2+ concentration and Curie-Weiss
78
temperature, θCW, are given in the Fig. 3.26. Notice that the Co2+ ions in the paramagnetic state are weakly antiferromagneticaly coupled (-3.5 K ≤ θCW ≤ -1.3 K). Since the extracted Co2+ concentrations are all close to the films nominal Co2+ concentration (~ 10%), we may assume that the amount of Co2+ ions that actually participate in the longrange FM ordering is only of the order of 1 % in every film. Based on this assumption, and the data of Fig. 3.25, we estimated that the number of Bohr magnetons per Co atom drops from ~ 1 µB/Co in the ZCO film to ~ 0.2µB/Co in the Al doped ZCO thin films.
3.4.2.3
Optical Bandgap Analysis of ZCO:Al Thin Films
The substitution of Co cations in the tetrahedral site of the wurtzite structure was confirmed by the optical transmission spectra of ZCO and Al doped ZCO thin films shown in Fig. 3.27. The characteristic optical absorption bands have been identified with d-d transition of high spin Co2+ 3d7 4F ion in tetrahedral oxygen coordination [23]. The transitions at 1.87, 2.02, and 2.19 eV, corresponding to the 4A2 to 4T1 (4P), 2E (2G), and 4
T1 (4F), respectively [24]. The orbital triplet final state is splitted into a single and
doublet by the trigonal component of the crystal field at the tetrahedral site [25]. The optical band gap of the ZCO and Al doped ZCO films were calculated by using the relation α 2 ∝ (hν − E g ) , where α is the absorption coefficient. The optical band gap gradually increases with the increase of Al content of 0.5 and 1.0 % by 38 and 54 meV, respectively with compare to the optical band gap energy (3.309 eV) of ZCO film.
79
(x10-3emu/mole-ZnO)-1 paramag. -1
χ
1.5 Zn0.9Co0.1 O:Alx H = 1 T (FC) H // film plane
1.2
x=0 x = 0.005 x = 0.01
0.9
0.6 Co
ΘCW
% 11.2(2) 8.0(2) 7.1(2)
K -3.5(2) -1.3(2) -1.5(2)
Al x 0 0.5 1.0
0.3
0.0
0
20
40
80
60
100
120
Temperature (K) Figure 3.26: The Curie-Weiss linear fit of the inverse paramagnetic susceptibility, χp-1, of the ZCO and Al: ZCO films for T ≤ 120 K. The obtained Co concentration, %, and CurieWeiss paramagnetic temperature, θCW, are also given.
Co10% Al0.5% Al1.0%
60
40
12 10 -1
6 4
2
20
8
9
α (x 10 cm )
Transmission (%)
80
2 0 3.0
0
3.1
3.2
3.3
3.4
3.5
3.6
3.7
Photon energy (eV)
1.6
2.0
2.4
2.8
3.2
Photon energy (eV) Figure 3.27: The optical transmission spectra of Zn0.9-xCo0.1O:Alx thin films
80
This increase of band gap in Al doped ZCO sample may be due to the Burstein-Moss (BM) shift [26], this energy band widening effect resulting from the increase of Fermi level in the conduction band of degenerate semiconductor.
3.4.2.4 Oxidation State of Co in ZCO:Al Thin Films The x-ray photoelectron spectroscopy (XPS) was used to characterize the charge state of Co ions in the films. Figure 3.28 shows a comparison of the Co2p core level photoemission spectrum of Zn0.9-xCo0.1O:Alx (x = 0, 0.005, and 0.01) thin films with the metallic cobalt (Co0) plate. It is determined from the XPS studies that the Co ions are in Co2+ formal oxidation state in the both ferromagnetic ZCO and ZCO:Al thin films. The binding energy of Co 2p3/2 core levels for Co-O bonding of the samples is in the range of 781.68 to 781.83 eV, and the energy difference of Co 2p3/2 and Co 2p1/2 between 15.55 to 15.7 eV. The comparison with Co0 and the binding energy difference between Co 2p3/2 and 2p1/2 excluded the possibility of the metallic cobalt cluster (Co0) in the material, because, if Co exists as metallic clusters in the thin films, the energy difference would have been 15.05 eV. On the other hand, if Co is surrounded by oxygen, these differences should be ≥ 15.5 eV [27]. Therefore, the chemical state of Co in our ZCO film can be considered in 2+ oxide state, not as elemental Co.
81
Intensity (arb. units)
2p3/2
2p1/2 Al 1.0%
Shake-up satellites Al 0.5%
Co 10%
Metalic Co
775
780
785
790
795
800
805
810
Binding Energy (eV) Figure 3.28: The XPS spectra of Zn0.9-xCo0.1O: Alx thin films, no signature of Co0 were detected.
82
3.4.3 Summary We obtained the room temperature ferromagnetic (FM) properties in PLD grown Co-doped and Al, Co co-doped ZnO (ZCO) thin films. The observed FM is the intrinsic property, due to the substitution of Co in the Zn lattice site of ZnO structure. In the 10% Co substituted sample, we obtained maximum saturation magnetization of ~ 1.0 μB/Co at room temperature, as well as maximum alloying of Zn1-xCoxO, without the formation of any secondary phase. Further increase (15 and 20%) of Co in ZnO, Raman studies clearly indicate that there was signature of the formation of ZnCo2O4, which is anti ferromagnetic in nature, moreover, Co-Co nearest neighboring alignment is also antiferromagnetic, and therefore it reduces the ferromagnetic properties for concentration over 10%. The origin of high-Tc ferromagnetism in ZCO films is due to the formation of BMP between the TM local magnetic moment and donor bound electrons through the CT transition, and the sufficient hybridization of the magnetic polarons. The reduction of saturation magnetization (Ms) in Al doped ZCO samples may be due to the destabilization of magnetic polarons by the free carriers. The optical transmission spectra and XPS analysis confirmed the Co2+ substitution at Zn site and no trace of metallic Co was found in our thin films up to the instruments detection limit.
83
3.4.4 References [1]
S. J. Pearton, C. R. Abernathy, M. E. Overberg, G. T. Thaler, and D. P. Norton N. Theodoropoulou and A. F. Hebard, Y. D. Park, F. Ren and J. Kim L. A. Boatner, J. Appl. Phys. 93, 1 (2003)
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Y. Z. Yoo, T. Fukumura, Z. Jin, K. Hasegawa, M. Kawasaki, P. Ahmet, T. Chikyow and H. Kainuma, J. Appl. Phys. 90, 90446 (2001)
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M. H. Kane, K. Shalini, C. J. Summers, R. Varatharajan, J. Nause, C. R. Vestal, Z. J. Zhang and I. T. Ferguson, J. Appl. Phys. 97, 023906 (2005)
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T. Dietl, H.Ohno, F.Matsukura, J. Cibert and D. Ferrand, Science 287, 1019 (2000)
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D. P. Norton, M. E. Overberg, S. J. Pearton, J. D. Budai, L. A. Boatner, F. M. Chisholm, J. S. Lee, Z. G. Kim, Y. D. Park, and R. G. Wilson, Appl. Phys. Lett. 83, 5488 (2003)
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Y. Belghazi, G. Schmerber, S. Colis, J. L. Rehspringer, A. Dinia, A. Berrada, Appl. Phys. Lett. 89, 122504 (2006)
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H. H. Huang, C. A. Yang, P. H. Huang, C. H. Lai, T. S. Chin, H. E. Huang, H. Y. Bor, and R. T. Huang, J. Appl. Phys. 101, 09H116 (2007)
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Y. Mastumoto, M. Murakami, T. Shono, T. Hasegawa, T. Fukumura, M. Kawasaki, P. Ahmet, T. Chikyow, S. Koshihara, and H. Koinuma, Science 291, 854 (2001)
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S. Yin, M. X. Xu, L. Yang, J. F. Lui, H. Rosner, H. Hahn, H. Gleiter, D. Schild, S. Doyle, T. Liu, T. D. Hu, E. T. Muromachi, and J. Z. Jiang, Phys. Rev. B 73, 224408 (2006)
[10] P. Sati, C. Deparis, C. Morhain, S. Schafer, and A. Stepanov, Phys. Rev. Lett. 98, 137204 (2007) [11] M. Venkatesan, P. Stamenov, L. S. Dorneles, R. D. Gunning, B. Bernoux, and J. M. D. Coye, Appl. Phys. Lett. 90, 242508 (2007) [12] X. H. Xu, H. J. Blythe, M. Ziese, A. J. Behan, J. R. Neal, A. Mokhtari, R. M. Ibrahim, A. M. Fox, and G. A. Gehring, New J. Phys. 8, 135 (2006) [13] J. Alaria, H. Bieber, S. Colis, G. Schmerber, and A. Dinia, Appl. Phys. Lett. 88, 112503 (2006) [14] X. C. Lui, E. W. Shi, Z. Z. Chen, H. W. Zhang, B. Xiao, and L. X. Song, Appl. Phys. Lett. 88, 252503 (2006)
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[15] H. Harima, J. Phys: Cond. Matt. 16, S5653 (2004) [16] K. R. Kittilstved, N. S. Norberg and D. R. Gamelin, Phys. Rev. Lett., 94, 147209 (2005) [17] K. Sato, H. K. Yoshida, Semiconductor Sci. Technology, 17, 367 (2002) [18] C. H. Park and D. J. Chadi, Phys. Rev. Lett., 94, 127204 (2005) [19] L. Yan, C. K. Yong and X. S. Rao, J. Appl. Phys., 96, 508 (2004) [20] K. R. Kittilstved, W. K. Liu, and D. R. Gamelin, Nat. Mat. 5, 291 (2006) [21] K. Nielsen, S. Bauer, M. Lubbe, S. T. B. Goennenwein, M. Opel, J. Simon, W. Mader, and R. Gross, Phys. Stat. Sol. (a) 203, 3581 (2006) [22] J. M. D. Coey, M. Venkatesan, and C. B. Fitzgerald, Nat. Mat. 4, 173 (2005) [23] P. Koidl, Phys. Rev. B 15, 2493 (1977) [24] K. J. Kim, Y. R. Park, Appl. Phys. Lett. 81, 1420 (2002) [25] M. Venkatesan, C. B. Fitzgerald, J. G. Lunney, and J. M. D. Coye, Phys. Rev. Lett. 93, 177206 (2004) [26] E. Burstein, Phys. Rev. 93, 632 (1954) [27] H. J. Lee, S. Y. Jeong, C. R. Cho, and C. H. Park, Appl. Phys. Lett. 81, 4020 (2002)
85
CHAPTER 4 STRUCTURAL, OPTICAL, AND MAGNETIC PROPERTIES OF Cu AND Mn DOPED ZnO THIN FILMS This chapter is devoted to the synthesis of PLD grown Cu and Mn doped ZnO thin films and the characterization of their structural, optical, and magnetic properties. The extensive structural and lattice dynamical properties were investigated by XRD, Raman scattering studies, and high-resolution TEM analysis. In Zn1-xCuxO (x = 0.010.05) thin films, Cu goes to Zn lattice site up to a maximum of 3 stoichiometric percentage, and beyond that it precipitates as Cu related secondary phase, e.g. CuO and Cu2O in ZnO lattice. The low temperature photoluminescence showed the basic ZnO excitonic structure and the splitting of the free excitons (A and B) due to Cu doping. The room temperature Vibrating Sample Magnetometer (VSM) measurements on Zn1-xCuxO thin films showed ferromagnetism with maximum saturation magnetization of 0.76 μB/Cu in 3% Cu doped ZnO thin films. Upon further increasing of Cu concentration, the saturation magnetization decreased abruptly, probably due to the precipitation of antiferromagnetic CuO and Cu2O phases. In the case of Zn1-xMnxO (x = 0.01-0.1) thin films, we have investigated multiphonon Raman scattering in the disordered lattice due to Mn doping. The optical band gap was found to increase in considerable amount. No ferromagnetism was detected in Mn doped samples.
86
4.1
Micro-structural and Ferromagnetic Properties of Cu doped ZnO Thin Films Thin films of Cu substituted ZnO were grown on Al2O3 (0001) substrate by
pulsed laser deposition. The micro-structural properties of Zn1-xCuxO (x = 1, 3, and 5%) thin films were carried out by XRD, high-resolution TEM, and Raman scattering studies. The 1% and 3% Cu doped samples showed nearly single crystalline feature, free from any secondary phase or defects. The XPS measurements confirmed the predominant Cu2+ oxidation state of these samples. The room temperature ferromagnetism was observed with maximum saturation magnetization of 0.76 μB/Cu in 3% Cu doped sample, which decreases as the Cu concentration increases.
4.1.1 Introduction Recent reports about the observation of room temperature ferromagnetism in copper doped ZnO have been taken with great interest by the scientific community. This is mostly because of the fact that the metallic copper (Cu), as well as all possible Cubased secondary phases, are non-ferromagnetic [1]. So, if any ferromagnetism is observed in a Cu-based system, then it will undoubtedly be the intrinsic property of the material. Despite the above interest, the Cu doped ZnO system is not beyond controversies. There are several contradicting reports, where some authors have confirmed [2, 3] the occurrence of FM in this system while others have ruled it out [4]. Even in studies, where room temperature ferromagnetism is reported, the effect of carrier type on the ferromagnetic properties is unclear [5]. Buchholz et al. found that p-type carriers are essential for realizing ferromagnetism in ZnO:Cu system and nonferromagnetism in n-type system [5]. In sharp contrast to this, Hou et al. reported
87
ferromagnetism in n-type ZnCuO films [6]. Moreover, it is difficult to synthesize high quality samples with controlled dopant concentrations because of the poor solubility of Cu in ZnO [7]. In the case of Cu doped ZnO, the Cu 3d states are strongly localized because of their nearly filled (3d states) character; and hence the empty Cu 3d states are located in the gap region without hybridizing with Zn 4s conduction band. This is the reason for the low solubility of Cu in ZnO lattice.
4.1.2 Experiments High quality Zn1-xCuxO (x = 0, 0.01, 0.03, and 0.05) thin films were grown on single crystalline (0001) Al2O3 substrate by pulsed laser deposition (PLD) technique using an Excimer laser (KrF, 248nm, 10Hz) with laser energy of 2.5 J/cm2. The ceramic targets of Zn1-xCuxO for PLD were prepared by solid state reaction mechanism. The deposition chamber was initially evacuated to 1 x 10-6 Torr base pressure. The substrate temperature was maintained at 650°C. The deposition was done under the oxygen pressure of 2 mTorr for 17 min, yielding films of about 0.8 μm thickness. The XRD measurements were carried out by Siemens D5000 X-ray difractometer (XRD) with CuKα (1.54Å) radiation. The Raman scattering studies were performed by using JobinYvon T64000 Triple-mate instrument. The radiation of 514.5 nm from Ar+ laser was used for the excitation to perform Raman measurements. A liquid nitrogen-cooled charge-coupled device (CCD) system was used to collect and process the data. The cross sectional HRTEM samples were prepared by mechanical process. The samples were attached (with glue) with a smooth surface, then mechanically thinned both side by using diamond papers (30 μm to 0.25 μm) and finally ion milled. The HRTEM measurements were carried out by using FEI Tecnai F20 system operated at 200 keV. The M-H loops 88
were obtained by using vibrating sample magnetometer (VSM) system with a maximum field of 6 kOe.
4.1.3 Results and discussions 4.1.3.1
XRD Analysis
The XRD spectra of PLD grown Zn1-xCuxO (x = 0, 0.01, 0.03, and 0.05) thin films on Al2O3 (0001) substrates are shown in Fig. 4.1. The films are highly c-axis (0002) oriented with the same wurtzite structure as in pure ZnO. The (0002) peak position in 2θ of ZnO at 34.54° sifted slightly (0.05° – 0.09°) leftward with increase of Cu concentration, suggest the decrease of lattice parameter. The ionic radii of divalent Cu in tetrahedral coordination (0.57 Å) is smaller than the divalent Zn in tetrahedral coordination (0.60 Å), as a result, there are changes in the cell parameters with Cu2+ substitution in the Zn lattice site [2]. At 5% Cu concentration the films showed additional peak due to the segregation of Cu related secondary phases of CuO and Cu2O.
4.1.3.2
Raman Scattering Analysis
The first order Raman spectra of Zn1-xCuxO thin films are shown in Fig. 4.2. Besides the strong Al2O3 substrate peaks, the two characteristic optical modes of wurtzite ZnO at 98.5 and 439 cm-1 corresponding to the E2low and E2high modes were observed [8]. According to the Raman selection rules only E2 and A1 (LO) modes can be observed in the back scattering geometrics of highly c-axis oriented thin films.
89
(0002)
(0004)
S
CuO
Intensity (arb. units)
Zn1-xCuxO Thin films
Cu2O
5%Cu
*
3%Cu
1%Cu
ZnO
30
35
40
45
50
55
60
65
70
75
80
Two theeta (degree) Figure 4.1: XRD spectrum of Zn1-xCuxO thin films, CuO and Cu2O related peaks in 5% Cu doped sample at 38.38 and 61.32, respectively.
low
S
Intensity (abr. units)
E2
Zn1-xCuxO Thin films high
E2
Ag M *
S
5%Cu
3%Cu
1%Cu
ZnO
100
200
300
400
500
600
700
-1
Raman shift (cm ) Figure 4.2: Raman spectra of Zn1-xCuxO thin films on Al2O3 substrate, the CuO related Ag mode at 296.8 cm-1 was observed in Zn0.95Cu0.05O thin film.
90
The absence of low intensity A1 (LO) signal in our spectra around 574 cm-1 may be due to the dominating Al2O3 signal at 576 cm-1. The Oxygen sub-lattice vibrational optical mode (E2high) was found to shift (2.6 cm-1) towards the lower frequency compare to the ZnO thin film. This shift is due to the confinement of optical phonons in a finite region. When Cu substituted the Zn lattice site in ZnO, it forms the ternary alloy of Zn1-xCuxO and the allowed region for the optical phonon becomes finite, compared to pure ZnO. The structural disorder due to atomic substitution breaks the translational symmetry of the allowed phonons of the host lattice and leads to k ≠ 0 phonons in the Raman line shape. The disorder induced effect causes phonon lineshapes becomes broadened and shifted towards the lower frequency giving by the relation [9, 10], ω(k) = A + B cos (πk), where A and B are the constants and k = k’ ± 2nπ/L; L is the correlation length and n is the reflective index of the material. The estimated L values corresponding to 1, 3, and 5% Cu doped ZnO are 24.5, 22.3, and 20.2 nm, respectively. The red shift of the E2high mode in Zn1-xCuxO thin films confirms the substitution of Cu2+ in Zn2+ lattice site [11]. The multi phonon optical mode (M) at 331.7 cm-1 in Zn1-xCuxO thin films is due to the braking of translational symmetry caused by the Cu doping in ZnO. In the 5% Cu doped sample a new peak appears at 296.8 cm-1 corresponding to the Ag mode related to CuO [7].
4.1.3.3
High Resolution TEM Analysis
The cross sectional HRTEM images are shown in Fig. 4.3. The interface of Al2O3/Zn0.99Cu0.01O clearly showed the epitaxial nature of the grown films [Fig. 4.3 (a)]. The fast Fourier transformation (FFT) of the selected area of the interface confirms the nearly perfect epitaxial growth of the film on the substrate (inset). The interface is
91
compositionally sharp and without any interfacial reaction and inter-diffusion. Fig. 4.3 (b) showed the HRTEM micrograph of the film, which is single crystalline, defects free, and the absence of any secondary phase or precipitation. The TEM analysis suggests that the Cu is directly substituted (partially) at the Zn lattice site in ZnO host lattice. The lattice fringes are around 0.52 nm and the FFT performed on the micrograph (inset) shows spots corresponding to the film only. However, the HRTEM micrograph of 5% Cu doped ZnO thin film (Fig. 4.3 (c)) showed polycrystalline nature, the grain boundaries are indicated by dotted line in the figure. Each crystalline grain has ordered (0002) orientations. Huge lattice dislocation was observed in this thin film (dotted circle). This dislocation may be due to the segregation of CuO nano phase in ZnO lattice [7]. This lattice image exhibits stacking fault in the atomic row along (0001) planes [Fig. 4.3 (d)].
4.1.3.4
X-ray Photoelectron Spectroscopy Analysis
The X-ray photoelectron spectroscopy (XPS) measurements were carried out to investigate the bonding nature and oxidation state of Cu in Zn1-xCuxO thin films (Fig. 4.4). The Cu 2p3/2 and 2p1/2 peaks observed at 933.21 and 953.04 eV, respectively in all the samples and the appearance of Cu2p satellite (~ 943 eV) suggest that the Cu ions are predominantly in Cu2+ (d9) oxidation state [2, 3]. Apart from the 933.2 eV peak in 5% Cu doped sample, there is a broad peak at the lower energy side of 2p3/2 position. After deconvolution (inset), the peak position was found at 932.17 eV, which belongs to the Cu1+ oxidation state of Cu in the thin film.
92
(a)
Zn0.99Cu0.01O
(b)
Al2O3
0.52 nm ¯
(c)
(d)
Zn0.95Cu0.05O
Figure 4.3: (a) High resolution TEM image showing the interface between Zn0.99Cu0.01O and substrate; (b) HRTEM micrograph of Zn0.99Cu0.01O film; (c) HRTEM image of polycrystalline Zn0.95Cu0.05O thin film; and (d) Lattice dislocation [dotted circle in 3(c)] exhibit stacking fault in the atomic row along (0001) planes. All insets show the FFT of the corresponding micrograph.
93
933.21
2p3/2
As Observed 2+ Cu 1+ Cu
Intensity (abr. units)
932.17
928
930
932
2p1/2 934
936
938
5% Cu
Binding energy (eV)
3% Cu
1% Cu
925
930
935
940
945
950
955
960
965
Binding Energy (eV)
Figure 4.4: XPS spectrum of Cu2p peaks of Zn1-xCuxO thin films, inset shows the deconvolution of Cu 2p3/2 of Zn0.95Cu0.05O thin film.
94
4.1.3.5
Room Temperature PL Analysis of Zn1-xCuxO Thin Films
The room temperature PL measurements of Zn1-xCuxO thin films are shown in Fig. 4.5. The narrowing of the near band edge (NBE) transition was found in room temperature photoluminescence (RTPL) measurements. The emission peak at 3.29 eV (UV) originates from the near-band-edge (NBE) transition in band gap of ZnO due to the recombination of free excitons through an exciton-exciton collision process. The RTPL spectrum shows that the UV emission is redshifted with increase of Cu concentration. This narrowing of optical band gap in Cu doped ZnO samples is due to the sp-d exchange interaction between the d electrons of transition metal and the band electrons of ZnO; the strength of this interaction strongly depends on the number of d electrons [12]. The s-d and p-d exchange give rise to negative and positive corrections to the conduction and valance band edges respectively, leading to the NBE narrowing [13]. The green emission (~ 2.60 eV) was observed in Cu doped (>1%) ZnO samples; surface defects, Cu impurities, and the oxygen vacancies are considered to be responsible for the green emissions in Zn1-xCuxO thin films [14]. The substitution of Cu in Zn lattice side donates two electrons for bond formation and become a neutral state of Cu2+ (3d9), with a deep acceptor level of 0.17-0.19 eV bellow the bottom of the conduction band [15, 16]. This Cu impurity center behaves like a trap for nonequilibrium holes or electrons. As a carrier is trapped, the Cu impurity center acquires a charge with respect to the lattice, and a longrange Coulomb field arises in addition to the short-range potential of the isoelectronic impurity. This field promotes electron or hole trapping onto a weakly localized hydrogen like orbital.
95
λ = 325 nm
FX
PL Intensity (arb. units)
@ 300 K
5% Cu 3% Cu 1% Cu ZnO
2.25
2.50
2.75
3.00
3.25
3.50
Photon Energy (eV)
Figure 4.5: PL spectra of Zn1-xCuxO thin film at 300 K; NBE transition red shifted due to Cu doping.
96
The valence electron is transferred to the d-orbital of the Cu2+ impurity, and the electronic subsystem of the ZnO lattice is excited, while the hole is localized at the Cu+ ion; consequently an exciton like state (3d10Cu+, h) is formed at the impurity center. The radiative recombination of this exciton gives rise to a leading of green emission in the spectrum [17- 19].
4.1.3.6
Magnetic Properties of Zn1-zCuxO Thin Films
Room temperature magnetic measurements on Zn1-xCuxO thin films were carried out using vibrating sample magnetometer (VSM). The M-H loops are shown in Fig. 4.6. All the films showed ferromagnetic nature at room temperature. The maximum saturation magnetization of 0.76 μB/Cu was observed in 3% Cu doped ZnO sample. The saturation magnetization decreased in further increase of Cu concentration. The coercive fields (Hc) of 1, 3, and 5% Cu doped samples are 171, 185, and 145 Oe respectively. The decrease of saturation magnetization in higher Cu concentration could be due to the increased number of Cu atom occupying adjacent cation position resulting the anti ferromagnetic alignment [20, 21] and also, contributions from anti ferromagnetic Cu2O [1] secondary phase. In order to understand the observed ferromagnetic moment of the films, it is essential to have an insight into the possible electronic configurations of Cu ions in the material. Metallic Cu (Co0) atoms have an outer cell electronic configuration of 3d104s1 and, hence, Cu+ and Cu2+ ions posses 3d10 and 3d9 configurations, respectively. In 3d10 configuration, all the d electrons are paired and, hence Cu+ ions do not posses any magnetic moment. On the other hand, in the case of Cu2+ ions with d9 configuration, one unpaired electron is available.
97
1.00
T = 300 K 0.75
Zn0.99Cu0.01O Zn0.97Cu0.03O
M (μΒ/Cu)
0.50
Zn0.95Cu0.05O
0.25 0.00 -0.25 -0.50 -0.75 -1.00 -6
-4
-2
0
2
4
6
Field (kOe)
Figure 4.6: M-H loop of the Zn1-xCuxO thin films at room temperature, maximum saturation magnetization (Ms) of 0.76 μB/Cu was observed in 3% Cu doped ZnO sample.
EC
~ 0.17eV
dxy
t2
dx2 −y2
Cu2+ (d9 ) ~ 3.37 eV
e
dxz
dyz
dz2 EV Figure 4.7: Schematic diagram and 3d orbitals splitting of Cu2+ (d9) in Zn2+ side of ZnO lattice under tetrahedral crystal field and spin-orbit interaction
98
This unpaired electron will give rise to a spin angular momentum of 1/2 which can result of the total magnetic moment (M) ~ 1.00 μ B corresponding to the Cu2+ (d9) valence state, as determined by local spin density approximation (LSDA) based calculations [21]. The observed ferromagnetism (0.76 μ B /Cu) is compare to the theoretical value, implying that Cu2+ ions in our Zn0.97Cu0.03O films are predominant (about 76%) in magnetically active Cu2+ state. Our results confirm that the partially substitution of Cu2+ ions in ZnO lattice can induce ferromagnetic ordering. According to the MCD measurements, Ando et al. [12] showed that the strength of the sp-d exchange interaction in ZnO-based DMS materials seems to depend on the number of 3d electrons. When the Cu2+ (3d9) ions substituted in the tetrahedral side in ZnO lattice, the energy band of the d-orbital splitted to t2 and e due to the tetrahedral crystal field (Td) and spin-orbit interaction. Moreover, the distortion of Td-field caused by substitution of Cu2+ in Zn2+ lattice site again splitted t2 and e as shown in Fig. 4.7. All of the spin-up states of the t2-orbitals are occupied. The up-spin p-electrons of the surrounding oxygen atoms cannot couple with the transition metal ion, whereas the down-spin p-electrons can do that. The exchange interaction through p-d orbital mixing may explain the ferromagnetic exchange interaction in ZnO:Cu thin films.
4.1.4 Summary In summary, thin films of Cu doped ZnO were grown on Al2O3 substrate by pulsed laser deposition. The XRD measurements showed the films are highly c-axis oriented and free from any secondary phase up to 3% of Cu substitution. Further increase of Cu concentration (5%), CuO and Cu2O related phases were detected. Raman scattering
99
studies confirms the substitution of Cu2+ in Zn2+ lattice side up to 3% and further increase of Cu in ZnO precipitates CuO secondary phase. The HRTEM provided evidence of nearly single crystalline, defects free, and epitaxial thin films up to 3% of Cu substitution. The XPS measurements indicate that the Cu ions are predominantly in Cu2+ oxidation state. Room temperature PL measurements confirmed the decrease of NBE transition in Zn1-xCuxO thin films. The strong sp-d exchange interaction in Zn1-xCuxO thin films results the NBE narrowing. The observed maximum saturation magnetization at room temperature is comparable with the theoretical value of the tetrahedrally coordinated Cu2+ in ZnO lattice. Among the total Cu atoms, ~ 76% of Cu ions are magnetically ordered in the system. The reduction of saturation magnetization (Ms) in 5% Cu doped sample may be due to the precipitation of anti-ferromagnetic CuO and Cu2O phase. The p-d exchange interaction could explain the origin of observed global magnetization in ZnO:Cu thin films.
100
4.1.5 References [1]
M. Wei, N. Braddon, D. Zhi, P. A. Midgley, S. K. Chen, M. G. Blamire, and J. L. MacManus-Driscoll, Appl. Phys. Lett. 86, 72514 (2005)
[2]
D. Chakraborti, J. Narayan, and J. T. Prater, Appl. Phys. Lett. 90, 062504 (2007)
[3]
T. S. Herng, S. P. Lau, S. F. Yu, H. Y. Yang, and X. H. Ji, J. S. Chen, N. Yasui and H. Inaba, J. Appl. Phys. 99, 086101 (2006)
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[5]
D. B. Buchholz, R. P. H. Chang, J. H. Song, and J. B. Ketterson, Appl. Phys. Lett. 87, 082504 (2005)
[6]
D. L. Hou, X. J. Ye, H. J. Meng, H. J. Zhou, X. L. Li, C. M. Zhen, and G. D. Tang, Appl. Phys. Lett. 90, 142502 (2007)
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C. Sudakar, J. S. Thakur, G. Lawes, R. Naik, and V. M. Naik, Phys. Rev. B 75, 054423 (2007)
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T. C. Damen, S. P. S. Porto and B. Tell, Phys. Rev. 142, 570 (1966)
[9]
P. Parayanthal and F. H. Pollak, Phys. Rev. Lett. 52, 1822 (1984)
[10]
H. Richter, Z. P. Wang and L. Ley, Solid State Commu. 39, 625 (1981)
[11]
K. Samanta, P. Bhattacharya, R. S. Katiyar, W. Iwamoto, P. G. Pagliuso, and C. Rettori, Phys. Rev. B 73, 245213 (2006)
[12]
K. Ando, H. Saito, Z. Jin, T. Fukumura, M. Kawasaki, Y. Matsumoto, and H. Koinuma, J. Appl. Phys. 89, 7284 (2001)
[13]
J. K. Furdyna, J. Appl. Phys. 64, R29 (1988)
[14]
H. Zhu, J. Iqbal, H. Xu, and D. Yu, J. Chem. Phys. 129, 124713 (2008)
[15]
Y. R. Lee, A. K. Ramdas, and R. L. Aggarwal, Phys. Rev. B 38, 10600 (1988)
[16]
J. K. Furdyna, J. Appl. Phys. 64, R29 (1988)
[17]
R. Dingle, Phys. Rev. Lett. 23, 579 (1969)
[18]
Ya. I. Alivov, M. V. Chukichev, and V. A. Nikitenko, Semiconductors 38, 31 (2004)
101
[19]
F. H. Su, J. Appl. Phys. 100, 013107 (2006)
[20]
K. Sato and H. Katayama-Yoshida, Jpn. J. Appl. Phys., Part 2 39, L555 (2000)
[21]
M. S. Park and B. I. Min, Phys. Rev. B 68, 224436 (2003)
102
4.2 Optical Properties of Zn1-xCuxO Thin Films Investigation of optical properties of PLD grown Cu doped (1, 3, and 5%) ZnO thin films were carried out by UV-VIS and PL measurements. XRD, HRTEM, and Raman scattering analysis confirms the substitution of Cu2+ in Zn2+ site of ZnO structure up to 3%, further increase of Cu concentration secondary phases of CuO, and Cu2O were appear. Low temperature PL measurement confirms the basic excitonic transitions of ZnO with additional transition at 3.376 eV due to Cu doping. Temperature dependent PL spectrum confirmed the existence of donor acceptor pair (DAP) band at 3.307 eV in Zn1xCuxO
(x = 0, 0.01, 0.03, and 0.05) thin films. The band edge transition was found to be
decreases in Cu doped samples.
4.2.1 Introduction The II-VI compound semiconductor ZnO with high optical band gap (3.37 eV) at 77K and high excitonic binding energy (60 meV) [1] remains the excitons stable even at room temperature. The p-type conduction in ZnO with hole concentration around 1019 cm-3 [2, 3] were successfully demonstrated by means of doping with group-V acceptors like N, P, and As. When ZnO doped with Al and Ga donor impurities, it becomes as a low-resistive transparent contact with high radiation, chemical, and thermal resistance, which is useful for the devices as photodetectors, Schottky-diodes, and sensors. Thus, the high hole and electron conduction, combined with a wide direct band gap makes ZnO an essential candidate for the development of semiconductor laser and emitters of visible and UV light. Gao et al. attempted to make the light emitting diode based on ZnO doped with donor and acceptor impurities but the diodes didn’t exhibit luminescence [4]. Moreover, it is necessary to pick up the suitable dopant to ensure that it would leave 103
emission spectrum and transparency of ZnO intact. The group-I element [Cu] is a very good dopant in ZnO because it simultaneously comprises of an electron donor state of Cu2+ (2T2) and corresponding acceptor Cu+ (3d10) level within the forbidden energy region in ZnO [5]. Therefore Cu acts as luminescence activators as well as a compensator of n-type materials are of considerable significance for II-VI compound semiconductor in general [5]. The emission spectra can be extended from UV to IR region depending on the concentration of Cu, defects in ZnO, and the excitation conditions [6, 7]. In order to design ZnO-based optoelectronic devices structure, the most important step is to realize the bandgap engineering to create the barrier layers and the quantum wells in the device heterostructures. We have analyzed the doping effect on the distinct photoluminescence property of nanocrystalline Zn1-xCuxO thin films.
4.2.2 Result and discussion 4.2.2.1
Optical Transmission Spectra Analysis
The optical transmission spectra of Zn1-xCuxO thin films in the range of 330 to 800 nm are shown in Fig. 4.8. The band edge transition of the films decreases from 3.276 eV to 3.263 eV with increase in Cu concentration up to 5% (inset). This reduction of optical band gap in Zn1-xCuxO thin films at room temperature may attribute to the sp-d exchange interaction between the band electrons and the localized d electron of the Cu2+ ions substituting Zn2+ ions [8]. The Cu substitution in ZnO host lattice creates impurity level in the forbidden energy region of ZnO. The strong sp-d exchange interaction was observed in the MCD measurements on PLD grown Zn1-xCuxO thin films [9].
104
90 ZnO 1%Cu 3%Cu 5%Cu
60 1.5
45 30 15
(α2) (x10 9 cm -2)
Transmission (%)
75
1.2
ZnO Cu 1% Cu 3% Cu 5%
0.9
0.6
0.3
0.0 3.20
3.22
0 1.6
2.0
3.24
3.26
3.28
hν (eV)
3.30
2.4
3.32
3.34
2.8
3.2
3.6
Photon energy (eV)
Figure 4.8: Optical transmission spectrum of Zn1-xCuxO thin films, the reduction of band edge transition with Cu content (inset)
105
The s-d and p-d exchange give rise to negative and positive corrections to the conduction and the valance band edges respectively, leading to the band gap narrowing [8].
4.2.2.2
Photoluminescence Analysis
The photoluminescence spectra of Zn1-xCuxO thin films at 77 K are shown in Fig. 4.9. The excitonic emissions at 3.366, 3.353, and 3.307 eV were observed corresponding to the free exciton (FX), donor bound exciton (D0X), and donor acceptor pair (DAP) recombination, respectively. The first and second LO phonon replica of the DAP band were observed at 3.235 and 3.163 eV, respectively. Apart from these transitions, a new peak at 3.367 eV appears due to Cu doping. This peak belongs to the B-free exciton (FXB), and this is due to the transition from conduction band to the B valance band [10]. The separation of between FXA and FXB is ~ 10 meV, which is within the reported value (9-15 meV) [1, 11]. Temperature dependent PL spectrum of 5% Cu doped ZnO thin film is shown in Fig. 4.10. Due to the increase of temperature, the intensity of free exciton (FX) at 3.366 eV increases and dominated over the bound exciton (D0X) because of the ionization of the impurities, that used to bound exciton at low temperature [12]. The DAP transition at 3.307 eV showed a different behavior with temperature, this peak blue shifted with the increase of temperature and a new peak emerges at the higher energy side. This feature is typically for DAP transition, and the new peak at higher energy is caused by band-toimpurity transition [13, 14]. The impurity involved in this transition must have high binding energy; keeping in mind, the higher binding energy of the acceptor in ZnO.
106
0
DX FXA
FXA = 3.367 eV
0
DX
FXB = 3.377 eV
FXA
PL Intensity (a. u)
FXB
77 K
FXB DAP
3.34
3.36
3.38
1LODAP
3.40
Photon energy (eV)
2LODAP
5% Cu
3% Cu 1% Cu ZnO 3.0
3.1
3.2
3.3
3.4
Photon energy (eV)
Figure 4.9: Photoluminescence spectrum of Zn1-xCuxO thin films at 77 K; inset shows the de-convolution of FXA and FXB transition.
0 Zn0.95Cu0.05O D X
DAP 0
PL Intensity (abr. units)
eA
λ = 325 nm
FX
* DAP 3.28
3.29
3.30
3.31
3.32
3.33
0
eA
PL Energy (eV)
77 K 100 K 120 K 140 K 160 K 180 K
1LODAP 2LODAP
3.10
3.15
3.20
3.25
3.30
3.35
3.40
3.45
Photon energy (eV)
Figure 4.10: Temperature dependent PL spectrum of Zn0.95Cu0.05O thin film, the appearance of eA0 at the higher energy side of DAP band with elevated temperature (inset)
107
This band-to-impurity transition is attributed to the conduction band-to-acceptor transition (eA0), in which the free electrons in the conduction band recombine with acceptors.
The
transition
energy
( E eA )
of
eA0
can
be
express
as
[15],
E eA = E g − E A + k B T / 2 , where Eg is the band gap energy, EA is the binding of the acceptor. The band gap energy at a particular temperature can be estimated from the FX transition energy and the free excitonic binding energy (60 meV). The acceptor binding energy at 160 K has been calculated as 105 meV, using EFX = 3.353 eV, EeA = 3.313 eV. This acceptor binding energy from the PL measurement is in good agreement with the value of 107 meV, reported earlier by Zhang et al. [12].
4.2.3 Summary We have observed a well-resolved PL spectrum of PLD grown Cu doped ZnO thin films on Al2O3 substrates. The free excitonic fine structure (FXA and FXB) was observed in Cu doped samples. With the increase of temperature the appearance of eA0 emission at the higher energy side of the peak at 3.307 eV confirmed the characteristics of DAP transition rather than accepter bound exciton. The near band-edge (NBE) transition was found to decreases with increase in Cu concentration; this reduction is due to the sp-d exchange interaction between the d electrons of Cu electron of ZnO lattice.
108
2+
state and the band
4.2.4 References [1]
D. G. Thomas, J. Phys. Chem. Solids 15, 86 (1960)
[2]
T. Yamamoto and H. Katayama-Yoshida, Jpn. J. Appl. Phys. 38, L166 (1999)
[3]
M. Joseph, H. Tabata, and T. Kawai, Jpn. J. Appl. Phys. 38, L1205 (1999)
[4]
X. Guo, H. Choi, H. Tabata, and T. Kawai, Jpn. J. Appl. Phys. 40, L177 (2001)
[5]
C. X. Xu, X. W. Sun, X. H. Zhang, L. Ke, and S. J. Chua, Nanotechnology 15, 856 (2004)
[6]
R. Dingle, Phys. Rev. Lett. 23, 579 (1969)
[7]
P. Dahan and V. Fleurov, Phys. Rev. B 57, 9690 (1998)
[8]
K. J. Kim and Y. R. Park, Appl. Phys. Lett. 81, 1420 (2001)
[9]
K. Ando, H. Saito, Z. Jin, T. Fukumura, M. Kawasaki, Y. Matsumoto, and H. Koinuma, J. Appl. Phys. 89, 7284 (2001)
[10]
A. Teke, Ü. Özgür, S. Doğan, X. Gu, H. Morkoç, B. Nemeth, J. Nause, and H.O. Everitt, Phys. Rev. B 70, 195207 (2004)
[11]
D. C. Reynolds, D. C. Look, B. Jogai, C. W. Litton, G. Cantwell, and W. C. Harsch, Phys. Rev. B 60, 2340 (1999)
[12]
B. P. Zhang, N. T. Binh, Y. Segawa, K. wakatsuki, and N. Usami, Appl. Phys. Lett. 83, 1635 (2003)
[13]
D. J. As, F. Schmilgus, C. Wang, B. Schottker, D. Schikora, and K. Lischka, Appl. Phys. Lett. 70, 1311 (1997)
[14]
J. F. Wang, D. Masugata, C. B. Oh, A. Omino, S. Seto, and M. Isshiki, Phys. Status Solidi A 193, 251 (2002)
[15]
H. Richter, Z. P. Wang and L. Ley, Solid State Commu. 39, 625 (1981)
109
4.3
Multi-phonon Raman Scattering in Mn doped ZnO The multi-phonon Raman scattering in Mn-doped (1-10%) ZnO was observed at
room temperature using 514.5 nm Ar+ laser as excitation source. The additional optical modes at 327, 332, 482, 532, and 680 cm-1 besides the ZnO related first order modes were identified as the second order Raman modes in the disordered lattice; however, the phonon mode at 327 and 680 cm-1 originated from the precipitation of the spinal secondary phase ZnMn2O4. Using the phonon confinement model in the Zn1-xMnxO alloy, we have calculated the crystalline grain size of 1, 3, 5, and 10% Mn doped samples as 31.8, 18.3, 15.9, and 14.1 nm, respectively. The room temperature optical transmission spectra in the spectral range from 300 nm to 800 nm shows a considerable increase of optical band gap (3.27 eV to 3.41 eV) due to the Mn doping in ZnO. The XRD spectrum of the Zn1-xMnxO ceramic targets and thin films of less than 10% Mn doped samples showed that the compounds are single phased and highly c-axis oriented with an increase of c-axis lattice parameter.
4.3.1 Result and Discussion 4.3.1.1
XRD Analysis
The structural analysis of Mn doped ZnO compound was carried out by XRD measurements. Figure 4.11 and Fig. 4.12 showed the XRD spectrum of Zn1-xMnxO (x = 1, 3, 5, and 10%) ceramic targets and thin films on (0001) oriented Al2O3 substrate. All of the peaks related to wurtzite ZnO structure were identified in Fig. 4.11 without any secondary phase up to the detection limit of the instrument.
110
(100) (200) (112)
(103)
(102)
(110)
(101)
(100) (002)
Intensity (arb. units)
Mn 10% Mn 5% Mn 3% Mn 1% ZnO
30
35
40
45
50
55
60
65
70
Two theeta (degree)
Intensity (arb. units)
(0002)
Figure 4.11: XRD spectrum of Zn1-xMnxO ceramic targets
Al2O3 10% Mn
5% Mn 3% Mn 1% Mn ZnO 30
35
40
45
50
55
Two theeta (degree) Figure 4.12: XRD spectrum of Zn1-xMnxO thin films on (0001) Al2O3 substrates
111
The XRD pattern of the thin films samples (Fig. 4.12) shows that the ZnO and Zn1xMnxO
are highly c-axis oriented on alumina substrate. It was found that due to the Mn
doping in ZnO, the (0002) peak shifted gradually towards lower diffracting angle as compared to ZnO. This phenomenon is understandable keeping in mind that the ionic radius of tetrahedrally coordinated Mn2+ (0.66 Å) is larger than Zn2+ (0.6 Å) [1].
4.3.1.2
Raman Scattering Analysis
The first order Raman scattering modes for ZnO and Zn1-xMnxO ceramic targets (used for thin film deposition) are shown in Fig. 4.13. We observed five normal modes at 99.9, 381, 438.5, 573, and 584 cm-1 corresponding to E2low, A1TO, E2high, A1LO, and E1LO, respectively [2], with the lower frequency shift of E2low mode (inset). Apart from the normal modes of vibration of ZnO, we observed five addition modes at 327 (I1), 332 (I2), 482 (I3), 525 (I4), and 680 (I5) cm-1; where I1 and I5 modes are present only in 10% Mn doped ZnO sample. The modes I2, I3, and I4 are assigned as multi-phonon scattering considering the two phonon process in the disorder lattice due to Mn doping. The modes are identified as (E2high – E2low), (A1TO + E2low), and (E2high + E2low), respectively. To identify the additional modes I1 and I5, we carried out Raman measurement of ZnMn2O4 target and compare with 10% Mn doped ZnO; the Fig. 4.14 clearly shows that these two additional modes originated from the precipitation of spinal secondary phase ZnMn2O4 in 10% Mn doped ZnO. The substitution of Mn2+ in the Zn2+ site of the ZnO has a direct effect on the optical (E2Low) mode.
112
Intensity (arb. units)
I4
Lo
A1 Lo E1
I5 80
90
E2L
100
110
120
-1
Raman shift (cm )
I1
E2H I3
10% Mn
I2 ATO 1 5% Mn 3% Mn 1% Mn ZnO
150
300
450
600
750
Raman shift (cm-1)
900
Intensity (arb. units)
Figure 4.13: Raman spectra of Zn1-xMnxO ceramic targets at room temperature, inset shows the shift of E2low mode towards lower frequency side
-1 Ia = 178 cm -1 Ib = 327 cm -1 Ic = 389 cm -1 Id = 680 cm Id
Ib
Mn 10%
Ic
Ia
ZnMn2O4
150
300
450
600
750
900
-1
Raman shift (cm ) Figure 4.14: Identification of secondary spinal phase ZnMn2O4 precipitated in 10% Mn doped ZnO
113
The structural disorder due to the atomic substitution breaks the translational symmetry of the zone centered (k = 0) optical phonons of the host lattice, which leads to the contribution of (k ≠ 0) phonons to the Raman line shape corresponding to the finite size effect. The disorder-induced effects (lower frequency shift and broadening) of E2Low mode (inset of Fig. 4.13) in Zn1-xMnxO can be explained by alloy potential fluctuations (APF) using a spatial correlation (SC) model [3, 4]. The consideration of a Gaussian attenuation factor exp(-k2L2/4), where L is the diameter of the correlation region, can be used to account for the k vector relaxation related to finite size effect [5] and structural disorder [6]. The Raman intensity at a frequency ω in a finite region can be written as [3, 5], 4πk 2 exp(−k 2 L2 / 4)dk I (ω ) ≅ ∫ 0 [ω − ω ( k )] 2 + [Γ / 2] 2 0 1
(1)
Where, k has the unit of 2π/a, ‘a’ is the lattice constant, and Γ0 is the undoped ZnO FWHM. Considering one dimensional linear chain model the dispersion relation for the E2low mode (Zn-sub lattice) of wurtzite ZnO can be written as,
ω(k) = A + B Cos (πk)
(2)
Where, A = 73.9 cm-1 and B = 26 cm-1 for the E2low phonon dispersion according to the ab initio phonon dispersion relation, calculated for ZnO [7] and k = k’ ± 2nπ/L, where n is the reflective index of the material. The estimated correlation lengths are 31.8, 18.3, 15.9, and 14.1 nm for the 1, 3, 5, and 10% Mn doped ZnO, respectively.
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4.3.1.3
Optical Transmission Spectra
The optical transmission spectra of Zn1-xMnxO thin films are shown in Fig. 4.15. The films are transparent in visible optical range and the transmittance maxima decreases upon increasing in Mn concentration. As Mn content increases, an additional broad absorption around 3 eV was observed. The optical band gap of the Zn1-xMnxO films were calculated by using the relation α 2 ∝ (hν − E g ) , where α is the absorption coefficient. The optical band gap increases with the increase of Mn content of 3% and above in the films up to 3.4 eV for 10% Mn. The band edge was decreased at Mn concentration for 1% as shown in inset of Fig. 4.13. This decrease in bandgap for low Mn concentration was also observed in Mn doped ZnSe and attributed to the sp-d interaction between transition metal and group VI anions [8]. The increase in band gap and the development of the midgap absorption by Mn doping were also repotted in ZnO [9] and other II-MnVI alloys [10]. The midgap absorption was assigned as 6A1 to 4T2 transition based on the reflectance spectra for ceramics specimens [11]. However, the spin forbidden d-d transition is unable to explain such a large absorption peak. As the absorption occurred in the wide energy range; therefore, lattice distortion due to Mn incorporation in ZnO, which might increase the oscillatory strength of d-d transition, is also ruled out. Fukumara et al. [10] assigned this absorption is due to the charge transfer transition between donor and/or acceptor ionization levels of Mn ions and the band continuum, as indicated in the calculated results of (ZnMn)S and (ZnMn)Se based on the Anderson impurity model.
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ZnO
60
3% Mn 5% Mn
1% Mn
ZnMn2O4
40
20
0
10% Mn
3.42
Optical bandgap (eV)
Transmission (%)
80
3.40 3.38 3.36 3.34 3.32 3.30 3.28 3.26 3.24
0
2
4
6
8
10
Mn concentration (%)
1.5
2.0
2.5
3.0
3.5
Photon energy (eV) Figure 4.15: Optical transmission spectra of Zn1-xMnxO and ZnMn2O4 thin films; insert shows the increase of optical bandgap energy due to Mn doping
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4.3.2 Summary We have studied in detail the structural and optical properties of Mn doped ZnO ceramic targets and thin films by XRD, Raman scattering, and optical transmission spectra. The multi-phonon peaks due to Mn doped ZnO disorder lattice and precipitation of ZnMn2O4 for higher Mn concentration (10%) were identified and assigned properly. The phonon confinement effect confirms the substitution of Mn2+ in Zn2+ site of the ZnO host lattice and the crystalline grain size decreases upon increasing Mn concentration. Optical absorption showed broad sub-bandgap absorption and increasing bandgap with increase in Mn concentration above 1%.
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R. B. Bylsma, W. M. Becker, J. Kossut, U. Debska, and D. Yoder-Short, Phys. Rev. B 33, 8207 (1986)
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