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Crystallization of Amorphous Silicon Thin Films Induced by Nanoparticle Seeds

Permanent Link: http://ufdc.ufl.edu/UFE0024344/00001

Material Information

Title: Crystallization of Amorphous Silicon Thin Films Induced by Nanoparticle Seeds
Physical Description: 1 online resource (133 p.)
Language: english
Creator: Kim, Taekon
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2009

Subjects

Subjects / Keywords: amorphous, crystallizaontion, excimer, solid
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: Crystallization of amorphous Si (a-Si) thin film has received extensive interest for their attractive applications into Si thin film transistors and Si based solar cells. Among various crystallization techniques, Solid phase crystallization (SPC) and Excimer laser crystallization (ELC) were investigated. Firstly, Solid phase crystallization (SPC) of amorphous silicon thin films deposited by the DC magnetron sputtering system with a modification in nucleation step was investigated at low temperature. The thin film consists of polycrystalline nanoparticles embedded in an amorphous matrix which can act as nuclei during crystallization, resulting in a lower thermal energy for the nucleation. The lowering energy barrier for nucleation would shorten the transition time from amorphous into polycrystalline silicon resulting from the reduction of incubation time and also lower the processing temperature spontaneously. In addition, a comprehensive study of the growth mechanism of the sputtered amorphous silicon thin films is presented during annealing. Samples were prepared with various substrate temperatures and RF power in order to optimize the crystallization of a-Si after the deposition. Also, the effects of annealing condition were examined. Low pressure N2 ambient during SPC promoted crystallization of a-Si thin films and the crystallinity. The low pressure annealing had a large impact on the crystallinity and growth behavior of subsequent films. In addition, the crystallinity, incubation time, the crystallized volume fraction and growth rate of the films annealed in a conventional furnace have been extensively studied by XRD and HRTEM. It was believed that crystalline Si nanoparticles would act as nuclei for growth of crystalline Si thin films, thus removing the high temperature requirement for nucleation, resulted in the improvement of the crystallization of a-Si. Secondly, the controlled Super Lateral Growth (SLG) can be obtained by nanoparticle induced crystallization (NIC) technique during laser annealing, which led to enhance the random super lateral growth (SLG) of Si thin films for the excimer laser crystallization (ELC). The crystallinity and surface information of the films irradiated by excimer laser have been studied by Raman spectroscopy, FESEM and AFM. Also, Transmission electron microscopy (TEM) was employed in order to obtain structural information. Polycrystalline Si nanoparticles, which have higher melting point than those of amorphous phase, would survive at high energy density of laser. In general, super lateral growth (SLG) occurs at vary narrow laser energy density region. Thus, it tends to be sensitive to laser energy density, which means not easy to control because of the characteristics of the mechanism of the SLG. In this study, poly-Si nanoparticles would act as nucleation seeds for the growth of the films during the solidification. Those nanoparticle seeds provided more probability to survive at higher density of energy compared to that without nanoparticle seeds, resulted in the large grain size distribution and the controlled super lateral growth (SLG), relatively independent of laser energy density.
General Note: In the series University of Florida Digital Collections.
General Note: Includes vita.
Bibliography: Includes bibliographical references.
Source of Description: Description based on online resource; title from PDF title page.
Source of Description: This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Statement of Responsibility: by Taekon Kim.
Thesis: Thesis (Ph.D.)--University of Florida, 2009.
Local: Adviser: Singh, Rajiv K.
Electronic Access: RESTRICTED TO UF STUDENTS, STAFF, FACULTY, AND ON-CAMPUS USE UNTIL 2009-11-30

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2009
System ID: UFE0024344:00001

Permanent Link: http://ufdc.ufl.edu/UFE0024344/00001

Material Information

Title: Crystallization of Amorphous Silicon Thin Films Induced by Nanoparticle Seeds
Physical Description: 1 online resource (133 p.)
Language: english
Creator: Kim, Taekon
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2009

Subjects

Subjects / Keywords: amorphous, crystallizaontion, excimer, solid
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: Crystallization of amorphous Si (a-Si) thin film has received extensive interest for their attractive applications into Si thin film transistors and Si based solar cells. Among various crystallization techniques, Solid phase crystallization (SPC) and Excimer laser crystallization (ELC) were investigated. Firstly, Solid phase crystallization (SPC) of amorphous silicon thin films deposited by the DC magnetron sputtering system with a modification in nucleation step was investigated at low temperature. The thin film consists of polycrystalline nanoparticles embedded in an amorphous matrix which can act as nuclei during crystallization, resulting in a lower thermal energy for the nucleation. The lowering energy barrier for nucleation would shorten the transition time from amorphous into polycrystalline silicon resulting from the reduction of incubation time and also lower the processing temperature spontaneously. In addition, a comprehensive study of the growth mechanism of the sputtered amorphous silicon thin films is presented during annealing. Samples were prepared with various substrate temperatures and RF power in order to optimize the crystallization of a-Si after the deposition. Also, the effects of annealing condition were examined. Low pressure N2 ambient during SPC promoted crystallization of a-Si thin films and the crystallinity. The low pressure annealing had a large impact on the crystallinity and growth behavior of subsequent films. In addition, the crystallinity, incubation time, the crystallized volume fraction and growth rate of the films annealed in a conventional furnace have been extensively studied by XRD and HRTEM. It was believed that crystalline Si nanoparticles would act as nuclei for growth of crystalline Si thin films, thus removing the high temperature requirement for nucleation, resulted in the improvement of the crystallization of a-Si. Secondly, the controlled Super Lateral Growth (SLG) can be obtained by nanoparticle induced crystallization (NIC) technique during laser annealing, which led to enhance the random super lateral growth (SLG) of Si thin films for the excimer laser crystallization (ELC). The crystallinity and surface information of the films irradiated by excimer laser have been studied by Raman spectroscopy, FESEM and AFM. Also, Transmission electron microscopy (TEM) was employed in order to obtain structural information. Polycrystalline Si nanoparticles, which have higher melting point than those of amorphous phase, would survive at high energy density of laser. In general, super lateral growth (SLG) occurs at vary narrow laser energy density region. Thus, it tends to be sensitive to laser energy density, which means not easy to control because of the characteristics of the mechanism of the SLG. In this study, poly-Si nanoparticles would act as nucleation seeds for the growth of the films during the solidification. Those nanoparticle seeds provided more probability to survive at higher density of energy compared to that without nanoparticle seeds, resulted in the large grain size distribution and the controlled super lateral growth (SLG), relatively independent of laser energy density.
General Note: In the series University of Florida Digital Collections.
General Note: Includes vita.
Bibliography: Includes bibliographical references.
Source of Description: Description based on online resource; title from PDF title page.
Source of Description: This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Statement of Responsibility: by Taekon Kim.
Thesis: Thesis (Ph.D.)--University of Florida, 2009.
Local: Adviser: Singh, Rajiv K.
Electronic Access: RESTRICTED TO UF STUDENTS, STAFF, FACULTY, AND ON-CAMPUS USE UNTIL 2009-11-30

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2009
System ID: UFE0024344:00001


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1 CRYSTALLIZATION OF AMORPHOUS SILICON THIN FILMS INDUCED BY NANOPARTICLE SEEDS By TAEKON KIM A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 2009

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2 2009 Taekon Kim

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3 To my parents and brother

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4 ACKNOWLEDGEMENTS I would like to acknowledge the many individuals that have made this work possible. First, I would like to thank my advisor, Dr. Rajiv K. Singh, for his guidance, support and encouragement for my research. The open environment that he has provided over the years that I have worked with him has made for a very dynamic and rewarding graduate school experience. I would also like to thank Dr. David Norton, Dr. Stephen Pearton, Dr. Brent Gila, and Dr. Fan Ren for kindly participating on my dissertation committee. I would like to also acknowledge all my friends and coworkers who made graduate school such a memorable time. Seoungyong Son, Jaeseok Lee, Myunghwan Oh, Purushotam Kumar, Sushant Gupta, Aniruddh Khanna, Junghoon Jang, Kerry Siebein and others provided the social perspective that is always necessary in life. I would also like to thank the administrative and support staff of the Department of Materials Science and Engineering for their efforts in making success possible. Finally, I would like to thank my parents, my brother and the rest of my family for the endless love and support they have provided over the many years of my extended education.

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5 TABLE OF CONTENTS page ACKNOWLEDGEMENTS .............................................................................................................4TABLE OF CONTENTS ............................................................................................................. ....5LIST OF TABLES ...........................................................................................................................7LIST OF FIGURES .........................................................................................................................8ABSTRACT ...................................................................................................................... .............12 CHAP TER 1 INTRODUCTION .................................................................................................................. 141.1 Silicon Based Photovoltaic ............................................................................................... 141.2 Synthesis of Polycrystalline Si Film ................................................................................. 191.3 Present Approach .......................................................................................................... ....202 LITERATURE REVIEW .......................................................................................................272.1 Solid Phase Crystallization (SPC) of a-Si ........................................................................ 272.1.1 Nucleation Theory for Solid Phase Crystallization ................................................ 282.1.1 Growth Theory for Solid Phase Crystallization .....................................................302.2 Excimer Laser Crystallization (ELC) of a-Si ................................................................... 312.2.1 Laser-Solid Interaction ...........................................................................................322.2.2 Mechanisms of Eximer Laser Crystallization ........................................................ 353 EXPERIMENTAL TECHNIQUES ........................................................................................443.1 Growth Techniques ...........................................................................................................443.1.1 DC Magnetron Sputtering System ..........................................................................443.1.2 Plasma Enhanced Chemical Vapor Deposition (PECVD) .....................................453.2 Thin Film Annealing .........................................................................................................463.2.1 In-situ low pressure N2 Annealing .........................................................................463.2.2 Excimer Laser Annealing (ELA) System ............................................................... 463.3 Analytical Techniques ......................................................................................................463.3.1 X-ray Diffraction (XRD) ........................................................................................ 463.3.2 Raman Spectroscopy ..............................................................................................473.3.3 Scanning Electron Microscopy (SEM) ................................................................... 483.3.4 Transmission Electron Microscopy (TEM) ............................................................ 48

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6 4 NANOPARTICLE INDUCED CRYSTALLI ZATION OF AMOR PHOUS SILICON THIN FILMS BY THE DC MAGN ETRON SPUTTERING SYSTEM ............................... 554.1 Introduction .......................................................................................................................554.2 Experimental .....................................................................................................................564.3 Results and Discussion .................................................................................................... .574.4 Summary ...........................................................................................................................605 POST ANNEALING EFFECTS OF AMORP HOUS SILICON THIN FILMS BY NANOPARTICLE INDUCED CRYSTALLIZATION ......................................................... 715.1 Introduction .......................................................................................................................715.2 Experimental .....................................................................................................................725.3 Results and Discussion .................................................................................................... .735.4 Summary ...........................................................................................................................806 EXIMER LASER CRYSTALLIZATION OF T HE AMORPHOUS SILICON THIN FLMS INDUCED BY NANOPARTICLES ........................................................................ 1016. 1 Introduction ....................................................................................................................1016. 2 Experimental ..................................................................................................................1026. 3 Results and Discussion .................................................................................................. 1036.4 Summary .........................................................................................................................1077 CONCLUSION .................................................................................................................... .1247.1 Solid Phase Crystallization (SPC) of the Seeded Films ................................................. 1247.2 Excimer Laser Crystallization (ELC) of the Seeded Films ............................................ 125REFERENCES .................................................................................................................... ........127BIOGRAPHICAL SKETCH .......................................................................................................133

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7 LIST OF TABLES Table page 1-1 Current use and current potentials of selected renewable energy source ..........................211-2 Advantages and disa dvantages of photovoltaics................................................................224-1 Deposition rate of all the samples with average value and standard deviation with different RF powers ...........................................................................................................625-1 Comparison of the full-width at half-m aximum (FWHM) of the diffraction peaks at (111) orientation between the plain and nanoparticle-induced samples when annealed in air or low pressure N2 ambient ...................................................................................... 81

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8 LIST OF FIGURES Figure page 1-1 Schematic of a simple conventional solar cell. .................................................................. 231-2 Schematic of a solar cell. ............................................................................................... ....241-3 Theoretical maximum efficiency as a function of semiconductor band gap for an AM1.5 global spectrum ..................................................................................................... 251-4 Best small area (0.5-5cm2) efficiency for various t echnologies measured under standard laboratory test conditions. ................................................................................... 262-1 Temperature dependence of the therma l epitaxial crystallization rate for ion implanted a-Si layers on (100) silicon substrate ................................................................ 382-2 Free energy diagram of amorphous and li quid Si with respect to crystal Si. .................... 392-3 Reflectivity and absorption coefficien t values as a function of wavelength for crystalline silicon at room temperature ..............................................................................402-4 Schematic of indirect interband tr ansition in silicon in cluding phonon momentum exchange ...................................................................................................................... ......412-5 Schematic of explosive crystallization process .................................................................. 422-6 Schematic of the super lateral growth (SLG) .................................................................... 433-1 Schematic of motions of elec trons in crossed E and B fields. ........................................... 503-2 DC magnetron sputtering system ....................................................................................... 513-3 Schematic of in-situ vacuum annealing system ................................................................. 523-4 Heat cycle in a conventional furnace during annealing .....................................................523-5 Schematic of setup for laser anneal of samples in seeded nucleation experiments ........... 533-6 Energy level diagram showing the states involved in Raman signal .................................544-1 Variation of the deposition rate as a f unction of substrate temperature for two RF powers of 150 and 200W ...................................................................................................634-2 Average deposition rate as a function of RF power at substrate temperature of 300 C ... 644-3 XRD intensity as a function of substr ate temperature for two RF powers of 150 and 200W after annealing .........................................................................................................65

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9 4-4 GIXD patterns of silicon thin films: (a ) samples without and (b) with nanoparticle seeds before and after annealing for 10hr s at different temp erature in a air ...................... 664-5 The full-width at half-maximum (F WHM) of the diffraction peaks at (111) orientation between samples with and without nanoparticle seeds af ter annealing in a air .......................................................................................................................................674-6 GIXD pattern of a-Si thin films crysta llized at (a) 530C, and (b) 700C for 10hrs on a plain and nanoparticle embedded substrate annealed in a air ......................................... 684-7 Cross section TEM images and SAD patterns of Si thin films after annealing in a air for 10hrs at 700C ............................................................................................................ ..694-8 Plan-view TEM micrographs of samples (a) without and (b) w ith nanoparticle seeds after annealing in a air at 700C for 10hrs .........................................................................705-1 GIXD patterns of silicon thin films: (a) plain, (b) nanoparticle embedded films before and after annealing for 10hrs at different temperature under N2 ambient .............. 825-2 The full-width at half-maximum (F WHM) of the diffraction peaks at (111) orientation between the plain and nanopa rticle-induced samples annealed in N2 ambient ....................................................................................................................... ........835-3 Average grain size as a function of the annealing time for the SPC annealed samples with and without the nanopartcle seeds .............................................................................845-4 Plan-view TEM micrographs of polycrystalline Si nanoparticles ..................................... 855-5 XRD and SAD patterns of polycrystalline nanoparticle ....................................................865-6 TEM images of plain samples without nanoparticle seeds after annealing low pressure N2 for 3hrs at 700C ............................................................................................ 875-7 TEM images of samples without nanopart icle seeds after annealing in low pressure N2 for 10hrs at 700C .........................................................................................................885-8 TEM image of the sample without nanopart icle seeds after annealing in low pressure N2 for 20hrs at 700C .........................................................................................................895-9 SAD patterns of samples without nanopart icle seeds after annealing in low pressure N2 for 10 and 20hrs at 700C, respectively ....................................................................... 905-10 TEM images of nanoparticle embedded samples after annealing in low pressure N2 at 550C for 10hrs ..................................................................................................................915-11 TEM images of nanoparticle embedded samples after annealing in low pressure N2 at 600C for 10hrs ..................................................................................................................92

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10 5-12 TEM images of nanoparticle embedded samples after annealing in low pressure N2 at 650C for 10hrs ..................................................................................................................935-13 TEM images of nanoparticle embedded samples after annealing in low pressure N2 at 700C for 10hrs ..................................................................................................................945-14 TEM images of nanoparticle embedded samples after annealing in low pressure N2 at 800C for 10hrs ..................................................................................................................955-15 SAD patterns of nanoparticle embedded samples after annealing in low pressure N2 for 10 and 20hrs at 700C, respectively .............................................................................965-16 Temperature dependency of the measured incubation time for crystallization samples without and with nanoparticle seeds .................................................................................. 975-17 Crystalline volume fraction as a function of annealing time at va rious temperatures ....... 985-18 Temperature dependence of the SPE grow th velocity in amorphous Si thin film deposited on poly-Si nanoparticle seeds ............................................................................ 995-19 Temperature dependency of the median crystallization time, t50 ................................... 1006-1 Raman spectroscopy of plain samples at various laser energy densities indicated ......... 1086-2 Raman spectroscopy of nanoparticle em bedded samples at different laser energy densities............................................................................................................................1096-3 Volume fraction as a f unction of various laser energy densities with and without nanoparticle seeds ............................................................................................................ 1106-4 SEM images of dispersed nanoparticles on the SiO2/glass substrate .............................. 1116-5 SEM images of plain samples anneal ed at laser energy density of 230mJ/cm2 ..............1126-6 SEM images of plain samples anneal ed at laser energy density of 320mJ/cm2 ..............1136-7 SEM image of the super lateral growth (SLG) annealed at laser energy density of 350mJ/cm2........................................................................................................................1146-8 SEM images of nanoparticle embedded sa mples annealed at laser energy density of 200mJ/cm2........................................................................................................................1156-9 SEM images of nanoparticle embedded sa mples annealed at laser energy density of 260mJ/cm2........................................................................................................................1166-10 SEM images of Si thin films irra diated by excimer laser at 290and 320mJ/cm2, respectively .................................................................................................................. ....117

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11 6-11 AFM images of samples without nanopart icle seeds annealed at laser energy density of 230 and 350mJ/cm2 .....................................................................................................1186-12 AFM images of nanoparticle embedded sa mples annealed at laser energy density of 200 and 290mJ/cm2 ..........................................................................................................1196-13 RMS as a function of laser energy density in samples with and without nanoparticle seeds ......................................................................................................................... ........1206-14 Grain size distribution as a function of various lase r energy densities with and without nanoparticle seeds ............................................................................................... 1216-15 TEM images of nanoparticle seeded samp les annealed at laser energy density of 290 and 320mJ/cm2 .................................................................................................................1226-16 SAD patterns of nanopartic le seeded samples annealed at laser energy density of 290 and 320 mJ/cm2 ................................................................................................................123

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12 Abstract of Dissertation Presen ted to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy CRYSTALLIZATION AND CHARACTERIZATI ON OF AMORPHOUS SILICON THIN FILMS INDUCED BY NAN OPARTICLES SEED By Taekon Kim May 2009 Chair: Rajiv. K. Singh Major: Materials Scie nce and Engineering Crystallization of amorphous Si (a-Si) thin f ilm has received extens ive interest for their attractive applications into Si thin film transistors and Si based solar cells. Among various crystallization techniques, Solid phase crystallization (SPC) and Excimer laser crystallization (ELC) were investigated. Firstly, Solid phase cr ystallization ( SPC) of amorphous silicon thin films deposited by the DC magnetron sputtering syst em with a modification in nucleation step was investigated at low temperature. The thin film consists of polycrystalline nanoparticles embedded in an amorphous matrix which can act as nuclei during crystall ization, resulting in a lower thermal energy for the nucleation. The lo wering energy barrier for nucleation would shorten the transition time from amorphous into polycrystalline silicon resulting from the reduction of incubation time and also lower th e processing temperature spontaneously. In addition, a comprehensive study of the growth mechanism of the sputtered amorphous silicon thin films is presented during annealing. Samp les were prepared with various substrate temperatures and RF power in order to optimi ze the crystallization of a-Si after the deposition. Also, the effects of annealing condi tion were examined. Low pressure N2 ambient during SPC promoted crystallization of a-Si thin films and the crystallinity. The low pressure annealing had a large impact on the crystallinity and growth behavior of subse quent films. In addition, the

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13 crystallinity, incubation time, the crystallized volume fraction and growth rate of the films annealed in a conventional furnace have been extensively studied by XRD and HRTEM. It was believed that crystalline Si nanopa rticles would act as nuclei for gr owth of crystalline Si thin films, thus removing the high temperature requirement for nucleat ion, resulted in the improvement of the crystallization of a-Si. Secondly, the controlled Super Lateral Growth (SLG) can be obtained by nanoparticle induced crystallization (NIC) t echnique during laser annealing, which led to enhance the random super lateral growth (SLG) of Si thin films for the excimer laser crystallization (ELC). The crystallinity and surface informati on of the films irradiated by excimer laser have been studied by Raman spectroscopy, FESEM and AFM. Also, Transmission electron microscopy (TEM) was employed in order to obtain structural informati on. Polycrystalline Si nanoparticles, which have higher melting point than those of amorphous pha se, would survive at high energy density of laser. In general, super lateral growth (SLG) occurs at vary narrow laser energy density region. Thus, it tends to be sensitive to laser energy dens ity, which means not easy to control because of the characteristics of the mechanism of the SLG. In this study, poly-Si na noparticles would act as nucleation seeds for the growth of the films during the solidification. Those nanoparticle seeds provided more probability to survive at higher density of energy compared to that without nanoparticle seeds, resulted in the large grain si ze distribution and the co ntrolled super lateral growth (SLG), relatively indepe ndent of laser energy density.

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14 CHAPTER 1 INTRODUCTION 1.1 Silicon Based Photovoltaic Photovoltaic has the hug e potential for future energy supply system. Recently, a considerable amount of money has been widely invested in the research, development and demonstration. Table 1-1 shows the technical and the theoretical po tentials of several renewable energy sources such as biomass, wind and solar en ergy.[1] Both potentials of PV electricity is high enough to contribute considerably to the abatement of the global CO2 problem. Photovoltaic (PV) is an important energy technology for many reasons such as environmental benefits. As a dome stic source of electricity, it contributes to the nation's energy security. As a relatively young, hightech industry, it helps to crea te jobs and strengthen the economy. As it costs increasingly less to produce and use, it becomes more affordable and available. Furthermore, few power-generati on technologies have as little impact on the environment as photovoltaics. As it quietly generates electricity from light, PV produces no air pollution or hazardous waste. It doesn't require liquid or gaseous fuels to be transported or combusted. And because its energy source sunl ight is free and abundant, PV systems can guarantee access to electric power. PV frees us from the cost and uncertainties surrounding energy supplies from politically volatile regions Table 1-2 lists some of advantages and disadvantages of photovoltaics. Therefore, Photovo ltaics is an excellently suitable solution for low power electricity supply in rural and remote areas in deve loping countries. Photovoltaics have a similar but smaller market in industria lized countries as well. All these solar energy technologies put the sun's energy to work fo r us in our homes, schools, businesses, and government buildings. They are being developed b ecause they are reliable, they have very few environmental impacts, and they make use of an abundant domestic energy resource: sunlight.

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15 Semiconductor solar cells fundamentally quite simply devices. Semiconductors have the capacity to absorb light and to deliver a portion of the energy of the absorbed photons to carriers of electrical current electrons and holes. A semiconductor diode separates and collects the carriers and conducts the generated electrical current preferentially in a specific direction. Thus, a solar cell is simply a semiconductor diode that has been carefully designed and constructed to efficiently absorb and convert light energy from the sun in to electrical energy. A simple conventional solar cell structure is depicted in Figure 1-1. All electromagnetic radiation, including sunlig ht, is composed of particles called photons, which carry specific amounts of energy determined by the spectral proper ties of their source. Photons also exhibit a wavelike character with the wavelength, being related to the photon energy, E, by hc E (1.1) Where h is Planks constant a nd c is the speed of light. Only photon with sufficient energy to create an electron-hole pair that is, those with energy grea ter than the semiconductor bandgap (Eg), will contribute to the energy conversion process. Thus, the spect ral nature of sunlight is an important consideration in the de sign of efficient solar cells. Solar cell are made of ma terials called semiconductors, which have weakly bonded electrons occupying a band of energy called the va lence band. When energy exceeding a certain threshold, called the band gap energy, is applie d to a valence electron, the bonds are broken and the electron is somewhat free move around in a new energy band called the conduction band where it can conduct electricity through the material. Thus the free electrons in the conduction band are separated from the valence band by the ba nd gap (measured in units of electron volts or eV). This energy need to free the electron can be supplied by photons, which are particles of

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16 light. Figure 1-2 shows the idealized relation be tween energy (vertical axis) and the spatial boundaries (horizontal axis). When the solar cell is exposed to sunlight photons hit valence electrons, breaking the bonds and exciting them to the conduction band. Th ere, a specially made selective contact that collects conduction-band electrons drives such el ectrons to the external circuit. The electrons lose their energy by doing work in the external circuit such as pumping water, spinning a fan, powering a sewing machine motor, a light bulb, or a computer. They are restored to the solar cell by the return loop of the circuit via a second selective contact, which returns them to the valence band with the same energy that they started w ith. The movement of these electrons in the external circuit and contacts is called the electric current. The poten tial at which the electrons are delivered to the external world is slightly less than the threshold energy that excited the electrons; that is, the band gap. Thus, in a material with a 1 eV band gap, electrons excited by a 2 eV photon or by a 3 eV photon will both still have a potential of slightly le ss than 1 eV (i.e. the electrons are delivered with an energy of 1 eV ). The electric power pr oduced is the product of the current times the voltage; that is, power is th e number of free electrons times their potential. Since only photon with hv > Eg can create electron-hole pairs and contribute to the output of the solar cell, it is obvious that the band gap determines ho w well the solar cell is coupled to the solar spectrum. A simple analysis can be performed to predict the maximum solar cell efficiency. More complete analyses of the theoretic al limits of solar cells are given elsewhere.[24] Assuming the maximum energy that can be extracted from an absorbed photon is Eg, the maximum efficiency can be expressed as Gdf AP E P IE q Ein G in incG G )( )/( 1 )(max (1-2)

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17 This is plotted in Figure 1-3 for an AM1.5 global spectrum and shows a maximum efficiency of 48% at about Eg=1.1 eV, close to the band gap of silicon. Of course, this is only a simple estimate and assumes that Voc=1/qEg and FF=1, which are obvious exaggerations. Perfect light trapping was also assumed so that Isc=Iinc, but that is a more r ealistic prospect. Under nonconcentrated solar illumination, th e actual maximum theoretical effi ciency for a silicon solar cell is approximately 30%. However, this simple ap proach does serve to demonstrate the important role the semiconductor band gap plays in determ ining solar cell perfor mance and shows that band gaps between 1.0 and 1.6 eV have nearly eq uivalent maximum theoretical efficiencies. Figure 1-4 shows the efficien cy for PV materials with band gap energy between 1.0 and 1.6 eV. From solid state physics we know that silic on is not the ideal material for photovoltaic conversion. It is a material with relatively low absorption of so lar radiation, and, therefore, a thick layer of silicon is required for efficien t absorption. Theoretically, light absorption is impeded because it requires a change of momentum Thus, materials with a direct band structure which gives very strong light absorption are more suitable for solar cell technology. They belong to the class of compound semiconductors like GaAs or InP, which are III-V compounds according to their position in the periodic table. Other important groups are II-VI and I-III-VII compounds, which, just like the elemental semic onductors, have four bonds per atom. However, there are problems with low efficiency and insuffi cient stability prevented further penetration of Copper Indium Diselenide (CIS), and Cadmium Telluride (CdTe). The new technology is based on the te rnary compound semiconductors CuInSe2, CuGaSe, CuInS2 and their multinary alloy Cu(In,Ga)(S,Se)2 (CIGS). The first results of single crystal work on CuInSe2 were extremely promising, but the co mplexity of the material looked complicated as a thin film technology.

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18 In spite of the complicated manufacture a nd the high cost, crysta lline silicon still dominates the market today and probably will con tinue to do so in the immediate future. This is mostly due to the fact that there is an a bundant supply of silicon as raw material, high efficiencies are feasible, the ecological impact is low, and silicon in its crystalline form has practically no degradation. Furthermore, detail ed models that take into account surface characteristics and the multi-reflections within the wafer show that absorption can be greatly enhanced.[5, 6] Therefore, a very thin layer of Si can offer a high degree of absorption of the solar spectrum-nearly as much as a thic k wafer and direct band gap materials. Besides, such a thin-film Si (TF-Si) solar ce ll offers many advantages that can lower the cost of generating solar electricity. A TF-Si cell offers bellows: (1) Reduced bulk recombination leading to lower dark current, higher Voc and higher FF of the device. Compared to a thick cell, a thin cell of the same mate rial quality can yield higher device performance. Likewise, for a comparable performance, TF-Si solar cell requires lower material quality than a thick cell. (2) Potential for low cost cells/module (3) Potential for lightwe ight photovoltaics, (4) Lower energy consumption for device fabrication (5) Potential for flexible solar cells. These advantages of TF-Si solar cells, in c oncurrence with the pe rformance advantage, make the very attractive for the future. Practic al realization of solar cells with the above advantages poses many challenges in both the de sign and device fabrication. These challenges include an efficient method for light-trapping to compensate for reduced thickness, and a lowcost substrate to support the thin film. Low-cost substrates generally imply materials that may not be compatible with the high temperature requ ired for formation and processing of Si film. This incompatibility can arise because of impurities in the substrate that can diffuse into the Si

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19 film, softening of the substrate, thermal mismatch, and less desirabl e electronic properties of the interface. 1.2 Synthesis of Polycrystalline Si Film Many researches have been perform ed to lowe r processing temperature in order to reduce the cost. Generally, there are two methods to obtain polycrystalline silicon thin films on insulators. One is a direct deposition method with the use of disilane (Si2H6) as a precursor or plasma enhanced chemical vapor deposition (PEC VD).[7, 8] In this method, the grain sizes are very small which is deleterious for electrical performance. Also, the surface is very rough, thus further hampering the properties of the devices. To engineer the polySi morphology a number of approaches are followed based manipulation of the poly-Si ch emical vapor deposition (CVD) process,[9, 10] on excimer laser crystallization (ELC), of a-Si films,[11, 12] or on solid phase crystallization (SPC) of a-Si.[ 13-20] Of these techniques, SPC is one of the most promising methods to obtain poly-Si films from the amor phous phase. Compared to CVD poly-Si films, poly-Si films obtained by SPC of a-Si films are characterized by a larger grain size, i.e., by a reduced density of grain boundaries, and indeed, some of the electri cal characteristics have been improved by using SPC poly-Si films.[13, 14] Un fortunately, the SPC te chnique significantly increases the crystallization temperature and process time which are deleterious for manufacturing. SPC of a-Si has r ecently met with renewed interest since the discovery that the presence of metal impurities strongly modifies th e crystal grain nucleation and growth kinetics, allowing considerable improvement of the electrical char acteristic.[19, 20] For excimer laser crystallization, in contrast to solid-phase crystallization (SPC), the silicon film is melted for very short time by laser annealing. This te chnique provides a low thermal budget and low defect density in grains due to melting and rapid regrowth, resulted in very fine grain distributi on in the range of 0.1~0.2 m.[21] However, the low defect density

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20 inside the grains have led to give ri se to high carrier mobilities (100~300cm2/Vsec).[22] Despite the large mobilities produced by th is method, there are several draw backs. The formation of very small grains gives rise to surface non-uniformity which leads to non uniform oxide growth. In addition, uniformity in processing is difficult due to small size of the laser beam and due to very sensitive to laser energy density fo r the super lateral growth (SLG). 1.3 Present Approach In order to fabricate very large grain sized poly cr ystalline Si thin films on insulators at low temperatures and with short processing times, we have investigated a new technique in this study. This technique is based on th e application of a nanoparticle seeds to induce the solid phase crystallization (SPC) of amorphous Si thin films. The application of a nanoparticle seed leads to improvement of the growth of the silicon film, resu lting in the formation of large grains. Grains larger than 2 m have been obtained. Studies to understand the nature of the growth process have been investigated in this disse rtation. Also, low pressure N2 annealing promoted the crystallization of a-Si thin films during SPC. In addition, excimer laser crysta llization (ELC) was car ried out to control the super lateral growth ((SLG) when nanoparticles were embe dded in amorphous Si matrix. Nanoparticle embedded samples irradiated by excimer laser show ed large grain size dist ribution, relatively not sensitive to laser energy density.

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21 Table 1-1 Current use and current potentials of selected renewable energy source Resource Current use Technical potential Theoretical potential Hydropower 9 50 147 Biomass energy 50 >276 2900 Solar energy 0.1 >1575 3900000 Wind energy 0.12 640 6000 Units: exajoule per year

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22 Table 1-2 Advantages and di sadvantages of photovoltaics Advantages of photovoltaics Disadvantages of photovoltaics Fuel source is vast and seentially infinite Fuel source is diffuse (sunlight is a relatively low-density energy) No emissions, no combustion or radiative fuel for disposal (does not contribute perceptibly to global climate change or pollution) High installation costs No moving parts (no wear) Ambient temperature operation (no high temperature corrostion or safety issues) High reliability in modules (>20years) Poorer reliability of auxiliary (balance of system) elements including storage Modular (small or large increments) Quick installation Can be integrated into new or existing building structure Can be installed at nearly any point of use Lack of widespread commercially available system integration and installation so far Daily output peak may match local demand Lack of economical efficient energy storage High public acceptance Excellent safety record

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23 Figure 1-1 Schematic of a simple conventional sola r cell. Creation of electron-hole pairs, eand h+, respectively, is depicted

PAGE 24

24 Figure 1-2 Schematic of a solar cell. Electrons are pumped by photons from the valence band to the conduction band. There they are extracted by a contact selectiv e to the conduction band (an n-doped semiconductor) at a highe r (free) energy and delivered to the outside world via wires, where they do some useful work, then are returned to the valence band at a lower (free) energy by a contact selective to the valence band (a ptype semiconductor)

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25 Figure 1-3 Theoretical maximum efficiency as a function of semiconductor band gap for an AM1.5 global spectrum

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26 Figure 1-4 Best small area (0.5-5cm2) efficiency for various technologies measured under standard laboratory test conditions. MJ c oncentrators are double junctions before 1995, and triple junctions after. a-Si repres ents stabilized efficiency after extended light soaking and are MJ after 1990

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27 CHAPTER 2 LITERATURE REVIEW In general, the for mation of polycrystal line Si films from the amorphous phase is controlled by nucleation and growth processes. The pre-existing nucleatio n seeds would affect the growth behavior which is important in cont rolling the microstructure of the films for solid phase crystallization (SPC). In addition, laser annealing of amorphous Si films has been widely investigated due to the excelle nt electrical properties. The nanoparticles embedded in amorphous matrix also would act as nuclea tion seeds for excimer laser crysta llization (ELC), resulted in the controlled super late ral growth (SLG). In this chapter, the literature for nuclea tion and growth mechanisms in solid phase crystallization (SPC) is reviewed. This is follow ed by a discussion of seed -based crystallization related to my research. Also, lit erature pertaining to excimer laser crystallization (ELC) and laser-solid interactions of Si films are intr oduced. The rest of the chapter is devoted to mechanisms of laser crystallization of thin films. 2.1 Solid Phase Crystallization (SPC) of a-Si Solid phase crystallization of a-Si on insulators such as glass has been w idely investigated.[23-29] The thermodynamics and kinetic s of solid phase crysta llization of Si films from amorphous phase has been explained by classical nuclea tion and growth theory. When amorphous Si films are annealed to a certain temperature, the film is transformed into thermodynamically stable crystalline phase through four steps, including incubation, nucleation, growth and steady state.[30-32] In this process, small crystallites as nucleation site are formed and then grow with time in amorphous matrix. In cas e of a-Si films with a crystalline layer, this transition occurs at the crystalline and amor phous interface referred to as solid-phase epitaxy (SPE). The growth rate of the cr ystallization is strongly dependent of temperature. It follows an

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28 Arrhenius behavior with an act ivation energy of 2.68.05eV over a growth rate range of more than 10 orders of magnitude.[30, 32] as shown in Figure 2-1 wh ich reported by Olson and Roth for a-Si on top of <100> oriented c-Si are shown.[30] Also, it is pres ented that the impurity content strongly affects the kinetic s of SPE, indicating dopants such as B, P, and As can promote the growth rate by more than one order of magni tude.[30, 33, 34] while contaminants such as C, O, F, or noble gases strongly decelerate the gr owth rate of films.[30, 35, 36] Moreover, the presence of n-type and p-type dopants also produces a compensating effect in the growth rate.[30, 37, 38] Another important result is the dependence of interface velocity on substrate orientation.[30, 39] Also, Spaepen and Turnbull ha ve reported a structur al model of thermal annealing in which the crystalline and amorphous interface resolved into a free energy minimum by the formation of terraces w ith a <111> orientation.[40] Thus some defects in the amorphous phase strongly affect for the crystallization process. Figure 2-2 shows the free energy diagram of a-Si (in both the relaxed and unrelaxed states) and of liquid Si with respect to crystalline Si.[41, 42] It is cl ear the a-Si has higher free energy than c-Si, and thus it exists as a kinetically meta -stable phase. The melting temperature of a-Si is predicted to be smaller than that of c-Si a nd to depend on the relaxa tion state. These last properties have indeed experime ntally verified by ultra-shor t pulsed laser annealing.[42, 43] 2.1.1 Nucleation Theory for Solid Phase Crystallization The crystallization of amorphous Si film s generally occurs through a nucleation and growth process.[29] These phenomena have be en explained by the cl assical nucleation and growth theory as mentioned previously.[23-29] A ccording to the classical theory, the nucleation of the films occurs beca use of the reduction of th e total free energy of the system. The Gibbs free energy change for this nucleation process, Gn, is determined by the difference in volume free energy change for atoms in the crystalline phase relative to that of th e amorphous phase and the

PAGE 29

29 amount of interface energy which is generated by the a-Si/c-Si inte rface, if the strain energy is small enough to be neglected. The energy change due to formation of an n-mer of a nucleating phase is given by[44] VGSVGGv n (2-1) where Gv is the free energy difference per unit vol ume due to the phase transformation, V is the volume of an n-mer of the nucleating phase, is the interfacial free energy between the cluster and the parent phase, S is the surface area of the cluster, G is the strain energy per unit volume associated with the formation of the cluster. In order to describe the cluster of arbitrary shape and configuration, atomic volume, and the shape factor, g( )=(2+cos )(1-cos 2)/4, are introduced. The volume of an n-mer cluster and surface area can be expressed by V=n and S=g( ) 2/3n2/3. Equation (2-1) can be simplified as follows: 3/23/2)( n gnGGv (2-2) When the first derivative of Equation (2-2) equa ls to zero, the cluster reaches the critical size which can be a nuclei. From this, we can derive the critical free energy, G*k. and the number of atoms in the critical cluster, k. 3]/)([)/1()27/8(vGg k (2-3) v kGk G 2/1* (2-4) For time dependent nucleation rate it was reported by Turnbull[45] as )/exp( )(* t JtJk S (2-5) where t is the time and is an incubation time. The incubation time can be derived by the assumption of random walk model in the critical region.[44] In this approximation, the incubation time, is defined as the time for a cluster to

PAGE 30

30 move from the subcritical size, k/2, to a supercritical size, k+ /2 along the Gibbs free energy barrier with the random walk distance, The value of is computed from the Gibbs free energy change equation by setting TkGGB n k (2-6) 2/ nk (2-7) Thus, 2/12 2])//(8[k n BnGTk (2-8) In the random walk theory, the time which is necessary for a part icle to displace the distance, in one dimension is given by k/2 (2-9) Therefore, the incubation time, can be expressed as follows: ])/*(/[82 2 k n kBnG Tk (2-10) 2.1.1 Growth Theory for Solid Phase Crystallization The grain growth mechanism in solid phase crys tallization of amorphous Si is the same as that of solid phase epitaxial grow th (SPE). However, this crysta llization needs a nucleation step before the growth process occurs. Due to this complication, an accurate measurement of the growth rate becomes very difficult. The growth rate is determined by fitting the crystallized volume and nucleation rate into the Avramis equation.[23] Another method to measure the growth rate is a direct observat ion of grains to grow with th e TEM.[24] The reported activation energies for the grain growth, (2.4~3.3eV), are similar to those for epitaxial regrowth, (2.4~2.9eV).[23-25]

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31 Grain growth mechanisms have been studied by many of researchers.[26-28] Nakamura et al. found that the grains of (110) surface plane ha ve a preferential growth in the <112> direction along twin boundaries. It results in the formation of dendrite sh aped grains. Noma et al. had investigated the grain growth pr ocess for differently oriented grains. The growth rates for the grains have the same trend as those of the epitaxial regrowth (v(110)>v(110)>v(111)). The nucleated grains are bounded by the slowest growing plane, (111), durin g the crystallization process. Therefore, the growth rate of grains is determined by the rate of forming the (111) plane. As mentioned in previous reviews of epitaxial gr owth, the (111) plane is difficult to grow due to more incorporated atoms for migration of the a-Si/c-Si interface than other planes. In (111) growth, micro-twins are formed on the inte rface to enhance the growth. An end point of twin plane provides atomic steps as nucleation sites for the growth.[28] According to Nomas model, the shapes of variously oriented gr ains are decided by a location of (111) twin planes w ith respect to film surface. For example, the <110> oriented surface grains have a preferential growth in the <211> direction and branch in other coplanar <221> directions. This results in a dendrite shape. On the other hand, when (111) twin planes are parallel to the film surface in <111> oriented grains, the shapes become disc type.[27] Based on the work by others, we know that the randomly nuc leated grains usually form micro-twins during the grain growth. From reviews in this section, it is known that the mechanism and the trend of solid phase growth of crystalline Si particle in a-Si thin films is the same as those of the epitaxial growth. 2.2 Excimer Laser Crystallization (ELC) of a-Si Since the first results relate d to laser annealing of ionimplanted semiconductors were obtained, many of papers in the field has b een reported enormously.[46-48] As mentioned previously, laser crystallization of poly-Si films results in the highest carrier mobility compared

PAGE 32

32 to those attained by any technique. However, al though researches on this technique have been actively performed,[12] the tr ansformation mechanisms are not yet well understood. Also, attempts at quantitative estimates of the effect s of high-power laser irradiation have been limited due to the lack of precise temperature a nd phase-dependent optical and thermophysical parameters.[49] Literature review s related to this topic follow. 2.2.1 Laser-Solid Interaction Electromagnetic radiation with wavelength ranging from ultrav iolet to infrared interacts exclusively with electrons, as atoms are t oo heavy to respond significantly to the high frequencies ( > 1013 Hz).[50] Therefore, the optical prope rties of matter are determined by the energy states of its valence electrons (bound or free). Bond el ectrons generally only weakly respond to the external electromagnetic wave and affect only its phase velocity. Free electrons, however, can be accelerated and th erefore extract energy from th e field. Since the field is periodically changing, th e oscillating electrons rera diate their kinetic energy (cause of reflection) or collide with the atoms, giving their energy to the lattice. Absorption of incident energy fundamentally dictates the resultant thermal state of the material and therefore is a suitab le point to begin an analysis of laser-solid interactions. The complex refractive index, n*, defines two quantities which descri be the degree of coupling of the incident radiation with the material, name ly reflectivity R and absorption coefficient It is defined as[51]: iknn (2.1) Where n, the real part, is the ratio of the pha se velocities in vacuum and the material. The extinction coefficient k describes the dampi ng of the light wave. For normal incidence, reflectivity and absorption are related to n and k by, respectively:

PAGE 33

33 22 22)1( )1( kn kn R (2.2) and k 4 (2.3) where is the wavelength of the radiation. As mentioned previously, the m echanisms involved in absorpti on of incident radiation in materials are defined by the electronic structure of the material, and therefore it is useful to discuss exclusively semiconductors. In semic onductors, five distinct mechanisms for the absorption of light can be identified.[48] 1) Photons with energy (h ) much less than the band-gap energy (Eg) can excite lattice vibrations directly. 2) Free or nearly free carriers can be ex cited by absorption of light with h < Eg; such carriers will always be present as a result of finite temperature and doping. 3) An induced metallic-like absorption due to free carriers generated by the laser radiation itself can occur. 4) For photon energies > Eg absorption will take place by dir ect and indirect (photon-assisted) excitation of electron-hole pairs. 5) Absorption induced by broken symmetry of the crystalline lattice is possible. The largest contributions to abso rption of laser radiation with h > Eg by crystalline or amorphous silicon are found in mechanisms (4) a nd (5), respectively. Mechanism (3), however, may be the cause of discrepancies of actual as compared to calculated absorption coefficient value.[52] Through the focus here is on absorption of ultra-violet radia tion (KrF excimer laser operating at 248nm) by amorphous or fine-grained polycrystalline silicon, the state of the material is not particularly cr itical in this instance because is saturated at ~106 cm-1 for this wavelength in silicon. This value is more characte ristic of metals, and is virtually independent of

PAGE 34

34 both temperatures[49] and state of the material, i.e., it is approximately the same for crystalline, amorphous, and liquid silicon. An analysis of ab sorption coefficient re ported by Jellison, et al.[53] reinforces this as no change in was observed over the temp erature range 300-1000K for 355nm Nd:YAG radiation. Reflectivity value for silicon depends on wa velength as well as crystalline state. Absorption coefficient and reflectivity values for crystalline silicon at room temperature as a function of wavelength are shown in Figure 2-3. From this figure it can be seen that the reflectivity demonstrates rela tively high values in the UV regime and represent R~65% for 248nm radiation. Through the reflectivity is relatively high in the UV regime, the absorption coefficient is at this maximum values. Silicon is an indirect band-gap material.[54] This means that the maximum of the valence band do not occur at the same point as the mini mum of the conduction band in k-space, where k is the wave vector. As mentioned prev iously, absorption of radiation with h Eg predominantly occurs by inter-band transitions, this case including those of an indirect nature. An indirect transition requires a change in both energy and momentum (the qua ntity k is proportional to the momentum). A two-step process is therefore required because the photon can provide a change in energy but not (significan tly) in momentum. Momentum is conserved thorugh phonon interaction as illustrated in Figure 2-4. Through a phonon can only absorb very small energies, it is able to absorb a large mome ntum when compared to an elec tron.[51] Also, although a broad spectrum of phonons is available, only those with the required momentum change are usable. Direct transitions are al so possible in an indirect gab materi al, if the photon energy is sufficient (h > 3.4 eV for Si). This explains in the hi gh absorption coefficient values observed for UV radiation (h ~ 5eV).

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35 When a beam of photons of energy h > Eg is absorbed in a semiconductor, excited carriers, which results in latti ce heating.[55], is a complicated process and an field of active research.[56] Excited carrier rela xation times on the order of pico seconds have been derived and confirmed using time-resolved investigation of the absorption coefficient. As incident radiation is converted to increasing lat tice temperatures, the thermophysic al properties of the material dictate temperature distribution and phase changes. This aspect of laser a nnealing area has been actively investigated.[57-61] 2.2.2 Mechanisms of Eximer Laser Crystallization The basic mechanism of laser heating pr oceeds through photon absorption and the subsequent rapid transfer of energy (within pi cosenconds) from the elec trons to the lattice.[6264] During laser crystallization of amorphous silic on thin films, the incident radiation leads to surface melting when sufficient energy employed. This liquid phase becomes unstable as rapid cooling takes place after the pulse is over. The mechanisms (nucle ation and solidif ication) that then take place and lead to the resultant micros tructure have been wide ly investigated.[12, 22, 47, 65, 66] Nucleation and solidific ation processes in pulsed laser crystallization of silicon thin films can be divided into three regimes, differe ntiated by whether complete melting of the film occurs. First, the energy density is not high enough to melt the surface film completely in partialmelting regime. After melting, ther e exists an unmelted crystalline layer at the bottom of the film. The final microstructure in this regi me results in a fine-grained c-Si In near-complete-melting regime, the SLG regime, where the ratio of grain size to film thickness can exceed a factor of ~10 is due to significant lateral regrowth from unmelted and discontinuous solid seeds at the bottom oxide interface. This is evident at the higher end of the partial-melting regime, where a small increase in the energy density produces an extremely sharp

PAGE 36

36 increase in the grain size. To date, these laterally grown poly-Si films correspond to the largest average grain sized poly-Si films produced by a single-pulse ELC process.[12, 21] Finally, in complete-melting regime, a sharp and sudden increase in the melt duration is strong indication of a transition to complete melting followed by super-cooling and transformation of the film. Before the onset of the solidification, substantial super-cooling followed by either transition or steady-state nuc leation, and then grow th of solids occurs. Typically, the microstructure in this regime is fi ne-grained Si. This regime can be compared with recent investigation on super-cooling of a liquid Si film by S. Stiffler et al.[67-70] It was found that when the cooling rate of the liquid Si exceeds about 3x1010 K/s, homogeneous nucleation of the solid takes place, resulting in a fine-grained Si microstructure. Based on three regimes, there are two mechanis ms for excimer laser crystallization of a-Si; the explosive crystallization and the super lateral growth (SLG). In the first case, where complete melting of the film does not occur, several potential theories have been proposed. Thompson et al.[ 71] and Lowndes et al.[72] support the following dynamic. At low energy densities the laser energy only melts a portion of the film, creating a thin liquid layer near the surface. As the liquid be gins to solidify as relatively large-grain polycrystalline silicon, the latent heat released raises the temper ature of the resolidified poly-Si above the temperature of the first order phase transition of a-Si to the metallic liquid, T -1. The underlying a-Si material then begi ns to melt. This new liquid is severely under-cooled compared to the poly-Si layer and therefore resolidifies as fine-grained poly-Si. Thus a thin liquid layer is presumed to propagate through the aSi material as a result of the re leased latent energy. This is the explosive recrystallization process. It is se lf-sustaining until eventually it is quenched by the energy required to raise the temperature of th e a-Si solid in front of the liquid to T -1. The final

PAGE 37

37 microstructure consists of a top layer of larg e-grained (50-100nm) polycry stalline silicon with a layer of fine-grained (1-10nm) poly-Si directly below it resulting from the explosive recrystallization pro cess as shown Figure 2-5. Thompson et al.[71] were able to support the explosive crystallization argument by use of transient reflec tance and conductance measurements. The other proposed mechanism for non-complete melting of Si films is by Im et al.[12] Im argues that in his films explosive crystall ization of a-Si occurs at the onset of the transformation, implying that pa rtial melting of explosively crystallized fine-grained Si is occurring rather than pa rtial melting of an amorphous film. It is suggested that early triggering of explosive crystallization may be attributed either to (i) the presence of microcrystalline clusters in the LPCVD samples, which was confirmed by analyzing solid-phase crystallization behavior and is absent in high-dose ion irradiated am orphized samples which called the Super Lateral Growth (SLG) regime and (ii) the possible pres ence of impurities, such as hydrogen. In this regime, almost complete melting of the film occurs to the extent that there is a discontinuous Si film composed of discrete islands. Growth from these clusters proceeds, resulting in an unusually large (300-400nm) grain size. If greater laser energy is used the resultant grain size returns to smaller values (~50-100nm), typi cal of lower-energy density irra diation as shown in Figure 26.[12]

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38 Figure 2-1 Temperature dependence of the thermal ep itaxial crystallization rate for ion implanted a-Si layers on (100) silicon substrate[30]

PAGE 39

39 Figure 2-2 Free energy diagram of amorphous and liqui d Si with respect to crystal Si. Diagrams for both relaxed (continuous line) and unrelax ed (dashed line) a-Si are shown. Note that the melting temperature of a-Si is different from that of crystal silicon and it depend on the relaxation state[41, 42]

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40 Figure 2-3 Reflectivity and absorption coeffici ent values as a function of wavelength for crystalline silicon at room temperature[50]

PAGE 41

41 Figure 2-4 Schematic of indirect interband transition in silicon including phonon momentum exchange[51]

PAGE 42

42 Figure 2-5 Schematic of explosive crystallization process

PAGE 43

43 Figure 2-6 Schematic of the s uper lateral growth (SLG)[12]

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44 CHAPTER 3 EXPERIMENTAL TECHNIQUES This chapter will cover the experim ental methods for thin film deposition, annealing techniques and characterization of the Si films. The first sections will describe the deposition systems for the thin films, and the procedures used to prep are the samples. All films used in this work were grown with the DC magnetron sputtering system and Plasma Enhanced Chemical Vapor Deposition (PECVD) system. In addition, annealing systems were introduced for solid phase crystallization and excimer laser crystallization. The following section will explain the techniques used to measure the properties of the thin films. Various analyzing t ools were employed to investigate the structural, electrical properties of the films. These techniques include X-ray diffraction (XRD), Raman Spectroscopy, Scanning electron microscopy (SEM), and Transmission electron microscopy (TEM). 3.1 Growth Techniques 3.1.1 DC Magnetron Sputtering System Sputter depo sition is one of the most widely used techniques for the fabrication of thin-film devices on the desired substrate. It is used primarily for the deposition of metal thin films, as well as the various related thin films which function as diffusion barriers, adhesion or orientation layers, or seed layers. Sputter deposition is usually carried out in diode plasma systems known as magnetrons, in which the cathode is sputtered by ion bombardment, and emits the atoms, which are then deposited on the desired substrate in the form of a thin film. The most widely used technology for sputter deposition is based on the magnetron cathode. Originally, physical sputter deposition utilized dc diodes, which were simply parallel plates powered by a power supply of several kilovolts in a working pressure of several tens to several hundreds of mTorr. The negative plate, also known as the cathode, was bombarded by ions from the plasma set up between these two plates, and cathode atoms we re dislodged from the metal surface. These atoms could then deposit on other surfaces inside the vacuum system, forming films.

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45 The principal type of system currently used for high-rate deposition of metals, alloys, and compounds is known as the magnetron cathode system. This type of tool uses magnetic confinement of electrons in the plasma, which results in a higher plasma density than in either the RF or DC diode systems. The higher plasma density reduces the discharge impedance and results in a much highercurrent, lower-voltage discharge. The electron confinement on a magnetron is due to the presence of orthogonal E and B fields at the cathode surface. This results in a classic E x B drift for electrons (the Hall effect), which gives rise to a sequence of cycloidal hopping steps parallel to the cathode face as shown in Figure 3-1. As a result, the secondary electrons which are emitted from the cathode because of ion bombardment are confined to the near vicinity of the cathode. In a magnetron, the electric field is always oriented normal to the surface of the cathode. The transverse magnetic field is configured so that the E x B drift paths form closed loops, in which the trapped, drifting electrons are constrained to circulate many times around the cathode face. A schematic of the DC magnetron sputtering deposition of Si used for this research is shown in Figure 3-2. A DC power source was used to sputte r the Cz-Si target, which was place 10cm below the substrate. A gas of Ar was supplied to the main chamber during deposition. Ar atoms were ionized by the electric field into Ar+ ions. Electrons from Ar ionization were accelerated and impacted other Ar atoms to form self-sustaining Ar+ plasma in the chamber. Ar+ ions bombarded the un-doped Cz-Si target to knock off atoms of Si from the target by momentum transfer. Finally, these Si atoms were deposited on the substrate. 3.1.2 Plasma Enhanced Chemical Vapor Deposition (PECVD) The STS310PC PECVD system was used for the deposition of silicon dioxide and amorphous Si films for excimer laser crystallization (ELC). The system was equipped with 13.56 MHz and 187.5 kHz frequencies and was capable of mixed frequency recipes. The temperature of the system was normally kept at 300 C. 2% silane gas was used for the deposition of a-Si.

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46 3.2 Thin Film Annealing 3.2.1 In-situ low pressure N2 Annealing Amorphous Si films deposited on SiO2/Si substrate were in-situ annealed in vacuum ambient with N2. A base pressure was 1.3x10-3 torr and working pressure was 1.2x10-1 torr with gas. A tungsten wire heater was placed under the substrate which was located in the center of the quartz tube as shown in Figure 3-3. Maximum temperature up to 1000C could be reached for this in-situ annealing with the ramping rate of 10C/min. Figure 3-4 shows the diagram of heat cycle of Si thin films during annealing in a conventional furnace. 3.2.2 Excimer Laser Annealing (ELA) System Laser annealing was perform ed with a Lambda Physik LPX 305 excimer laser operating with KrF (248nm) and pulse width ~25 ns (FWHM). The beam was collimated using a long focal length (200cm) lens. The beam was then focused using a spherical lens with a 25cm focal length. Distance from the lens to the sample holder was varied to control the energy density at the sample surface. Discharge voltage was also varied in the laser control system to alter the beam energy and hence energy density at the sample. Total energy was measured immediately after the focusing lens to incorporate energy loss through the lenses. Schematic diagrams of the laser setup are given for each experiment illustrated in Figure 3-5. 3.3 Analytical Techniques 3.3.1 X-ray Diffraction (XRD) X-ra y diffraction (XRD) is a non-destructive tool for analyzing material properties such as crystallinity, strain and grain size. Also, the phase identification and orientation can be easily analyzed with this technique. The incident X-ray interact with the periodic crystal planes and results in constructive and destructive interface. Only the condition satisfying the Braggs Law gives the constructive interference.

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47 Braggs Law: n = 2dsin Where n is the order of diffraction, is the x-ray wavelength, d is the interplanar spacing of crystal planes, and is the incident angle of x-ray. In this research Philips APD 3720 X-ra y and Philips Xpert MRD high resolution diffractometers were used. The Philips APD 3720 system using Cu K ( =1.5405) was employed for -2 scans to check for preferred orientation of crystalline and identify polycrystalline Si nanoparticles. The domain size can be obtained by using Scherrers formula from -2 scans: Scherrers formula: BB t cos 89.0 where t is thickness of crystallite, is the wavelength, B is FWHM (full width at half maximum), and B is the Bragg angle. However, the APD 3720 uses theta-2theta Bragg-rentano setup and this hinders a geometric freedom. Thus, the Philips Xpert was used to get further information from crystallized Si thin films. The Philips Xpert high resolution diffractometer uses 4circle goniometer and gives various setting like omega, phi and Chi scan (Figure 3-4). The glazing incident x-ray diffraction (GIXD) was used to inve stigate the crystallinity of the surface of Si thin films 3.3.2 Raman Spectroscopy Ram an spectroscopy is a spectroscopic technique used in condensed matter physic and chemistry to study vibrational, rotational, and other low-frequency modes in a system. It relies on inelastic scattering, or Raman scattering, of monochrom atic light, usually from a laser in the visible, near infrared, or near ultraviolet range. The laser light interacts with phonons or other excitations in the system resulting in the energy of the laser photons being shifted up or down. The shift in energy gives information about the phonon modes in the system. The Raman effect occurs when light impinges upon a molecule and interacts with the electron cloud of the bonds of that molecule. The incident photon excites the molecule into a virtual state. For the spontaneous Raman effect, the

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48 molecule will be excited from the ground state to a virtual energy state, and relax into a vibrational excited state, which generates Stokes Raman scattering. If the molecule was already in an elevated vibrational energy state, the Raman scattering is then called anti-Stokes Raman scattering as shown Figure 3-6. A change in the molecule polarization potential or amount of deformation of the electron cloud with respect to the vibrational coordinate is required for the molecule to exhibit the Raman effect. The amount of the polarizability change will determine the Raman scattering intensity, whereas the Raman shift is equal to the vibration al level that is involved. 3.3.3 Scanning Electron Microscopy (SEM) Field Em ission Scanning Electron Microscopy (FESEM) was used to examine the surface structure of the poly-Si thin films annealed by a conventional furnace and excimer laser. A JEOL JSM 6400 operating at 15kV and a working distance of 15mm was used to obtain the surface information. 3.3.4 Transmission Electron Microscopy (TEM) High Resolut ion Transmission Electron Microscopy (HRTEM) was dominantly carried out in order to study the microstructure of the crystallized Si thin films in this work. A JEOL 2010F tungsten filament 200kV TEM with point to point resolution of 2.3A was used to obtain the structural properties of the films. The preparation of the samples for TEM analysis depended on the desired geometry of sample observation. In general, there are two methods: plan view and cross-sectional view. In this dissertation, both methods were ca rried out to examine the structure information. For TEM sample preparation, the sample was cut into 3mm diameter discs. And then planview sample preparation began with mechanical polishing of the back of the sample with Alumina powder (size of ~100 m) until it was ~50 m thick with two types of sand papers after glued with anti-etching wax to a circle holder. The sample was carefully removed from the circle holder, and mounted onto Teflon flame cylinder with paraffin on edge of backside of the sample in order to make a sample electron transparent. The sample was then slowly etched with mixture of 25 % Hydrofluoric

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49 acid (HF) and 75% Nitric acid (HNO3) until a small hole appeared. After etching the sample was soaked into Heptane solution for a few hours in order to completely remove the residue of the protecting wax.

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50 Figure 3-1 Schematic of motions of electrons in crossed E and B fiel ds. The vertical electric field E is consistent with the presence of a cat hode located at the botto m of the figure. The magnetic field is oriented pe rpendicularly to the page

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51 Figure 3-2 DC magnetron sputtering system

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52 Figure 3-3 Schematic of in-sit u vacuum annealing system Figure 3-4 Heat cycle in a conve ntional furnace during annealing

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53 Figure 3-5 Schematic of setup for laser anneal of samples in seeded nucleation experiments

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54 Figure 3-6 Energy level diagram showing the states involved in Raman signal

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55 CHAPTER 4 NANOPARTICLE INDUCED CRYSTALLIZATION OF AMORPHOUS SILICON THIN FI LMS BY THE DC MAGNETRON SPUTTERING SYSTEM 4.1 Introduction Polycrystalline silicon has been widely investig ated for their applications in the fabrication of thin film transistors (TFTs) and solar cell devices.[73] For this application, large grain size and the high quality of polycrystalline silicon ar e required to achieve high electron-hole mobility, which is critical for better perf ormance.[74] Furthermore, low temperature processing, less than 600C, is highly desirable for compatibility with lo w cost substrate such as glass. In general, high quality polycrystalline sili con thin films can be obtained by the crystallizat ion of amorphous silicon using PECVD or sputtering system fo llowed by low temperature thermal annealing. Alternatively, poly-Si th in films can be deposited directly at below 600C, provided that low deposition rates are sufficiently employed.[75] However, the grain size of as-deposited polycrystalline silicon films is mu ch smaller than that of the crystallized silicon films. Moreover, the surface roughness associated with as-deposited polycrystalline Si films is generally higher than that of recrystallized silicon films.[76] Solid phase crystallization (SPC) is a widely used method because of not only its simplicity and low cost but also its ability to produce a smooth interface and excellent uniformity with a high reproducibility.[77] SPC process at 600C, how ever, usually requires a long annealing time of over 20-60hrs to complete the transformation to poly-Si films with large grain size, making it difficult for manufacturing.[77, 78] In addition, poly-Si films made in this process have a high density of in-grain defects due to low crystallization temperatures, which may deteriorate the electrical prope rties.[77-79] Thus, various methods have been employed to shorten the crystallization time as well as im prove the performance for the poly-Si based devices.[66, 79-81]

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56 Recently, various kinds of methods using seed templates have been employed to enhance the crystallization of amorphous si licon thin film. Crystallites embedded in a-Si layer,[82] metal nanocrystals embedded in a-Si layers,[83] and na nowires embedded in a-Si layers.[84] In fact, these synthetic methods have produced the large grains of poly-Si at low temperature below 600C, and reported the improved electron mobility and the controlled or ientation of silicon films. Nevertheless they sometimes induced so me drawbacks such as mechanical damages induced by stress, randomly or iented grains, and metal c ontaminations. High quality nanoparticles, in this respect, can act as inte resting nucleation seeds when imbedded in a-Si matrix during solid-phase crystallization. In this work, a new modification in the nucle ation step, which can lead to enhance the crystallization of amorphous si licon thin film, is proposed fo r the fabrication of photovoltaic silicon based devices. 4.2 Experimental The high crystalline silicon nanoparticles in a-Si matrix were in situ synthesized using a two-step scheme. First, polycrystalline Si nanoparticles were dispersed with ethanol using ultrasonication at a power of 220W for 1hr. LFSiO2 (low frequency silicon dioxide) layer was deposited on Si substrate by PECVD unde r the condition of 2% silane gas (SiH4, 400sccm) and nitrus (N2O, 1420sccm). The working pressure in the chamber was 550mTorr and the substrate temperature was maintained at 300C. The disper sed nanoparticle solution was uniformly spincoated on SiO2/Si substrate at 2000rpm for a few seconds And then Reactive Ion Etching (RIE) was carried out in order to remove native ox ide from the surface of nanoparticles under the condition of CHF3 (40sccm), O2 (10sccm) gas flow for few seconds. DC power of 100W whereas RF power of 200W was used.

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57 Subsequently, amorphous silicon thin film s were deposited by the DC magnetron sputtering system. The base pressure obtained in deposition chamber was 6.8x10-8 Pa. Nondoped Cz-Si target and the mixture of Ar and H2 gases were used for the deposition. A number of sets of samples were prepared with various substrate temperatures in the rage of 200-400C, for two RF powers 150 (low power) and 200W (high power) in order to optimize the condition for the deposition of a-Si. The sputtering pressure was 3x10-1 Pa. An 800nm thick amorphous Si was deposited on the spin-coated substrate (Nanoparticles/SiO2/Si). Finally, the Solid-phase crystallization (SPC ) at different temper atures for 1-20hrs was carried out in a conventional furnace under an air condition. 4.3 Results and Discussion Figure 4-1 shows the deposition ra te as a function of substrat e temperature for the two RF power conditions 150 and 200W, respectively. Tw o main parameters, which are the substrate temperature and RF power, can change directly the film growth mech anism. They control precursor mobility and surface reactions. Furtherm ore, the RF power also changes the energy of species impinging the substrate. In the low temperature region, the deposition rate increased with the substrate temperature, while the deposition rate decreased with the substrate temperature in the high temperature region. As shown in Fig. 4-1, the deposition rate is larger for films deposited with higher RF power in the low temperature region presumably due to an increase in electrons and energetic ions, which enhance both the ionization rate of argon atoms and the sputtering yield, and the subsequent Si flux at the growing surf ace. The films growth rate is governed by the flux of adatoms arriving at the su bstrate. This explains the increase in the deposition rate with increasing de position RF power as shown in Fig. 4-1 and 4-2. In the high temperature region, however, the deposition rate became larger in films deposited with lower RF power. The temperature corresponding to the maximum value of the deposition rate depends on

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58 the RF power. At 150W, the depos ition rate reached its maximum value at 350 C. However, at 200W, the temperature for the maximum deposition rate was lowered to 300 C. The reactivity at the growing surface is related to the hydrogen surface coverage. An increase of the substrate temperature leads to the reduction of the hydr ogen surface coverage. Thermally stimulated hydrogen desorption has been observed by infrar ed absorption spect roscopy (IRAS) in a temperature range up to 350 C.[85] Therefore, as shown in Figure 4-1, the substrate temperature Tmax corresponding to maximal deposition rates are 350 and 300 C for two used RF powers of 150 and 200W, respectively. Above these temperatur es, the deposition rate becomes controlled by the surface reaction rather th an by the incoming flux of adatoms. In this high temperature region two features may take place: the hydrogen desorption from the bulk and therefore a change of the surface reaction probability. At 350 C or above, hydrogen starts to diffuse from the a-Si:H bulk outward, which induces an incr ease in the hydrogen su rface coverage of the growing film and a reducti on of the deposition rate. Figure 4-2 shows the average deposition rate as a functi on of RF power. The average deposition rate gradually increased as RF power increased. Also, deposition rate for each sample with the average value and standa rd deviation is summarized in Table 1. This result is in good agreement to the work of Webb[84] and Savvide ss[86] where an increase in deposition rate is attributed to an increase in the density of Ar ions when the power is increased. For analysis, X-ray diffraction (XRD) was empl oyed after annealing the deposited films to measure the crystallinity of silicon films dependi ng on different temperature and annealing time. Figure 4-3 shows the variation of XRD intensity with substrate temperature for two different RF power conditions 150 and 200W after annealing at 700C for 10hrs. As the substrate temperature increased, the intensity gradua lly increased. As can be seen in Fig 4-3, the

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59 intensity reached its maximum value at a subs trate temperature equal to 300C for each RF power 150 and 200W, while decreased at 400C. Thus, the optimized conditions of substrate temperature of 300C and RF pow er of 200W was obtained to pr oduce high quality of poly-Si thin films in our study. Figure 4-4 shows the grazing incident X-ray diffraction (GIXD) prof iles of the silicon thin films before and after annealing in air at different temperatur es for 10hrs. No obvious diffraction peaks were observed before annea ling in both samples implying the as-deposited silicon thin films to be amorphous. Also, nanopa rticles embedded in amorphous matrix before annealing dont affect the intensity of XRD. Nanoparticles embedded samples, however, had significantly higher crystallinity under the same annealing co ndition compared with those crystallized conventionally illust rated in Fig 4-4. (b). Those fi gures show diffraction peaks at 28.44 and 47.30 representing (111) and (220) orientations of the cr ystallized Si thin film for both samples, with and without th e nanoparticle seeds, respectively, indicating the crystallization of the thin films has initiated. These peaks show that crystalline silicon has been formed during solid phase crystallization (SPC) and that the films are of polycrystalline phase. Figure 4-5 indicates the full-wi dth at half-maximum (FWHM) of the diffraction peaks at (111) as a function of annealing temperature. This shows the eff ect of nanoparticle seeds on the films crystallinity for various annealing temperatur es. It can be also observed from Fig. 4-5 that the nanoparticles embedded in a-Si matrix affected the crystallinity of th e poly-Si film. As the annealing temperatures increased, FWHM decreased and the peak intensity of the films increased gradually. It also shows that FWHM was effectively minimized in the sample with nanoparticle seeds after annealing. Poly-Si thin films can be fabricated with relatively good crystallinity by the nanoparticle induced crystallization (NIC) technique.

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60 The comparison between the thin films with and without the nanoparticles after solid phase crystallization (SPC) at 530C and 700C for 10hrs is shown in Fig. 4-6(a) and (b), respectively. Annealing temperature was increased from 450C w ith 20C intervals to c onfirm the effect of nanoparticles for the crystallization as nuc lei in amorphous matrix The sample with nanoparticles embedded in amorphous Si matrix has initiated the crystallization at 530C whereas the other showed no peaks at (111) and (220) or ientations as evident from Fig. 4-6(a). This shows the transition from amorphous to pol ycrystalline Si in samples with nanoparticle seeds occurred at lower temperature in our pr ocess. Also, the difference in intensity when annealed at 700C for 10hrs was ma ximized compared to that of the plain sample illustrated in Fig. 4-5(b). Figure 4-7 shows cross sectional TEM images of the crystallized Si thin film after annealing for 3hrs at 700C. The thickness of the film was about 800nm. Crystallites in the film were observed. Also, SAD patterns of the films indicate the film was of polycrystalline phase as shown in Fig 4-7 (d). Figure 4-8 shows the typical grai n size in films which were an nealed at 700C for 10hrs in air. In previous result, the difference of inte nsity between samples was maximized at 700C for 10hrs. In this condition, the av erage grain size of the nanopartic le induced sample (~300nm) was about 3 times larger than that of plain sample (~100nm). This indicate s the importance of our technique in terms of not only the processing time reduction but also the fabrication of high quality thin films with large grains. Thus, great improvement in both of the crystallization temperature and transition time can be simulta neously achieved by this seed based approach. 4.4 Summary We investigated the SPC growth behavior of un-doped a-Si thin film s prepared by the DC sputtering system after annealing in air when crys talline nanoparticles were directly employed as

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61 the seed nuclei. For the growth mechanism of a-Si:H films, the influence of substrate temperature is divided into two main processes. At first, in low temperature region the deposition rate is controlled by the flux of arriving adatoms from the plasma Secondly, at high temperature region the deposition rate is controlled by the surface reactions where the hydrogen surface coverage plays a key role. The optimum substrat e temperature and RF power to produce the a-Si thin film for the crystallization were obtained. High quality polycrystalline Si thin films were fabricated by the nanoparticle-induced crystalliz ation at relatively low temperature and short time. XRD and HRTEM analysis showed that our modification in the nucleation step promoted the crystal growth process by reduc ing the nucleation time, and then enhanced the crystallinity of the thin films. The measured grain size in the nanoparticle embedded sample was relatively higher compared to that of th e crystallized Si films without nanoparticle seeds. Based on our observations, we suggest that the uniformly disp ersed layer of high quality Si nanoparticles can be promising seed nuclei for the SPC process to achieve a high quality poly-Si thin film for photovoltaic Si based devices.

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62 Table 4-1 Deposition rate of all the samples w ith average value and standard deviation with different RF powers RF power (W) Deposition rate (/min) Average deposition rate (/min) Standard deviation Sample 1 Sample 2 Sample 3 100 24.5 24.9 25.1 24.8 0.3 150 26.9 26.1 26.5 26.5 0.4 200 30.1 29.4 30.5 30.0 0.6 250 34.9 34.8 35.9 35.2 0.7

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63 Figure 4-1 Variation of the depos ition rate as a function of subs trate temperature for two RF powers of 150 and 200W

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64 Figure 4-2 Average deposition rate as a function of RF power at substrate temperature of 300 C

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65 Figure 4-3 XRD intensity as a function of subs trate temperature for two RF powers of 150 and 200W after annealing

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66 Figure 4-4 GIXD patterns of silicon thin films: (a) samples without and (b) with nanoparticle seeds before and after annealing for 1 0hrs at different temperature in a air

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67 Figure 4-5 The full-width at half-maximum (FWHM) of the diffraction peak s at (111) orientation between samples with and without nanopart icle seeds after an nealing in a air

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68 Figure 4-6 GIXD pattern of a-Si thin films crystallized at (a) 530C, and (b) 700C for 10hrs on a plain and nanoparticle embedded substrate annealed in a air

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69 Figure 4-7 Cross section TEM images and SAD pattern s of Si thin films after annealing in a air for 10hrs at 700C

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70 Figure 4-8 Plan-view TEM micrographs of sample s (a) without and (b) with nanoparticle seeds after annealing in a air at 700C for 10hrs

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71 CAHPTER 5 POST ANNEALING EFFECTS OF AMORPHOUS SILICON THIN FILMS BY NANOPAR TICLE INDUCED CRYSTALLIZATION 5.1 Introduction Recently, the crystallization of amorphous silico n (a-Si) has been widely investigated for its application in the Si-based microelectronic devices.[87-89] However, for these applications, large-grained poly-Si is required to achieve hi gh electron-hole mobility, which is critical for increased performance.[74] Low temperature processing (<600C) is highly desirable for compatibility with low cost substrates such as glass. In general, high quality polycrystalline silicon thin films can be simply obtained by th e crystallization of am orphous silicon (a-Si) deposited using plasma-enhanced chemical vapor deposition (PECVD) or sputtering followed by low temperature thermal annealing. Alternatively, pol y-Si thin films can be also deposited directly below 600C, provided that deposition rates are sufficientl y low.[75] However, the grain size of this as-deposited polycryst alline silicon films is typically much smaller than that of the crystallized silicon films. Moreover, the surface roughness asso ciated with as-deposited polycrystalline Si films is gene rally higher than that of crystallized silicon films.[76] Various methods for crystallization of amorpho us Si have been widely employed.[22, 78, 79, 90-94] Solid phase crystallization (SPC) is wi dely used method of fa bricating po ly-Si films because of simplicity and low cost as well as th e ability to produce a smooth interface with the substrate, excellent uniformity, and high reprodu cibility.[80] However, SPC process at 600C usually requires a long anneali ng time of 20-60 hrs to complete the transformation to poly-Si films with large grain size.[80, 82] In addition, poly-Si films made in this process have a high density of defects in grains which may deteriorat e the electrical properties.[82, 83] Thus, various methods should be employed to shorten the crysta llization time and enhance the film quality as well as improve the performance of pol y-Si based devices.[20, 83, 95, 96]

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72 Recently, various kinds of processing techniqu es using seed templates have been employed to enhance the crystallization of a-Si thin f ilms. Specifically crystallites embedded in a-Si layer,[97] metal nanocrystals embedded in a-Si laye rs,[98] and nanowires embedded in a-Si layers[99] have been used to enhance the crys tallization. In fact, these methods have produced large-grained poly-Si at te mperatures below 600C with and reported improved electron mobility. Therefore, high quality nanoparticles, in this respect, can presumably act as nucleation seeds during solid-phase crystalliz ation (SPC) when directly imbe dded in an a-Si matrix. Also, for the poly-Si films by SPC, the surface oxygen suppr esses the migration of Si atoms on the a-Si surface.[100] Thus, low pressure N2 annealing was carried out to examine the effect of oxygen atoms at the surface and cr ystallization mechanism. In this work, a new modification in the nucleation step to enhance the crystallization of an a-Si thin film before annealing in low pressure N2 ambient is proposed for the fabrication of high-quality poly-Si films for photovoltaic devices. 5.2 Experimental Poly-Si nanoparticles in amorphous Si matrix were synthesize d in situ using a two-step scheme. First, polycrystalline Si nanoparticles were dispersed in ethanol using ultra-sonication at a power of 220W for 1hr. Next, a low frequency silicon dioxide (SiO2) layer was deposited on crystalline Si substrate by PECVD using 2% silane (SiH4) and nitrus gases (N2O) flowing at 400 and 1420 sccm, respectively. The working pressu re in the chamber was 550 mTorr and the substrate temperature was maintained at 300C. The dispersed nanopartic le solution was spincoated on SiO2/Si substrate at 2000rpm for a fe w seconds. Subsequently, reactive ion etching (RIE) for 10s using CHF3 and O2 gases flowing at 40 and 10sccm, respectively, was carried out to remove native oxide from the surface of nanopa rticles. DC power of 100W whereas Rf power of 200W was used for the etching.

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73 Subsequently, a-Si thin films were depos ited by the direct current (DC) magnetron sputtering system. The base pressure obtained in deposition chamber was 6.8x10-8 Pa, whereas sputtering pressure was 3x10-1 Pa. A non-doped Cz-Si target a nd the mixture of Ar and H2 gases were used for the deposition. RF power and DC voltage were 200W and 125V, respectively. An 800nm thick amorphous Si was deposited on th e spin-coated substrate at 300C. Finally, SPC process at different temperatures for 1-20hrs wa s carried out in a conventional furnace with low pressure (1.2x10-1 Torr) N2 after wet-etching of native oxide on the surface with HF solution. For analysis, X-ray diffraction (XRD) was used after annealing to measure the crystallinity and grain size of silicon films depending on various temperature and time. Also, the samples were analyzed by means of high resolution tran smission electron microscopy (HRTEM) in order to determine the structure, the crystallized vo lume fraction and growth rate of the films. 5.3 Results and Discussion Figure 5-1 shows the grazing incident x-ray di ffraction (GIXD) profiles of the silicon thin films before and after annea ling in low pressure (1.2x10-1 Torr) N2 ambient at different temperatures for 10hrs. No obvious diffraction peaks were observed before annealing in both samples which means the as-deposited silicon thin films to be amorphous and nanoparticles embedded in a-Si matrix didnt affect the cr ystalline phase as shown in Fig. 5-1(b). Nanoparticles embedded samples, however, had signi ficantly higher crystallinity after annealing compared with those crystallized conventionally Those figures show di ffraction peaks at 28.44 and 47.30 representing (111) and (220) preferred orientations of the crystallized Si thin film for both samples with and without th e nanoparticles, respec tively, indicating the crystallization of the thin films has initiated.

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74 Figure 5-2 shows the full-widt h at half-maximum (FWHM) of the (111) reflection in samples with and without spin-coated nanoparticles after annealing for diffe rent temperatures in low pressure (1.3x10-3 Torr) N2 ambient for 10h. Samples with nanoparticles showed smaller FWHM (higher crystallin ity) than those without nanoparticles. Comparison of the FHMW of the diffraction peaks at (111) orient ation between with and without nanoparticles when annealed in air or low-pressure N2 ambient is summarized in Table 5-1. From this data the nanoparticle embedded samples annealed in low pressure N2 ambient show the lowest FMHW intensity indicating the highest qual ity film. Based on this result, the optimization of annealing conditions was determined. Also, this allows investigating th e seed effect on the films crystallinity for the various annealing temperatures. It was believed that the nanoparticles embedded in a-Si matrix and annealing in low pressure N2 ambient affected the crystallinity of the poly-Si films. As the annealing temperature increased, FWHM decreased a nd the peak intensity of the films increased. Thus, this means that poly-Si thin films can be fabricated with rela tively good crystallinity by this nanoparticle induced crystallizati on (NIC) technique in low pressure N2 ambient. In general, the average grain size can be evaluated from the breadth of the X-ray Bragg diffraction peaks with Sc herrers formula.[101] cos 94.7 FWHM GSIZE (5-1) where GSIZE is the grain size (nm), FWHM (degree) is the full-width at half-maximum of diffraction peaks, is the diffraction Bragg angle. A nd Figure 5-3 shows the annealing time dependence of the average grain size for the thin films after SPC with and without the nanoparticle seeds. It can be seen that the average grain size increases fast first and then slow as the annealing time prolongs, and the average grain size for the films with the nanoparticle seeds increases much faster than t hose without the pretreatment fo r the same SPC annealing time,

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75 which indicates that the nanopartic le seeds can promote the crystall ization of the a-Si thin films during the SPC process. The structure of both nanopartic les and crystallized silicon thin films with and without nanoparticle seeds was extensively studied by hi gh resolution transmission electron microscopy (HRTEM). All samples were examined in a JEOL 2010F electron microscope, operating at the accelerating voltage of 200kV. Figure 5-4 and 5-5 show plan-view TEM images, XRD and SAD patterns of polycrystalline Si nanoparticles respectively. The size of nanoparticles was in the range of 15~20nm as shown in Figure 5-4. Also, the diffr action peaks at 28.44 and 44.30 representing (111) and (002) reflection of the poly-Si nanoparticle were eviden t illustrated in Figure 5-5(a). From the selected area diffraction (SAD) pa tterns, the silicon nanoparticles were of polycrystalline phase as shown in Figure 5-5( b). Also, spots in SAD patterns show the high quality of poly-Si nanoparticles. The crystallized sample with na noparticle seeds showed similar SAD patterns compared to those of polycrystal line nanoparticles illust rated in Figure 5-15, indicating the crystallization was induced by nanopa rticle seeds. Figure 5-6, 5-7 and 5-8 show th e plan-view TEM images of polycrystalline Si thin films without nanoparticle seeds afte r annealing in low pressure N2 ambient for 3, 10, and 20hrs at 700C, respectively. From the Fig. 5-6, some crys tallites, of which the size was less than ~20nm, were observed indicating the random nucleation wa s initiated. Also, the random growth of the films has occurred during the cr ystallization depending on anneali ng time. The average grain size was in the range of 150~1000nm as shown in Fig. 5-7 and 5-8, respectively. Figure 5-9 shows the SAD patterns of polySi samples without nanoparticles after annealing for 10, 20h at 700C, respectively. As th e annealing time increase d, the ring shape of

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76 SAD patterns became much clear, indicating the im provement of the crystallinity of the films depending on annealing time. Figure 5-10, 5-11, 5-12, 5-13 a nd 5-14 show the plan-view TEM images of poly-Si thin films after annealing in low pressure N2 at 550, 600, 650, 700 and 800 C for 10hrs, respectively. As annealing temperature increased, the grain size of samples with nanoparticle seeds increased. In general, the grain size of polycrystalline silicon thin films fabricated by SPC is finally decided by the competition between nucle ation rate and growth rate of crystallites,[102] in high temperature region, it will produce a lot of nuclei in the a-Si thin films which will inhibit the further grain growth process in the growth stage and reduce the grain size in the steady state. However, nanoparticle seeded samples weren t affected due to th e pre-existing nuclei. Figure 5-16 shows the experiment al incubation time as a function of temperature. As the thickness of the amorphous Si layer increases, th e incubation time becomes longer. Also, the incubation time decreases when the annealing temperature increases. The incubation times depends on nucleation and growth processes. If the annealing temperature increases, the nucleation and growth process becomes fast. The temperature dependency can be expressed by an Arrhenius relationship: kT QTexp0 (5-2) Where 0 is a pre-exponen tial factor and QT is the activation energy for the incubation time. From the slope of the Arrhenius curve in Fig. 5-16, the activation energy of 2.9 eV can be calculated. The activation energy is also useful thermodynamic parameter which can be used to estimate the interface energy and the Gibbs free energy for forming a critical nucleus. In this study, the crystalline fraction was define d as the ratio of crystallized parts to the whole area in the observed TEM images. Crys tallized volume fraction was obtained though

PAGE 77

77 several TEM images. Since the sizes of the observed grains were larger than the film thickness, the crystalline area fraction could be assumed to be the crystalline volume fraction. Generally, the annealing time dependence of the crystalline fraction is expr essed by Avramis equation[103, 104] exp11 0 m ct (5-3) Where is the crystalline volume fraction, t is transformation duration and 0 is a lag time for nucleation and c is a characteristic time. 33/1 2 drrns g c (5-4) where rg is the growth rate and rns is the steady-state nucleation rate. As discussed above, nanoparticles embedded in a-Si matrix can promote the nucleation and growth during solid phase crysta llization (SPC). In other words, the activation energy for nucleation is lowered. From the equation (5 -3) as transformation time (t) increases, the crystalline volume fraction ( ) increases. In our study, nucleat ion time is not required which results in the reduction of incubation time si nce nanoparticles can act as nuclei for the crystallization of a-Si thin f ilms. Crystallization from the surface would be due to the poly-Si surface having no oxygen atoms and Si atoms on surf ace having bigger free energies while in the case of samples annealed in air, the surface oxy gen suppresses the migration of Si atoms on the a-Si surface. This indicates that the migration of Si atoms is controlled by the surface oxide and pushes Ar atoms out of the films from the surf ace. Thus crystals grow on the surface and the diffusion of Ar is important.[100] Figure 5-17 shows the annealing time dependence of the crystallized volume fraction of the Si thin films at various annealing temper atures. From the figure, it can be seen that

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78 crystallization began after a certain incubation time, and a longer annealing time was required for the film to be completely crystallized as anne aling temperature decreased. It was found that the time to be fully crystallized was significantly reduced and the measured crystallized volume fraction of the Si films was relatively enhanced in the crystallized films induced by nanoparticles illustrated in Fig. 5-17(b). This indicates that nuc leation for the crystallization was successively modified by our approach, which effectively led to the reduction of the transition time from amorphous to polycrystalline phase. Figure 5-18 shows the behavior of the growth rate of poly-Si thin film depending on temperatures indicated. The lower growth rate in our experiments can arise from the small size of poly-Si nanoparticles and cr ystalline phase used as the seed nucl ei for the crystallization of a-Si. According to the thermally activated growth th eory,[105] the solid-phase growth velocity ( ) is given by )],/exp(1)[/exp(/ kTg kTEhkTca a )/exp(/ kTEhga ca when, kTgca (5-5) Where is the distance across the interface of crystalline and amorphous phases, Ea is the activation energy for the crystallization, and gca is the free energy difference between crystalline and amorphous phases. Here, Ea is usually interpreted as the energy involved in bond breaking or bond arrangement responsible for SPE growth and it can be obtained from the slope of the plot in Fig. 5-18. Thus, Ea = 3.2 eV which is close to the value reported by R. Iverson for epitaxial crystallization.[23] Equation (5-4) also indicates that the SPE growth velocity is proportional to the free energy difference between the crystallin e and amorphous phase gca. According to Gibb-Thomson effects,[106] the free energy of nanosized particle of phase faced with phase is raised by amount of 2 Vm/r, where is / the interfacial energy, Vm is the

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79 molar volume of and the r is the radius of the nanosi zed particle. This size effect can be applied to our study where the poly-Si nanoparticle s have average size of 15~20nm, leading to a larger free energy than that of a flat Si substrate. In addition to, grain defects in the polycrystalline phase of nanopartic le increases the free energy, resu lted in the reduction of the growth rate. As a result, the free energy differe nce between the crystalline and amorphous phase gca become smaller when c-Si nanoparticles are em ployed as the seed for the crystallization of a-Si, and the SPE growth velocity is decelerat ed according to Equation (5-5) compared to SPE with the single crystalline template. The crystallization is a comb ination of nucleation and grow th. Therefore, the median crystallization time, t50, in which 50% of a-Si has crystallized, was determined from the fitted volume fraction curve. The temperature dependence of t50 is shown in Figure 5-19. The activation energy for the crystalliz ation of samples with nanoparticle seeds was relatively lower than those without nanoparticle seeds. The ac tivation energy for the crystallization of the nanoparticle seeded sample is 2.6 eV. It is very close to the activation energy for the incubation time, which is 2.9 eV. Actually, the incubation time is related not only to the nucleation process, but also to the growth process. Theref ore, both activation energi es are supposed to be very close in value. In general, the solid phase crystallization ( SPC) of the amorphous Si thin films proceeds through four steps, including incubation, nuc leation, growth and steady state.[107, 108] Nanoparticle seeds can reduce the incubation time, the a-Si thin films with nanoparticle seeds start nucleation for short time, and fewer annealing time is needed for the crystals to achieve the same average grain size for the thin films with nanoparticle seeds than those without nanoparticles. In this study, crystalline nanopart icles can act as nuclei wh ich led to the reduction

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80 of the nucleation time and lower crystallization temperature due to lower activation energy after annealing in low pressure N2 ambient. Consequently, the crystallites induced by na noparticles grew much faster than those crystallized conventionally under the same a nnealing conditions, resulting in the higher crystallized volume fraction and growth rate. 5.4 Summary We investigated the SPC growth behavior of un-doped a-Si thin films prepared by the DC sputtering system when crystalline nanoparticles were directly employed as the seed nuclei after annealing with low pressure (1.2x10-1 Torr) N2 ambient. Relatively highe r quality polycrystalline Si thin films were fabricated by the nanoparticle-induced crystalliz ation (NIC) with low pressure N2 at relatively lower temperature and shorter tim e compared to those annealed in air. This indicates that crystallization from the surface would be due to the poly-Si surface having no oxygen atoms and Si atoms on surface having bigger fr ee energies while in the case of samples annealed in air, the surface oxygen suppresses the migration of Si atoms on the a-Si surface. XRD and HRTEM analysis showed that our modi fication in the nucleation step promotes the crystal growth process by reducing the incubation time, and then improves the crystallinity of the thin films. The activation energy for incubation was lower, while the measured crystallized volume fraction and growth rate in the nanopart icle embedded sample were relatively higher compared to those of the crystallized Si films without nanoparticle seeds. Based on our observations, we suggest that the uniformly dispersed layer of high quality Si nanoparticles can be promising seed nuclei for the SPC process to achieve a high quality poly-Si thin film for photovoltaic Si based devices.

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81 Table 5-1 Comparison of the full-width at half-maximum (FWHM) of the diffraction peaks at (111) orientation between the plain and na noparticle-induced samples when annealed in air or low pressure N2 ambient Temp. ( C) Plain sample Annealed with N2 Nanoparticle sample Annealed with N2 600 2.39 1.81 1.36 1.02 700 1.42 1.19 1.19 0.98 800 1.26 1.01 1.06 0.82

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82 Figure 5-1 GIXD patterns of silicon thin films: (a ) plain, (b) nanoparticle embedded films before and after annealing for 10hrs at different temperature under N2 ambient

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83 Figure 5-2 The full-width at half-maximum (FWHM) of the diffraction peak s at (111) orientation between the plain and nanoparticle-i nduced samples annealed in N2 ambient

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84 Figure 5-3 Average grain size as a function of th e annealing time for the SPC annealed samples with and without the nanopartcle seeds

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85 Figure 5-4 Plan-view TEM micrographs of polycrystalline Si nanoparticles

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86 Figure 5-5 XRD and SAD patterns of polycrystalline nanoparticle

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87 Figure 5-6 TEM images of plain samples without nanoparticle seeds after annealing low pressure N2 for 3hrs at 700C

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88 Figure 5-7 TEM images of samples without nanopart icle seeds after annealing in low pressure N2 for 10hrs at 700C

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89 Figure 5-8 TEM image of the sample without nanopa rticle seeds after annealing in low pressure N2 for 20hrs at 700C

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90 Figure 5-9 SAD patterns of samples without nanopa rticle seeds after annealing in low pressure N2 for 10 and 20hrs at 700C, respectively

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91 Figure 5-10 TEM images of nanoparticle embedded samples after annealing in low pressure N2 at 550C for 10hrs

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92 Figure 5-11 TEM images of nanoparticle embedded samples after annealing in low pressure N2 at 600C for 10hrs

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93 Figure 5-12 TEM images of nanoparticle embedded samples after annealing in low pressure N2 at 650C for 10hrs

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94 Figure 5-13 TEM images of nanoparticle embedded samples after annealing in low pressure N2 at 700C for 10hrs

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95 Figure 5-14 TEM images of nanoparticle embedded samples after annealing in low pressure N2 at 800C for 10hrs

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96 Figure 5-15 SAD patterns of nanoparticle embedde d samples after annealing in low pressure N2 for 10 and 20hrs at 700C, respectively

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97 Figure 5-16 Temperature dependency of the m easured incubation time for crystallization samples without and with nanoparticle seeds

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98 Figure 5-17 Crystalline volume fraction as a functi on of annealing time at various temperatures indicated

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99 Figure 5-18 Temperature dependence of the SPE gr owth velocity in amorphous Si thin film deposited on poly-Si nanoparticle seeds

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100 Figure 5-19 Temperature dependency of th e median crystallization time, t50

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101 CHAPTER 6 EXIMER LASER CRYSTALLIZATION OF TH E AMORPHOUS SIL ICON THIN FLMS INDUCED BY NANOPARTICLES 6. 1 Introduction Recently, excimer laser crystallization (ELC) ha s widely investigated to fabricate high performance polycrystalline silicon thin film transistor.[109-111] Excimer lasers typically operate in the ultraviolet and he nce photons are absorbed by the silicon thin films within a few nanometers of the surface. Melting and solidifying occur on a nanosecond time scale, often without affecting the underlying substrate. This technique en ables the use of inexpensive substrates, such as glass, which are highly preferab le for low cost, large-ar ea electronics devices. Excimer laser crystallization (ELC ) of amorphous silicon films is a well-established method for producing large grain polycrystalli ne silicon films on glass subs trates. Although the application has primarily been focused on TFTs, its adapta bility for photovoltaics (PV) has attracted considerable interest.[11 2-114] However, due to the random lo cation of the grain boundaries thin film transistors (TFTs) made in these films s how poor device-to-device uniformity.[93, 115] The necessity of sufficiently large m ono-crystalline islands, the fabrica tion of such devices, however, would require a method to control their location.[116, 117] With ELC, grains having a diameter exceeding the film thickness can be obtained due to the super-la teral growth (SLG) phenomenon.[12] SLG occurs very sensitively depending on laser power density which means difficult to control. This refers to the process in which the film is nearly completely melted such grains grow laterally from the few isolated solid portions remaining at the interface. In this study, by artificially controlling SLG from these solid seeds, nanoparticles embedded in a-Si matrix, the larg e grains is also easil y obtained without the sensitivity of laser power density.

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102 6. 2 Experimental The poly-Si nanoparticles in a-Si matrix were in situ synthe sized using a two-step scheme. First, polycrystalline Si nanopart icles were dispersed with etha nol using ultra-sonication at a power of 220W for 1hr. LFSiO2 (low frequency silicon dioxide) layer was deposited on glass substrate by the plasma enhanced chemical vapor deposition (PECVD) under the condition of 2% silane gas (SiH4, 400sccm) and nitrus (N2O, 1420sccm). The working pressure in the chamber was 550mTorr and the substrate temperat ure was maintained at 300 C. The dispersed nanoparticle solution was uni formly spin-coated on SiO2/Si substrate at 2000rpm for a few seconds. And then Reactive Ion Etching (RIE) was carried out to remove native oxide from the surface of nanoparticles unde r the condition of CHF3 (40sccm), O2 (10sccm) gas flow for 10seconds. DC power of 100W whereas Rf power of 200W was used. Subsequently, amorphous silicon thin films were de posited by PECVD. The base pressure obtained in deposition chamber was 8.4x10-1 mTorr. 2% silane (SiH4, 2000sccm) and Ar (390sccm) gases were used for the deposition. RF power (13.56 MHz) was 30W. The working pressure was 800 mTorr. A 200nm thick a-Si was deposited on the spin-coated substrate at 300C. For excimer laser crystallization (ELC), we have performed single shot irradiation of a-Si thin film (200 nm thick) deposited by PECVD. And the samples were dehydrated at 500C for few hours in a vacuum condition (1.3x10-3 Torr). Subsequently, laser annealing was performe d with a Lambda Physik LPX 305 excimer laser operating with KrF (248nm) and pulse wi dth ~25ns (FWHM). The beam was collimated using a long focal length (200cm) lens. The beam was then focused using a spherical lens with a 25cm focal length. Distance from the lens to the sample holder was varied to control the energy density at the sample surface. Laser processi ng of the seeds embedded samples was done. All

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103 samples were irradiated in air with a single pu lse for each sample. As mentioned previously, the laser beam impinges on the surface of the sample so the nanoparticle seeds in contact with the film interface may affect nuclea tion during the solidification. Gi ven this arrangement, it is worthy to note that complete melting of the film is necessary from the seeds to potentially affect nucleation in the molten film. Laser energy densi ties used for this set of samples ranged from 100 to 450mJ/cm2. Though higher energy densities were initially examined, it was found that values exceeding 500mJ/cm2 caused ablative loss of the film. 6. 3 Results and Discussion Figure 6-1 shows Raman spectra of the depos ited films as a function of laser energy density. No Raman peak is observed in the ca se of the as-deposited sample. Evidence of a crystalline phase started to show after an energy density of 200mJ/cm2. Only above 200mJ/cm2, a peak around 520cm-1 started to appear that indicates the initiation of crystallinity in the film. The crystalline peak became sharp after 320mJ/cm2. However, the peak occured at 518cm-1, which was lower by a few wave numbers than th e number observed for monocrystalline silicon. This observation has been attributed to phonon c onfinement,[118] possibly due to the presence of nanocrystals embedded in a-Si:H environment. Stress-induced effects are also reported to cause this behavior.[118] The Raman peak intensity increases up to 320mJ/cm2 and decreases at 350mJ/cm2. This decrease could be due to the in duced surface damage by high energy laser beam. Similar results of decreasing Raman intens ity peak due to higher-e nergy ion implantation in semiconducting films were reported by ma ny workers.[119] Also, amorphous phase was characterized by shoulder ar ound the Raman shift of 480cm-1.[120] We could not observe shoulders for more than that of 200mJ/cm2. We studied the laser crystallization of a-Si thin films having nanoparticles as seed layer on glass as a function of the laser energy density. Figure 6-2 illust rates the Raman spectrum of the

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104 samples as a function of laser energy. As the laser energy increased from 200 to 380mJ/cm2, Raman intensity also increased. Similar results were observed in samples without nanoparticle seeds. But, in the samples with nanoparticl e seeds, the amorphous phase could be observed around 480cm-1 in all samples. It indicates that the cr ystallization in samples with nanoparticle seeds did not fully complete at the lower energy. At same time, it is also observed that in theses samples the crystallization was possible at the energy levels as low as 200mJ/cm2 in contrast to the samples having no nanoparticle seeds, wher e higher energy is needed to initiate crystallization. Although the samples of (c) and (d ) were irradiated with the laser of different energy densities, both of them had the simila r crystalline volume fr action of 72% could be concluded that the laser energy of 290mJ/cm2 is the optimized laser crystallization energy for aSi with nanoparticle seeds. The crystalline volume fractions of samples were calculated from the integrated intensities of the Raman peaks, with Gaussian fits for the amorphous peak (Ia) and Lorentzian fits for the crystalline peaks (Ic). The calculation was done as proposed by Tsu et al. with crystalline volume fraction (Xc) given by Eq. (6-1), where is the ratio of the backscattering cross-sections amorphous and crystalline phases,[120] a c c CII I X )( (6-1) The selection of a value for is complex due to its dependency on absorption coefficient of amorphous and crystalline silicon.[ 121] Different energy densities re sult in different grain sizes, which further complicates the analysis, since the absorption coefficient changes too.[122] has been calculated to be between 0.8 and 0.9, the most widely used value being 0.8 for mixed phase silicon, especially for excimer laser crystallized silicon.[121, 123, 124] For this analysis too, was taken to be 0.8.

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105 Figure 6-3 shows the crystalline volume fraction of poly-Si f ilms for different laser energy density. The crystalline volume fraction increased as the laser energy density increased. Moreover, the crystalline volume fraction was maxi mized in samples with nanoparticle seeds. In case of the samples without nanoparticles, the crystalline volume frac tion was 66% at laser energy density of 320mJ/cm2. This means that only 66% region had the electrical-band structure of crystalline silicon and the residual 34% region had a band structure associated with disordered bonding states. The crystalline volum e fraction increased to 72% in samples with nanoparticle seeds. An important reason of the high crystall ine volume fraction close to 72% was the increase in the grain size. Figure 6-4, 6-5, 6-6 and 6-7 show the field emission scanning electron microscopy (FESEM) images of nanopartic les dispersed on the SiO2/glass and samples without nanoparticle seeds after Secco etching. A selective wet chemi cal etch method was used to examine the surface information. The solution is the so called Secco etchant, which is widely used to investigate defects in the semiconductor industry. The co mposition of the solution is as follows: HF + K2Cr2O7+H2O in a ratio of HF : H2O = 2 : 1 with 44g K2Cr2O7 dissolved in 1l of the of H2O. It etches defect on all surfaces. From Fig. 6-5 and 6-6 as laser energy density increased, the grain size of the films slightly increased Also, the super lateral growth (SLG) can be observed and the maximum grain size was 1~4 m as shown in Figure 6-7. Figure 6-8 and 6-9 show the FESEM images of nanoparticle seeded samples after Secco etching. As laser energy density increases, aver age grain size of films increases. Moreover, grains induced from nanoparticles were relatively larger than t hose with no nanoparticle seeds. In addition to, Figure 6-10 shows the grains induced by nanoparticle seeds at energy densities of 290 and 320mJ/cm2, respectively. The grain size of films didnt decrease in sample

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106 with nanoparticle seeds, while th e grain size of plain samples ra pidly decreased after SLG region due to the complete melting. This result indicates that the SLG can be easily controlled, even at the high energy density which led to the comp lete melting of the films, since poly-Si nanoparticles, which has higher melting point than that of amorphous Si phase, in amorphous matrix can act as nucleation seeds. Figure 6-11 and 6-12 show the AFM images of samples with and without nanoparticles after laser annealing at various laser energy densities. Substan tial mass transport occurs during the growth of the disk structures induced by th e lateral solidification.[ 125] This suggests that lateral growth plays a major role in the formati on of poly-Si grains even in the partial melting regime, as already reported for SP irradiated samples.[126] RMS comparison as a function of laser ener gy density in sample s with and without nanoparticle seeds is shown in Figure 6-13. Maxi mum values of RMS were observed at the laser energy density where SLG occurred in both samp les. Moreover, nanoparticle seeded samples have higher roughness compared to that of non-seeded samples which mean more applicable for photovoltaic device due to the cap ability of light trapping. Figure 6-14 shows average maximum grain size as a function of laser energy density for samples without and with nanoparticle seeds crysta llized by excimer laser irradiation. The grain size changes significantly in the range of energy dens ity between 250 and 320mJ/cm2, and has the maximum value at 290mJ/cm2 for both films, whereas the grai n sizes are slightly different, being 800nm and 650nm for the plain and nano particle seeded f ilms, respectively. Figure 6-15 and 6-16 show TEM images and SAD patterns of samples with nanoparticle seeds annealed at laser energy density of 290 and 320mJ/cm2, respectively. Single crystalline phase was observed at laser energy density of 290mJ/cm2 as shown in Figure 16(a).

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107 6.4 Summary In this study, the investigation of the ELC of a-Si films by PECVD was carried out to prepare large grains at a low process temperatur e for a glass substrate. We reported the excimer laser crystallization of a-Si f ilms deposited on glass with or without nanoparticle seeds. Nanoparticles can act as nucleati on seeds led to SLG. Laser with energy density starting from 200mJ/cm2 was employed to bring crystallization in a-Si films w ith or without nanoparticle seeds. Films did not fully crystallize at energy density of laser less than 260mJ/cm2. The value of maximum crystal volume fraction in samples with nanoparticles seeds was calculated as 72% at the high energy density of 290mJ/cm2. The presence of SLG in poly-Si thin films was confined to a very narrow energy density range, thus expl aining the difficulties to reveal SLG in such a material. However, nanoparticle induced poly-Si films was not relatively sensitive to the narrow range of laser energy density for SLG, and also showed the larger grain size of the films compared those of samples without nanoparticle seeds. High quality poly-Si thin films were successi vely obtained by the nanoparticle induced crystallization technique. This result shows low temperature polySi is promising for fabricating large grains on a glass substrate for la rge-area microelectronic applications.

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108 Figure 6-1 Raman spectroscopy of plain samples at various laser energy densities indicated

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109 Figure 6-2 Raman spectroscopy of nanoparticle embedded samples at different laser energy densities: (a) 200mJ/cm2 (b) 260mJ/cm2 (c) 290mJ/cm2 (d) 320mJ/cm2

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110 Figure 6-3 Volume fraction as a function of various laser ener gy densities with and without nanoparticle seeds

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111 Figure 6-4 SEM images of dispersed nanoparticles on the SiO2/glass substrate

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112 Figure 6-5 SEM images of plain samples anne aled at laser energy density of 230mJ/cm2

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113 Figure 6-6 SEM images of plain samples anne aled at laser energy density of 320mJ/cm2

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114 Figure 6-7 SEM image of the super lateral growth (SLG) annealed at laser energy density of 350mJ/cm2

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115 Figure 6-8 SEM images of nanoparticle embedded samples annealed at laser energy density of 200mJ/cm2

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116 Figure 6-9 SEM images of nanoparticle embedded samples annealed at laser energy density of 260mJ/cm2

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117 Figure 6-10 SEM images of Si thin films i rradiated by excimer laser at 290and 320mJ/cm2, respectively

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118 Figure 6-11 AFM images of samples without nanopa rticle seeds annealed at laser energy density of 230 and 350mJ/cm2

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119 Figure 6-12 AFM images of nanoparticle embedded samples annealed at laser energy density of 200 and 290mJ/cm2

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120 Figure 6-13 RMS as a function of laser energy dens ity in samples with and without nanoparticle seeds

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121 Figure 6-14 Grain size distribution as a function of various la ser energy densities with and without nanoparticle seeds

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122 Figure 6-15 TEM images of nanopar ticle seeded samples annealed at laser energy density of 290 and 320mJ/cm2

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123 Figure 6-16 SAD patterns of nanopar ticle seeded samples annealed at laser energy density of 290 and 320 mJ/cm2

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124 CHAPTER 7 CONCLUSION The research presented in this dissertation focused on the crystallization of a morphous silicon thin film embedded with nanoparticle seeds. Two different t ypes of crystallization techniques, which are solid phase crystallization (SPC) and excimer laser crystallization (ELC), were carried out to product high quality poly-Si thin films. Firstly, a-Si films deposited by the DC magnetron sputtering system were investig ated for solid phase crystallization (SPC). Nanoparticle seeded Si thin films were anneal ed in a conventional furnace to understand the behavior of the crystallization and obtain a hi gh quality films. For excimer laser crystallization (ELC), a-Si thin films with nanoparticle seeds were deposited by PECVD followed by the dehydration process in order to remove the na tive oxide on the surface of nanoparticles. Film properties were examined by various characterizati on tools. Relatively large improvement of the crystallization of a-Si films was observed. 7.1 Solid Phase Crystallization (SPC) of the Seeded Films We investigated the SPC growth behavior of undoped a-Si thin films prepared by the DC sputtering system when crystalline nanoparticles were directly employed as the seed nuclei. High quality polycrystalline Si thin films were fabric ated by the nanoparticle-induced crystallization at relatively low temperature and short time. XRD and HRTEM analysis showed that our modification in the nucleation step promotes the crystal growth pr ocess by reducing the nucleation time, and then improves the crystallinity of the thin films. The measured crystallized volume fraction in the nanoparticle-i nduced sample is relatively higher compared to that of the crystallized Si films without nanoparticles. Based on our observations, we suggest that the uniformly dispersed layer of high quality Si nanoparticles can be promising seed nuclei for th e SPC process to achieve a high quality poly-Si

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125 thin film for photovoltaic Si based devices. Also, nanoparticle seeds influences the formation of nucleation at the interface between a-Si matrix and nanoparticle. The interfacial energy of a nucleus at the contacted interface with the s eed is reduced by forming a low energy boundary with the seed like a special grain boundary be tween two grains. Some special misorientation between two contacted crystals can reduce the interfacial energy dramatically compared to random misorientation between them because the result is two crystals matched very well without dangling bonds. According to our results, th e preferred texture of grains resulting from the crystallization is <1 11>. Compared to this crystalliz ation, nanoparticle seed indeced crystallization has a much shorter lag time for the crystallization. Furthermore, its crystallinity is much better than that of the plain sample, as shown in the results. Therefore, this technique to crystallize amorphous Si thin films on insulators can meet simultaneously the demands which the LCD (Liquid Crystal Display) industries are desperately asking to improve the performance of thin film transistors (TFT) and to make a process compatible with a low melting temperature glass substrate. It can reduc e the thermal budget of the process by inducing crystallizat ion at lower annealing temperature for a much shorter period. In addition to that, it can provide much better crystallinity, which improves the electrical properties of TFT, such as high driving current low leakage current, and better oxide integrity. 7.2 Excimer Laser Crystallization (ELC) of the Seeded Films We reported the excimer laser crystallization (ELC) of a-Si films deposited on glass with or without nanoparticle seeds. Nanoparticles can act as nucleation seeds led to SLG. Laser with energy density starting from 200mJ/cm2 was employed to bring crystallization in a-Si films with or without nanoparticle seeds. Films did not fully crystallize at energy dens ity of laser less than 260mJ/cm2. The value of maximum crystal volume frac tion in samples with nanoparticles seeds was calculated as 72% at the high energy density of 290mJ/cm2. The presence of SLG in poly-Si

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126 thin films was confined to a very narrow energy density range, thus explai ning the difficulties to reveal SLG in such a material. However, nanopar ticle induced poly-Si fi lms was not relatively sensitive to the narrow range of laser energy density for SLG. Thus, the controlled the super lateral growth (SLG) can be obtained by excimer laser crystallization with nanoparticle seeds. Subsequently, high quality poly-Si thin films were successively obtained by the nanoparticle induced crystalliza tion technique. This result show s low temperature poly-Si is promising for fabricating large grains on a gl ass substrate for large-area microelectronic applications

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133 BIOGRAPHICAL SKETCH Taekon Kim was born on February 23, 1977, in Seoul, South Korea. After graduating from high school in 1995, he attended Hong-ik University and earned a bachelors degree in materials science and engineering in February 2002. In August 2004, he enrolled at the University of Florida in the Depa rtment of Materials Science and Engineering to pursue Ph.D under the advisement of Dr. Rajiv K. Singh. His main research involved crystallization and characteriza tion of amorphous Si thin films for solar cell. He is co-author of more than 10 journal and conference papers.