Optoelectronics Materials And Techniques P Predeep

Optoelectronics Materials And Techniques P Predeep

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  OPTOELECTRONICS - MATERIALS AND TECHNIQUES Editedby Padmanabhan Predeep
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  Optoelectronics - Materialsand Techniques Edited by Padmanabhan Predeep Published by InTech Janeza Trdine 9, 51000 Rijeka, Croatia Copyright © 2011 InTech All chapters are Open Access articles distributed under the Creative Commons Non Commercial Share Alike Attribution 3.0 license, which permits to copy, distribute, transmit, and adapt the work in any medium, so long as the original work is properly cited. After this work has been published by InTech, authors have the right to republish it, in whole or part, in any publication of which they are the author, and to make other personal use of the work. Any republication, referencing or personal use of the work must explicitly identify the original source. Statements and opinions expressed in the chapters are these of the individual contributors and not necessarily those of the editors or publisher. No responsibility is accepted for the accuracy of information contained in the published articles. The publisher assumes no responsibility for any damage or injury to persons or property arising out of the use of any materials, instructions, methods or ideas contained in the book. Publishing Process Manager Mirna Cvijic Technical Editor Teodora Smiljanic Cover Designer Jan Hyrat Image Copyright Triff, 2010. Used under license from Shutterstock.com First published September, 2011 Printed in Croatia A free online edition of this book is available at www.intechopen.com Additional hard copies can be obtained from orders@intechweb.org Optoelectronics - Materials and Techniques, Edited by Padmanabhan Predeep p. cm. ISBN 978-953-307-276-0
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  free online editionsof InTech Books and Journals can be found at www.intechopen.com
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  Contents Preface IX Part 1Inorganic Optoelectronic Materials 1 Chapter 1 Optoelectronic Properties of Amorphous Silicon the Role of Hydrogen: from Experiment to Modeling 3 Franco Gaspari Chapter 2 Silicon–Rich Silicon Oxide Thin Films Fabricated by Electro-Chemical Method 27 Pham Van Hoi, Do Thuy Chi, Bui Huy and Nguyen Thuy Van Chapter 3 Silicon Oxide (SiOx, 0<x<2): a Challenging Material for Optoelectronics 55 Nicolae Tomozeiu Chapter 4 Gallium Nitride: An Overview of Structural Defects 99 Fong Kwong Yam, Li Li Low, Sue Ann Oh, and Zainuriah Hassan Chapter 5 Cuprous Oxide (Cu2O): A Unique System Hosting Various Excitonic Matter and Exhibiting Large Third-Order Nonlinear Optical Responses 137 Joon I. Jang Chapter 6 Optoelectronic Properties of ZnSe, ITO, TiO2 and ZnO Thin Films 165 S. Venkatachalam, H. Nanjo, K. Kawasaki, H. Hayashi, T. Ebina and D. Mangalaraj Part 2 Polymer Optoelectronic Materials 185 Chapter 7 Side-Chain Multifunctional Photoresponsive Polymeric Materials 187 Luigi Angiolini, Tiziana Benelli, Loris Giorgini, Attilio Golemme, Elisabetta Salatelli and Roberto Termine
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  VI Contents Chapter 8Ladder Polysiloxanes for Optoelectronic Applications 211 Zhongjie Ren, Shouke Yan and Rongben Zhang Chapter 9 Synthesis of Aromatic-Ring-Layered Polymers 235 Yasuhiro Morisaki and Yoshiki Chujo Chapter 10 Nanomorphologies in Conjugated Polymer Solutions and Films for Application in Optoelectronics, Resolved by Multiscale Computation 261 Cheng K. Lee and Chi C. Hua Part 3 Techniques and Characterization 285 Chapter 11 Optoelectronic Techniques for Surface Characterization of Fabrics 287 Michel Tourlonias, Marie-Ange Bueno and Laurent Bigué Chapter 12 Optoelectronic Circuits for Control of Lightwaves and Microwaves 313 Takahide Sakamoto Chapter 13 An Analytical Solution for Inhomogeneous Strain Fields Within Wurtzite GaN Cylinders Under Compression Test 337 X. X. Wei Chapter 14 Application of Quaternary AlInGaN- Based Alloys for Light Emission Devices 355 Sara C. P. Rodrigues, Guilherme M. Sipahi, Luísa Scolfaro and Eronides F. da Silva Jr. Chapter 15 Air Exposure Improvement of Optical Properties of Hydrogenated Nanostructured Silicon Thin Films for Optoelectronic Application 375 Atif Mossad Ali Chapter 16 Fabrication and Characterization of As Doped p-Type ZnO Films Grown by Magnetron Sputtering 393 J.C. Fan, C.C. Ling and Z. Xie Chapter 17 Light Intensity Fluctuations and Blueshift 421 Moon Kyu Choi Chapter 18 Self-Similarity in Semiconductors: Electronic and Optical Properties 435 L. M. Gaggero-Sager, E. Pujals, D. S. Díaz-Guerrero and J. Escorcia-García
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  Contents VII Chapter 19Long-Term Convergence of Bulk- and Nano-Crystal Properties 459 Sergei L. Pyshkin and John Ballato Chapter 20 Micro-Raman Studies of Li Doped and Undoped ZnO Needle Crystals 477 R. Jothilakshmi
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  To my father;but for his unrelenting efforts I would not have made it to this day.
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  Preface Optoelectronics - Materialsand Techniques is the first part of an edited anthology on the multifaceted areas of optoelectronics contributed by a selected group of authors including promising novices to experts in the field, where are discussed related materials and techniques. Photonics and optoelectronics are making an impact multiple times the semiconductor revolution made on the quality of our life. In telecommunication, entertainment devices, computational techniques, clean energy harvesting, medical instrumentation, materials and device characterization and scores of other areas of R&D the science of optics and electronics get coupled by fine technology advances to make incredibly large strides. The technology of light has advanced to a stage where disciplines sans boundaries are finding it indispensable. In this context this book would be of importance to researchers from materials scientists to device designers and fabricators. Photonics is to optics like electronics is to electricity. Photonics sculpts light like a sculptor does with granite. Light is beings squeezed, cut into the pieces, reconstructed back and the like. Currently optics is undergoing revolutionary changes and photonics is going to be the next centuries’ technology. Globally, countries are vying with each other in formulating their technology initiatives so as to ensure that they should not miss the “Photonics Bus” as many of them missed the semiconductor revolution in the last century. Data transfer and communication technology are going to unimaginable heights by the idea of photonic crystals - the idea optical scientists copied from mother nature’s work in nanotechnology in blooming beautiful colors and patterns on objects of desire like butterfly wings and peacock feathers With the emergence of photonics and laser technology, optoelectronics seems to be losing its identity and is often mixed up with photonics. Photonics draws from and contributes to several other fields, such as quantum electronics and modern optics. In this era of great mix up of disciplines and multidisciplinary research, it is not surprising that such mix of closely connected players like electrons and photons refuses to be confined to narrow boundaries of sub disciplines. Naturally the articles in this anthology also have their boundaries diffused over the diverse optical phenomena of optoelectronics and photonics. Readers are advised to bear this in mind when looking for titles of this book.
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  X Preface I amproud to present this collection of carefully selected peer reviewed high quality articles on various optoelectronic and photonic materials and techniques and would like to thank to the authors for their wonderful efforts. Stake holders of the ongoing optoelectronic and photonics revolution such as researchers, academics and scientists are sure to find this collection of essays enormously useful. July 2011 P. Predeep Professor Laboratory for Unconventional Electronics & Photonics Department of Physics National Institute of Technology Calicut India
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  Part 1 Inorganic OptoelectronicMaterials
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  1 Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling Franco Gaspari University of Ontario Institute of Technology, Canada 1. Introduction Amorphous silicon, and its more useful alloy form, hydrogenated amorphous silicon (a-Si:H), has been the subject of investigation for more than three decades. A-Si:H is a low- cost, efficient material which is used extensively for electronic devices. Indeed, most recent electronic device textbooks contain a comprehensive review of the physics of amorphous materials and amorphous silicon in particular (Baranovski, 2006; Kasap, 2005; Street, 2000). The advantages of a-Si:H are particularly evident when considering the photovoltaic application context for the preparation of solar cells: in fact, a-Si:H has a large optical absorption coefficient (about 0.5 micron of the material will absorb 90% of the incident sunlight); the energy gap can be modulated to allow for near optimum conversion efficiency for sunlight; it can be alloyed with other elements (carbon, germanium) to create multi- junction structures with increased energy conversion efficiency for sunlight. Finally, it is plentiful and can be deposited on a variety of materials (at low temperature, over large areas, and on flexible substrates). However, the presence of metastable defects in a-Si:H adversely affects the performance of photovoltaic cells and thin film transistors. Electrical conductivity, photoconductivity and luminescence degradation have been linked to defect formation, such as dangling bonds (DBs) in the a-Si:H film (Akkaya & Aktas, 1995; Street, 1980). Staebler and Wronski (1977) found that defects can be created by illuminating a-Si:H. The creation of these light-induced defects (LID) is therefore referred to as the Staebler-Wronski (SW) effect. The presence of these defects, or dangling bonds, is the major factor responsible for the deterioration of the optical and electronic properties of a-Si:H. On the other hand, these defects are metastable and can be cured. Indeed, we could define a SW process that can be described as a two-step reversible process: i. Exposure to sunlight leads to an increase in the density of states (dangling bonds) in the energy gap of a-Si:H; this represents the SW effect proper; ii. Subsequent annealing at elevated temperatures (150-200 OC) reduces the density of states back to the original value, thus restoring the optoelectronic properties. It has been shown experimentally that both optical and electronic properties of amorphous silicon, such as refractive index, optical gap, absorption coefficient, electron and hole
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  Optoelectronics - Materialsand Techniques 4 mobility, etc., are strongly dependent on hydrogen content, in terms of both hydrogen concentration and hydrogen dynamics (diffusion) under various conditions - see, for instance, (Searle, 1998) and references therein. The investigation of such dynamics, including the relation with defect creation and annealing, is crucial for assessing the appropriate solutions to achieve better control of the defects and, consequently, better optoelectronic performances. There exists a large amount of articles and review papers or books that address the basic properties of a-Si:H, including analysis of the structural, optical and electronic properties; description of a variety of experimental methods used for the growth of a-Si:H films; and correlation between growth parameters and film quality. In this chapter a summary of the basic properties and historical issues related to a-Si:H and its applications in optoelectronics is presented in section 2. A more exhaustive description of the basic properties of a-Si:H is provided by the references in this section. Section 3 will focus on the role of hydrogen in relation to the optoelectronic properties and defect dynamics in a-Si:H, and will examine some of the prominent models of hydrogen diffusion also used to describe the SW process dynamics. Section 4 will describe the use of tritium, an isotope of hydrogen, as an experimental probe that can be used as a reference by such models. Finally, section 5 will present the results of an integrated experimental and theoretical approach aimed at developing a proper model of the dynamics inherent to a-Si:H and amorphous materials in general. Future work necessary to achieve a proper description of these dynamic processes will be indicated in the Conclusion section. 2. Properties of a-Si:H There exist several preparation methods for a-Si:H films. Early work on evaporated and sputtered a-Si:H lead to poor quality films, and it is now widely accepted that Radio Frequency (RF) Glow Discharge produces the best quality material, although other more recent methods claim similar or better results. A comprehensive review of the advantages and disadvantages of the different methods employed to grow a-Si:H can be found in the books edited by Searle (1998) and Street (1991). In general, it is desirable that a hydrogen plasma be employed to help the formation of Si— Hn ion radicals; hence, methods based on plasma-enhanced chemical vapour deposition (PECVD) techniques are usually preferred. The ions produced in the plasma region are directed via an electric field towards a substrate, where film growth takes place. A common characteristic of these PECVD techniques is the possibility of tuning the system using several parameters, which might be mutually dependent on or independent of each other, like partial gas pressure, electrode bias, substrate bias, flow rates, gas mixtures, substrate temperature, and any other adjustable parameter. A review of plasma deposition of a-Si:H can also be found in (Bruno et al.,1995). If the goal of current research in this sector is the understanding and prediction of the properties of a-S:H, it is crucial that the dependence of physical properties on preparation conditions be fully examined. This requires the development of experimental and predictive tools applicable to size scales ranging from the atomic to the macroscopic levels. Both Searle (1998) and Street (1991) provide an exhaustive review of the structural, optical and electronic properties of a-Si:H, and point out the still unresolved issues. In the following subsection, the basic properties of a-Si:H are presented, with a focus on the role of hydrogen.
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  Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling 5 2.1 Structure and Density of States (DOS) In order to understand the implication of the amorphous structure of a-Si:H on its opto- electronic properties, it is useful to examine the structure of amorphous silicon in comparison to its crystalline form (c-Si). Crystalline silicon is characterized by the well known diamond (or tetrahedral) structure, with bond length of 23.3 nm and bond angle of 109.5o. As a matter of fact, the amorphous form shows very small changes from the crystalline parameters, with a ± 10% deviation in bond length, and a ± 5% deviation in bond angle. These small changes make it possible to maintain a relatively good short range order (within the first 2-3 nearest neighbours); however, the accumulation of structural stress, due to the progressive compounding of small deviations, eventually leads to bond breaking and the appearance of dangling bonds. Figure 1 shows simple 2-d schematics of the formation of dangling bonds: a 2-d square crystal (1a) is slightly distorted (1b, top, center atom). The distortions become more marked as the network is extended, and eventually a dangling bond (DB) appears to relieve the structural stress (top right quadrant of figure 1c: this is usually also accompanied by under-coordinated and over-coordinated bonds). Fig. 1. (a to c) 2-d schematics of formation of dangling bonds due to long range disorder. The negative effects of the dangling bonds on the opto-electronic properties of a-Si can be effectively removed by hydrogenation; that is, hydrogen atoms are introduced to passivate (bond to) the dangling bonds; see, for instance, (Kasap, 2005; Street 1991, 2000). Hydrogen atoms incorporated into the films satisfy the covalent bonds at defects and microvoids and also allow the lattice to relax, thereby reducing the density of localized states by several orders of magnitude. Figure 2(b) show a 3-d representation of amorphous silicon with dangling bonds passivated by hydrogen atoms. A crystalline structure is also shown for comparison in figure 2(a). The differences and similarities between the crystalline silicon and amorphous silicon structures are evident when we examine the radial distribution functions (RDF) for the two structures, as shown in Figure 3. The amorphous structure still shows ordered, crystalline features for the first 3 nearest neighbors. The first neighbor also maintains the crystalline sharpness for the peak, while the progressive deviations from the crystalline structure are evident in the spreading of the peaks for the second and third nearest neighbor.
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  Optoelectronics - Materialsand Techniques 6 (a) (b) Fig. 2. A 3-d computer model representation of c-Si (a), and a-Si (b) with dangling bonds passivated by hydrogen atoms (red balls). The role of hydrogen in determining the degree of disorder is also the subject of numerous studies. For instance, O’Leary et al. (1996), by using optical absorption data, and by investigating how the modeling parameters vary with the bonded hydrogen concentration, suggest that bonded hydrogen helps decreasing the amount of disorder, and has an impact on the optical absorption spectrum. More recently, Ukpong (2007) studied the chemically-induced disorder-to-order transition in hydrogenated amorphous silicon as a function of hydrogen concentration, CH. The author identifies three stages, associated with low CH, medium CH, and high CH, that describe the changes in the stress and structure parameters. Rui et al (2005) investigated the effect of hydrogen plasma annealing on the micro-structural transition from disorder to order in amorphous silicon films. They found that there exist two steps for the reaction between atomic hydrogen and Si network, and show that the hydrogen plasma treatment conditions strongly influence the microstructures of the amorphous Si films The disorder inherent in the amorphous structure and the presence of dangling bonds has a crucial impact also on the electronic density of states (DOS) of amorphous silicon. Figure 4 shows a simple schematic representation of the electronic DOS of a-Si:H. Fig. 3. Radial Distribution Function of crystalline silicon (left) and amorphous silicon (right) [From: Laaziri et al., 1999]
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  Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling 7 Fig. 4. A schematic representation of the electronic density of states, g(E), of a-Si:H. VB indicates the valence band and CB the conduction band. The dashed, red, vertical lines show the mobility edges, which are defined as the energy level separating extended (non- localized) states from localized states. The main features in Figure 4 can be summarized as follows: i. The localized tail states are a consequence of the disorder inherent to the amorphous structure. Several studies have examined the role and the extent of tails states in a-Si:H. In particular, a characteristic energy, EU, or Urbach energy, gives the measure of the width of the tail states; hence, it is also referred to as the Urbach tail width - see, for instance, (Ley, 1983). The characteristic width for the band tail states is about 50 meV for the valence band tail states and about 25 meV for the conduction band tail states – see, for instance, (Cody, 1981) and the relevant articles cited in the reviews in (Street, 1991) and (Searle, 1998). Furthermore, the tail states width has been associated with a “degree” of disorder, with the implication that the optoelectronic properties of a-Si:H are also dependent on its value. ii. The localized defect states in the middle of the gap are associated with the formation of DBs. Different models have been proposed to identify the percentage and the energy levels of neutral DBs vs. the positive and negative DBs. Indeed, a DB is identified not only by the fact that the bond is unsatisfied, but also by its net charge, which is determined by the number of electrons sharing the dangling bond, i.e., no electrons imply a positive DB (D+), a single electron makes the bond a neutral one (D0), while the presence of two electrons lead to negative DBs (D-). One of the most interesting and
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  Optoelectronics - Materialsand Techniques 8 utilized models, describing the energy distribution of the three types of defects, is the so-called defect-pool model (Powell & Deane, 1996). iii. The localized states in the band tails become delocalized at a critical boundary called the mobility edge. A mobility gap is then defined as the energy separation between the two mobility edges of the conduction and valence bands. 2.2 Optical properties of a-Si:H A-Si:H can be described as a direct band-gap semiconductor. The original study of Tauc et al. (1966), in which the distributions of electronic states are assumed to be exactly square- root in character, terminating abruptly at the respective band edges, leads to a simple analysis of optical absorption and luminescence experiments. Optical absorption and luminescence occur by transition of electrons and holes between electronic states such as conduction and valence bands, tail states, and gap states. Tauc’s relation (Tauc, 1966) describes the dependence of the optical absorption constant, α, on the energy gap as: ( )2 G α ω B ω E = − = = (1) Where B is a constant, ħω is the photon energy and EG is the optical gap. The empirical determination of the optical gap EG can then be achieved by plotting α ω = vs. ω = , which is known as Tauc’s plot (a schematic illustration of Tauc’s plot is shown in Figure 5). However, the presence of localized tail states extending from the conduction and valence bands into the energy gap makes the determination of an optical gap unclear. For instance, Malik & O’Leary (2004) and Thevaril & O’Leary (2010) have addressed the fact that in amorphous semiconductors considerable deviations from square-root distributions of electronic states occur. They claim that the presence of tail states introduces a corresponding tail in the imaginary part of the dielectric function, ( ) 2 ω ε = , which makes the optical gap difficult to determine, i.e., it introduces a considerable amount of uncertainty into the Tauc optical gap determination procedure. Nevertheless, it is still a common procedure to determine the optical gap by using Tauc’s relation, although two different methods have been used to obtain a value for the gap. The first simply extrapolates the high energy, linear section of the plot of α ω = vs. ω = , and takes the intercept with the x-axis as the value of the optical gap, as shown in Figure 5. The second chooses the photon energy at which the absorption coefficient is equal to 104 cm−1, defined as E04, as the optical gap. The characteristic values for the band gap of a-Si:H determined from Tauc’s plot range from ~1.7 eV to ~1.9 eV. The variations in gap value are due to preparation conditions, but it is well accepted that the main parameter responsible for the value of the optical gap is the hydrogen content (CH). Indeed, there are numerous studies that have investigated the dependence of the optical gap and other optical parameters, like absorption coefficient and refractive index, on CH. Earlier studies can be found in the references in (Searle, 1998) and (Street, 1991). In summary, it has been shown that the optical band-gap of a-Si:H tends to increase with hydrogen content; see also, for instance, (Daouahi et al., 2001; Gaspari et al, 1993).
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  Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling 9 Fig. 5. A schematic illustration of a Tauc’s plot. The extrapolation of the high energy linear portion is used to determine the optical gap EG 2.3 Electronic properties of a-Si:H A-Si:H electronic properties also exhibit a strong dependence on the hydrogen bonding and content. For example, dark conductivity in a-Si:H can be described by two main processes. The first is the standard extended states conduction process, described by the relation (Mott, 1983) A 0 B E σ=σ exp k T ⎛ ⎞ − ⎜ ⎟ ⎝ ⎠ (2) where σ and σ0 are the electrical conductivity and a prefactor, respectively, and EA, kB and T are the activation energy, the Boltzmann constant and the temperature, respectively. EA is given by either EC− EF or EF− EV, depending on whether electrons or holes are considered, with EC and EV being the conduction band and valence band edges respectively. The second conduction process is referred to as variable-range hopping (VRH) conduction, a well known process in amorphous materials in general. This conduction process is associated with hopping within tail states, and is characterized by the following temperature dependence (Mott, 1983): h h0 1 4 B σ =σ exp T ⎛ ⎞ − ⎜ ⎟ ⎜ ⎟ ⎝ ⎠ (3) where σh and σh0 are the electrical conductivity and a prefactor, respectively, for variable range hopping. It has been shown that tail states are also subject to the SW effect (Longeaud et al., 2000). The authors also state that passivation of these DBs in the tail states is related to hydrogen reservoirs. As mentioned before, hydrogen will also influence tail states by reducing the amount of disorder, and by relaxing the structure; furthermore, defects states will tend to shift the Fermi level, thus influencing the activation energy in the dark conductivity. Therefore the hydrogen content plays a fundamental role in determining the conduction processes, as it does for the optical gap.
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  Optoelectronics - Materialsand Techniques 10 3. The role of hydrogen in a-Si:H As previously indicated, the presence of hydrogen during the growth of a-Si:H has a dramatic effect on the optoelectronic properties of a-Si. It is a well established fact that the presence of hydrogen atoms reduces the DB density, both at the center of the gap and in the tail states, thus reducing also the EU values. Furthermore, the optical gap increases with hydrogen content. However, several questions are still unanswered: for instance, it is still unclear whether there is a direct relation between hydrogen-content and optical gap, or whether such increase is also due to a structural reordering, leading to a less disordered structure, as indicated by the Urbach width. More importantly, the role of hydrogen dynamics during the defect creation and defect passivation phases of the SW process is still a matter of debate. The nature of hydrogen bonds, the hydrogen distribution, and hydrogen mobility represent crucial parameters in addressing these issues. 3.1 Hydrogen bonding The most effective characterization of hydrogen content and hydrogen bonding is provided by the vibrational density of states (VDOS), obtained experimentally via transmission and Raman infrared spectroscopy. Fourier Transform Infrared Spectroscopy (FTIR) has become in fact one of the routine modes of investigation to determine the quality of the a-Si:H film (Searle, 1998; Street, 1991; and references therein). Investigations on the correct interpretation of crucial features in the infrared (IR) spectrum, such as the nature of the stretching modes at about 2000 cm-1, the roles of chains and microvoids, the distinction among different poly-hydride bonds - i.e. Si—H2 vs. Si—H3 vs. (Si—Hn)m - became crucial in order to achieve a better understanding of the role of hydrogen atoms both in the determination of the basic film properties (energy gap, Fermi level, etc.) and in the dynamics of creation and annealing of defects. For instance, early infrared spectroscopy (Jeffrey et al. 1979; Knights & Lujan, 1979; Zanzucchi et al., 1977), primarily of evaporated and sputtered a-Si:H, associated poly- hydride bonding with poor film properties, but Street & Tsai (1988) and Kato & Aoki (1985) showed that that was not the case. A model predicting the various modes of vibration for silicon and hydrogen atoms in a-Si:H was developed by Lucovski et al. (1989). Recently, the correct interpretation of the various modes, in particular the stretching modes between 1950 and 2150 cm-1, has been questioned (Smets & van de Sanden, 2007); however the frequency assignments by Lucovski still provide an excellent reference for the investigation of a-Si:H. 3.2 Hydrogen diffusion models Several models have been proposed to describe the dynamics of hydrogen diffusion within the amorphous silicon network relative to the Staebler-Wronski effect. Furthermore, many techniques have been employed to generate a realistic computational model of a-Si:H. In particular, molecular dynamics (MD) has become one of the more powerful and frequently used tools for the correlation of the microscopic characteristic of materials with their macroscopic properties, observed experimentally. In order to underline the variety of models and approaches used to analyze hydrogen diffusion, a summary of some of the most
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  Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling 11 important models and experimental studies introduced over the past 20 years is presented below. It should be noted that the following summary represents only a fraction of the publications on this subject, and it is not intended to be exhaustive, but rather to provide a sense of the diversity in the approaches to the problem. Santos et al. (1991, 1993) present first experimental evidence of light-induced hydrogen motion in undoped a-Si:H, obtained from diffusion experiments under illumination. A definite increase in diffusion was observed for the illuminated samples. The authors speculate that the recombination of e-h pairs releases energy and may induce excitation of hydrogen (H) from a Si—H bond. Another important conclusion is that there is an electronic nature to hydrogen motion in a-Si:H. Up to at least 275 oC, the H diffusion is not purely a thermal process but is dominated by the concentration of carriers. Jackson & Tsai (1992) consider hydrogen bonding in terms of a density of states. Bonding in a given configuration is equivalent to occupancy of the state. The barriers to configuration changes are equated with the energy required to reach transport energy. The main conclusions are that there is a range of possibilities: one extreme is the case in which hydrogen is predominantly bonded on void surfaces and the transport energy is substantially different in a-Si than in c-Si; the other extreme is that hydrogen predominantly resides in platelets structures and the transport energy is similar to c-Si. The actual case depends on deposition conditions. Also, Jackson et al. (1993) show that, at relatively high hydrogen concentration, hydrogen atoms reside mainly in clusters. The energy of the clusters depends on the number of hydrogen pairs within the cluster. Annealing has the effect to shift the hydrogen in more stable clusters. Van de Walle & Street (1994, 1995) investigate, using first principle pseudo-potential calculations, the bonding energetics and the diffusion mechanism of Si—H bonds in general and in amorphous silicon. The main conclusions are as follows: i- it is favorable for a hydrogen atom to move from a DB site to a bond centered (BC) site (bond-centered between two silicon atoms); ii- not only is this a favorable path, but the energy levels which are introduced into the band-gap open the way for carrier-enhanced dissociation; iii- the main path is that by which H stays at approximately the same distance from the original silicon atom, i.e., it moves along the direction of its wagging mode into a BC site. Finally, the motion of hydrogen atoms through a-Si can be described by a diffusion coefficient: ( ) H 0 A D D exp E /kT = − (4) Where EA is the activation energy of the diffusion process. Biswas et al. (1998) agree that H diffusion involves some type of hydrogen interstitial state, but that the exact nature of the diffusion mechanism is not well identified. They propose a type of H diffusion motion more consistent with energetics calculations and experimental evidence. The authors use a tight-binding model to calculate the energy to break a Si—H bond and place the hydrogen atom at nearby or distant silicon sites. They find that hydrogen is very reactive, and can form a new Si—H bond by breaking a Si—Si bond. This is represented by: * * x y x y Si—H Si —Si Si Si —H Si + → + + (5) These configurations typically consist of two DBs (Si* and Siy *). Remarkably, the energy of reaction is low even when the Si—H bond being broken is not a weak one. The authors
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  Optoelectronics - Materialsand Techniques 12 identify a manifold of transport states, which depends on the bond length deviation of Si— Si bonds, through the a-Si:H network. They propose that diffusion of hydrogen proceeds by the hydrogen atom breaking and reforming Si—Si bonds in the network, with the diffusing hydrogen carrying a transporting DB along. The basic conclusion is that hydrogen motion is very reactive and does not rely on the existing DBs in the network. It is also somewhat different from the hydrogen motion through bond-centered sites, which according to these authors is a less reactive process since it implies that the Si—Si bond must stretch outwards to accommodate the hydrogen atom. The calculation of the energy barriers is complex, but the authors set an upper limit of 0.8 eV and a likely value of 0.5 eV. One of the most important and popular models for hydrogen diffusion is the one proposed by Branz (1999) with the Hydrogen Collision Model (HCM). In this model, DBs are created when recombination of light induced carriers stimulates emission of mobile hydrogen from Si-H bonds according to: Si H DB Si H /DB − → + − (6) The basic process is described by the following steps: 1. The mobile hydrogen atom goes to a Si—Si bond 2. The bond is broken, forming a temporary Si-H and a DB 3. The hydrogen atom hops to another Si—Si bond, again breaking the bond, while the previous bond reconstructs itself. 4. The mobile hydrogen atom continues to hop (it can be proven that its binding energy to the various bonds it breaks on its way is weaker than regular Si-H). 5. Eventually, the mobile hydrogen atom re-traps to Si-H through one of two mechanisms, described below: The first is a normal re-trapping to an immobile DB, given by Si-H/DB+DB→ Si-H, that is, the inverse process of eq. [6]. Basically, one can see this phenomenon as an H jumping to an ordinary DB, or as the formation of a bond between the mobile DB - that accompanies the mobile H – and the immobile DB. In both cases, no net DB results from the process. The second mechanism can be described as a re-trapping process of the form Si-H/DB+Si- H/DB→ M(Si-H)2. This is far less frequent than normal re-trapping. It represents the collision of 2 mobile H atoms (and their accompanying mobile DBs) that associate into a metastable complex, M(Si-H)2, containing a pair of Si-H bonds in close proximity. The meaning of this process is that there is a net formation of 2 DBs (the ones left behind by the original Si-H bonds), thus resulting in the SW effect. All this can be summarized by the following reaction: ( )2 2Si H 2DB 2Si H /DB 2DB M Si H − → + − → + − (7) Experimental studies by Cheong et al. (2000) examine one of the main and most controversial assumptions of the HC model, namely, that the photo-generated mobile hydrogen atoms can move a long distance at room temperature. They devised an experiment to observe hydrogen motion at room temperature, since detection with traditional methods such as IR and deuterium tracing is inadequate. By using the high sensitivity of the Raman spectrum of electro-chromic amorphous tungsten (a-WO3) to hydrogen insertion the authors were able to detect the long-range motion of hydrogen at room temperature.
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  Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling 13 The floating bond (FB) model as described, for instance, in (Biswas & Pan 2003), proposes an alternative explanation for the H diffusion process. To put it simply, compared with the hydrogen-collision model proposed by Branz, the creation of DBs is mediated by floating bonds rather than hydrogen atoms. The proponents of the FB model point out that the emission rate of mobile hydrogen should be larger than the creation rate of a pair of DB and FB and that the mobility of movable hydrogen should be faster than that of FB, leading to the dominance of the Branz mechanism for DB creation. However, one should note that the possibility of the DB creation by the mobile hydrogen in the case of the HC model is very small, but DB-FB pair creation directly leads to the creation of DB. As mentioned previously, the hydrogen distribution is also an important parameter in determining the dynamics of the SW effect. For instance, Tuttle & Adams (1997) show that the energetic and properties of H-atoms must be analyzed considering also their phases, i.e., dilute or clustered. According to the authors, the relative ratio of these phases and their distribution has an important role in determining the properties of a-Si:H. This is a fundamental fact that needs to be taken into account, if a model has to be used to simulate processes connected with hydrogen dynamics, including the testing of the models outlined above that have come to prominence as explanations of the Staebler-Wronski effect. Gaspari et al. (2010) have examined the hydrogen distribution in simulated samples, obtained by ab-initio Molecular Dynamics (AIMD), by examining the H-H radial distribution function. It was noted that the H-structure and its distribution within the underlying silicon network is crucial in determining the properties of a-Si:H and for finding whether the sample possesses high quality characteristics for photovoltaic or micro-electronic applications. These findings are in agreement with results reported in (Tuttle & Adams, 1997), and indicate that the dilute vs. clustered distribution ratio, combined with a proper interatomic distance, plays a major role in determining the properties of a-Si:H. More recent studies continue to refine old models and propose new ones; however the focus has now shifted on the computational aspect of the modeling and, in particular, the realism of the model structure. This topic will be discussed in section 5. In the following section, a unique procedure providing a novel experimental and theoretical analysis regarding dangling bond formation and annealing in a-Si:H is presented. The approach employed is the incorporation of tritium into a-Si:H. 4. Tritiated amorphous silicon In order to shed light on the role of hydrogen in defect dynamics, it would be desirable to be able to control the evolution of DBs and correlate it with hydrogen dynamics. However, a quantitative study in which the density of DBs is changed usually involves changing deposition conditions, or high-temperature annealing, or damaging the material with high- energy particles or light (Danesh et al., 2005; Schneider & Schröder, 1990; Sholz et al., 1994). These procedures modify, to varying degree, other structural properties of the material, making it difficult to isolate the effect of DBs on the optoelectronic properties of a-Si : H. An alternative approach to investigating the properties of a-Si:H has been to substitute hydrogen with one if its isotopes, tritium, and use the effects of the radioactive decay process of tritium as a means to follow the dynamics of defect creation and annealing, and their impact on the opto-electronic properties (Costea et al., 2000; Gaspari et al., 2000; Kherani et al., 2008; Kosteski et al., 2000, 2003, 2005; Zukotynski et al., 2002).
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  Optoelectronics - Materialsand Techniques 14 4.1 Dangling bond formation due to tritium decay Tritium decay is described by the formula: 3 + - T He + + ν β → , where − β is the beta-particle emitted by the Tritium atom and ν is an anti-neutrino. The half-life of tritium is 12.3 years and the mean energy of the beta particle is 5.69 keV. The bonding of tritium in amorphous semiconductors is stable at room temperature (Sidhu et al., 1999). Because of the direct-like nature of the energy gap in amorphous semiconductors, radiative recombination of electron-hole pairs is allowed. Furthermore, the nuclear decay process which leads to the emission of a β particle from a Si-T site will produce a 3He atom which will diffuse out and leave behind a dangling bond (or recombination center, as shown in Figure 6). This process has led to a novel way of investigating the properties of amorphous silicon and amorphous semiconductors in general; see, for instance, (Kosteski et al., 2000, 2003; Zukotynski et al., 2002). It should be noted that tritium decay has side effects that might in turn have an impact on DB formation (i.e., He recoil, e-h pair recombination). These issues have been addressed and it was proven that the impact of the secondary effects is negligible. A detailed analysis of the relative impact of electron-hole pair recombination is presented in the following sub-section. Fig. 6. The mechanism of tritium decay and dangling bond formation: I – Tritium is bonded to a Si atom and shares 2 covalently bonded electrons; II – Tritium decays into 3He+ and emits a high energy electron (β emission); III – The 3He+ ion captures both electrons shared with the Si atom and leaves behind a positive dangling bond (DB+); IV – The positive DB can become a neutral DB by capturing in turn an electron generated by collisions between the original β particle and the silicon matrix. An average of 1500 electron-hole pairs per β emission is generated via this process (Kosteski et al., 2000). 4.2 Effect of electron-hole pairs on dangling bonds The tritium decay process results in the emission of a β particle with a mean energy of 5.7 keV. The β decay transforms the tritium atom into 3He+ with a, maximum, recoil energy of about 3 eV. Furthermore, electron-hole (e-h) pairs will be created in the material as the β particle transfers its energy to the silicon matrix. It has been shown that electrons with keV energies are unlikely to produce atomic displacements (Dubeau et al., 1996). They lose energy in collisions with electrons in the material, through which they pass, creating e-h pairs that thermalize by shedding ‘Raman’ phonons. In the case of a-Si:H, a pair is created for every 4.3 eV of energy lost. It can also be shown that Helium recoil is expected to create little damage (Street et al., 1979; Stutzmann, 1991).
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  Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling 15 On the other hand, e-h pairs recombination has also been associated with the creation of DBs (Yelon et al., 2000) and with the mechanism of hydrogen diffusion in a-Si:H (Branz et al., 1993, 1999; Cheong et al., 2000; Santos et al., 1991, 1992). The relative impact of e-h pairs on DB formation can be estimated as follows: An estimate of the number of e-h pairs generated by tritium decay in an a-Si:H:T film is given by the formula (Kosteski et al., 2000): Tv avg G n E / = λ ε (8) Where λ is the decay constant of tritium (λ = 1.78 x 10-9 s-1), Eavg is the average energy of a beta particle (5.7 keV) and ε is the energy needed to produce an electron-hole pair (4.3 eV); nTv is the volume density of atomic tritium. Assuming nTv~5 at.%, as confirmed by IR measurements for standard tritiated samples used in the investigation (Gaspari et al., 2000; Sidhu at al., 1999), we obtain: 15 3 1 G 5.9 x 10 cm s − − = (9) The effects of electron beam irradiation on a-Si:H has been extensively investigated (Schneider & Schroder, 1990; Scholz et al., 1994). In particular, the authors examined the effect of keV electrons – in the range 1 to 30 keV – and determined that metastable states were formed as a result of the irradiation. The results of their work can be summarized as follows: • The annealing kinetics of these defects is the same as for those produced by light. • The DB creation kinetics is linear with time over more than two decades in time. • The approach to saturation goes from linear to saturated with a change of less than a factor of 2 in the total number of metastable defects. • There is no sub-linear regime, in contrast to light-induced defect creation and annealing. • The saturation defect density is independent of dose rate over three orders of magnitude and independent of primary electron energy from less than 1 to more than 30 keV . • Finally, the defect creation rate and saturation are independent of temperature from 170 oC to room temperature. Yelon et al. (2000) applied the hydrogen collision model developed by Branz to explain the origin of metastable states due to ion bombardment. Since the work by Schroder and co- workers deals in the most part with room temperature 20 keV irradiation, at 170 mW cm-2, of samples, d=0.6 μm thick, the authors have compared their model to the above-mentioned experimental conditions. The intensity of the e-beam corresponds to a flux of Φ = 5.3 x 1013 cm-2 s-1 electrons striking the sample. In this case, the initial measured creation rate of dangling bonds is R0 = 2 x 1015 cm-3s-1. They obtained for the volume-averaged pair generation rate the number: 20 3 1 p R ΦN /d 6 x 10 cm s − − = = (10) which is comparable to that of light degradation at the smaller intensities for which smaller DB creation rates and sub-linear (Ndb ∝ t1/3) kinetics are observed.
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  Optoelectronics - Materialsand Techniques 16 By applying Yelon’s treatment to e-h pairs generated by beta particles originating from tritium decay, it can be shown that the number of dangling bond due to e-h recombination is 2 orders of magnitude less than the number due to the direct conversion of a Si—T bond into a DB. This is shown in Figure 7. T DB e-h DB Fig. 7. DB generation rates due to direct conversion of a Si—T bond into a DB and to e-h pair recombination. In summary, the studies of the properties of a-Si:H:T thin films have established the following: Creation of DBs, or induced changes to the host structure, by other causes associated with Tritium decay (i.e., irradiation damage from β particles, e-h pair recombination, He recoil) is negligible. This is of particular relevance, since it makes it possible to eliminate any side effects of Tritium decay from the analysis, in particular electron irradiation. Calculations have shown that DBs due to electron irradiation caused by Tritium decay account for less than 1% of the total number of dangling bonds created by the release of the He atom from the Si bond. Annealing recovers the photoluminescence (PL) and the electrical properties of the tritiated samples (see figures 8a and 8b). This is the fundamental aspect of the SW process in a-Si:H:T since the dangling bonds are created following a well known nuclear process, in which the Tritium atom is replaced by an Helium atom which diffuses away from the original bond. This clearly shows that the annealing process cannot be a simple reversal of the formation process of DBs. Therefore the passivation of DBs appears to be independent of the DB formation process and linked to the a-Si:H structure and bonding configuration. The significance of point 2 should be emphasized: many models of the SW effect, in particular earlier ones (Adler, 1984; Stutzmann et al., 1985), are based on the assumption that the annealing mechanism is basically the reverse process of the dangling bond formation mechanism. Although some fundamental common aspect might be present in both phases of the SW process, it is possible to focus on each phase independently. In
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  Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling 17 particular, the objective is to establish the relationship between the structure of a-Si:H (i.e. void fraction, hydrogen bonding, etc.) and the curing process (i.e., is it due to H diffusion and/or a network readjustment during annealing? If H diffusion is involved, how is it mediated?). A number of fundamental issues remain unresolved: microscopic atom dynamics, for instance, influences atomic structure, chemical bonding, diffusion and vibrations, and are difficult to study both experimentally and theoretically. Furthermore, it is not clear why the number of DBs measured by electron spin resonance (ESR) is one order of magnitude less than expected, as noted by Whitaker et al. (2004) and Ju et al. (2007). (a) (b) Fig. 8. (a) Decay of PL signal for a tritiated amorphous silicon film. The black curve was taken one hour after deposition, and the decay vs. time is shown by the coloured curves, labelled in the figure by the post-deposition times at which the spectra were recorded. (b) PL spectrum taken for the same sample after annealing. The open dots show the PL signal (almost completely quenched) approximately two weeks after deposition. The filled circles show the recovered PL signal after annealing at 150 oC for one hour. 4.3 The stretched exponential behaviour One of the most interesting observations for tritiated a-Si:H is that the increase in DBs density can be described by a stretched-exponential function, as shown in figure 9 (Kosteski et al., 2003). This indicates that the expected exponential growth of DBs due to tritium decay has a competing process of DB passivation that leads eventually to saturation and a steady- state distribution of DBs. One hypothesis is that hydrogen atoms that diffuse from nearby sites annihilate the tritium induced defects. In this case, hydrogen detachment from such a site should not be accompanied by the creation of a new defect. This might be possible if the hydrogen atom comes from a large reservoir of paired hydrogen sites such as double-hydrogen complexes (Branz, 1999) or hydrogenated vacancies (Zhang Branz, 2001). It should be noted that a similar concept has been proposed for DBs in tail states (Longeau, 2000). The stretched exponential behaviour has been reported for a number of different conditions for a-Si:H, for instance: Danesh et al. (2005) have investigated a-Si:H degradation for 18 MeV electron beam bombardment. They observed that “the annealing kinetics obeys the stretched-exponential law which is known to be typical for a-Si:H films underwent the light- induced degradation.”
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  Optoelectronics - Materialsand Techniques 18 The stretched exponential behaviour has also been investigated, among others, by Morigaki Hikita (2007) and Morigaki et al. (2008) for light-induced defects. Morigaki considers three possible annealing mechanisms, i.e., hydrogen dissociation from two closely located Si-H bonds, dissociation from a hydrogen molecule, and reconstruction of silicon network (Zukotynski et al., 2002). Although the dissociation values for the first two processes are higher than the activation energy of thermal annealing (~1 eV), Morigaki argues that the values are still reasonable for activation energies higher than 1 eV. It is not clear, however, why the same stretched exponential behaviour and similar saturation values are observed for tritiated samples, unexposed to light, or for high energy electron bombardment. The stretched exponential behaviour observed for tritiated amorphous silicon is consistent with the reported values for light induced defects. The major difference between the defect re-equilibration dynamics in a-Si:H:T vs. illuminated a-Si:H is that in the former case there is no assistance from external excitation due to photons. Yet, it appears that the a-Si:H:T structure still adjusts to the increase in the density of dangling bonds. As previously mentioned, this indicates that a common process, related to a combination of network reconstruction and H-diffusion, should be at the basis of the annealing phase, regardless of the DB formation mechanism. The hydrogen bonding structure and the hydrogen diffusion process are linked by the activation energies required to bring a hydrogen atom into the mobility band, as indicated, for instance, in (Powell Deane, 1996). As the above mentioned authors point out in their “defect pool” model, hydrogen transitions involve silicon DB defects, the defect energy and the Fermi energy. The hydrogen density of states (HDOS) is a crucial factor in this model, and is at the basis of the defect equilibration dynamics. N0 = 5.6x10 16 cm -3 Nsat= 6.9x10 17 cm -3 t = 1x10 6 sec α = 0.76 Fig. 9. Dangling bond rate of creation for a-Si:H:T measured by ESR (Kosteski, 2003). The dashed line is the theoretical prediction for one dangling bond per tritium decay. The solid line is a stretched exponential fit to the experimental data. The parameters in the inset refer to the stretched exponential formula: ( ) ( )α d 0 sat 0 N = N + N - N × 1- exp -t τ ⎡ ⎤ ⎢ ⎥ ⎣ ⎦ .
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  Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling 19 5. Ab Initio molecular dynamics The main reason for the development of so many models and interpretations of hydrogen mobility is due to the difficulty of obtaining a clear microscopic picture of the dynamic processes via experimental techniques. The only way to investigate the trajectories of hydrogen atoms in a-Si:H is to extensively model it numerically. Many theoretical techniques have been employed to generate a realistic model of a-Si:H, including classical MD (Ishimaru, 2002; Izumi et al., 2005), tight-binding MD (Singh et al., 2004; Tuttle Adams, 1997), and Ab-Initio MD (Su Pantelides, 2002). Molecular dynamics has become one of the more powerful and frequently used tools for the correlation of the microscopic characteristic of materials with their macroscopic properties, observed experimentally. Some authors (Abtew Drabold, 2006) have applied the model to simulate the dynamics of hydrogen diffusion, in particular as a way of analyzing possible mechanisms behind the Staebler-Wronski (SW) effect. However, classical MD is not sufficiently accurate to describe the covalent bonding and forces in semiconductors, and a tight-binding approach; even DFT-LDA based MD (Singh et al., 2004) is not sufficiently transferable to non-crystalline systems. For amorphous Si, with no translational symmetry, presence of non-saturated Si bonds, hydrogen switching between host Si atoms and diffusion, the most accurate approach is first principles (or ab-initio) MD. Although AIMD has to be used with a smaller unit cell and the simulation time is shorter compared to, for instance, tight-binding, it is extremely accurate in the description of the interatomic potential, forces, and inharmonicity of the interaction. We have seen that the microscopic details of disordering, hydrogen migration and bonding within the amorphous silicon network are crucial for the understanding of a-Si:H, and for the improvement of the overall quality of the material. However, hydrogen migration and other complex properties must be investigated within a realistic structure to obtain useful information. Kupchak et al. (2008) and Gaspari et al. (2009) have shown that the commonly used radial distribution function (RDF) is not a sufficient validation parameter, and that the recreation of physically consistent vibrational spectra does, on the other hand, represent a sound validation protocol. The authors have used AIMD to model bulk a-Si:H under various conditions. At various key points in simulations, detailed pictures of measurable quantities related to the electronic structure have been calculated, assuring that the model remain close to nature. In AIMD one treats atoms classically, but the potential is calculated quantum mechanically using DFT. The solutions to the Schrödinger equation are approximated numerically. In terms of computing operations, these calculations involve mainly matrix operations and fast Fourier transforms (FFT). Details of this approach can be found in (Gaspari et al., 2009; Kupchak et al., 2008). This approach has led to reproducing all the fundamental vibrational signals in a-Si:H. Computer visualization was used to demonstrate some of the most important features, like the existence of poly-hydride bonding. The vibrational spectra obtained for the “good” samples were consistent with experimental data and showed also time stability, as shown in Figures 10 and 11. As previously mentioned, Gaspari et al. (2009) have examined the appropriateness of using RDF as a validation parameter for assessing the realism of a simulated structure. The authors compared a variety of a-Si:H samples with H content from 0 to 20%, and have found
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  Optoelectronics - Materialsand Techniques 20 that the RDF practically does not depend on the amount of hydrogen in the sample. Furthermore, all the calculated RDF agree reasonably well with the most recent and accurate RDF measurement for a-Si with no hydrogen. This reflects the fact that the most probable distance between neighboring atoms is equal to a sum of the atoms’ covalent radii. Even when hydrogen passivates the dangling bonds, this does not modify the Si–Si bond length. On the other hand, atomic vibrations do depend on microscopic bonding (bonds), their angular distribution, distortion or breaking. In fact, the experimental measurements demonstrate a variety of spectral features that obviously require microscopic theoretical interpretation. Furthermore, in order to further verify the validity of the model, the authors have also studied the special case of metastable Si-H-Si bonds, observed experimentally by Darwich et al. (1995), and have confirmed Darwich’s claim within experimental error. Gaspari et al. (2009) indicate that the decrease in the vibrational frequency with respect to that of a stable mono-hydride bond is due to the sharing of the hydrogen electron density between two Si atoms. This decreases the Si–H bond strength, increases the bond length and results in reduction of the vibrational frequency. Therefore, the band in the 1500-1800 cm-1 region can be interpreted as the signature of hydrogen metastable bonds, including the TCB bond, with variations in the frequency due to the different overlap between the H and the Si electron wave functions. Fig. 10. Hydrogen stretch vibrations for a-Si64-H10 system at high frequency (Kupchak et al., 2008). The solid black line shows all H-associated stretching vibrations, including dihydride modes (blue, short dash) and monohydride modes (red, long dash). Note the very close agreement with data by Lucovsky et al. (1989).
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  Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling 21 Fig. 11. Time dependent frequencies for a “good” sample. Note the absence of vibrations between the two main modes (2000 cm-1 and 640 cm-1) indicating stability of the bonds. The colour scale is related to the peak intensity, that is, white represents the strongest signal (peak), while black represents no vibrational signal (Kupchak et al., 2008). The investigation led by this author has proven that in order to validate the simulation of complex structure, bonding, and diffusion, a protocol needs to be established for the verification of the “realism” of the simulated models. Using hydrogenated amorphous silicon as an example, Gaspari et al. (2009, 2010) have unambiguously demonstrated that reproduction of the radial distribution function, used commonly in numerical simulations, is not sufficient and must be complemented with verification of other, more complex, macroscopic properties. By focusing on the vibrational modes of the amorphous system, it was proven that the vibrational spectra represent a crucial testing tool for non-crystalline materials because of their complexity and sensitive link to structure and bonding configuration. Successful reproduction of all the experimentally observed vibrational features for a-Si:H has proven the validity of the algorithm and indicates that hydrogen structure and dynamics are extremely sensitive to the parameters of the model. In order to correctly apply a numerical model to extract such important macroscopic parameters as density of states, optical gaps, and migration dynamics, the accuracy should be verified first by the derivation of the standard vibrational modes and comparison with experimental observation. Indeed, the importance of hydrogen distribution and its connection to hydrogen mobility is demonstrated by recent investigations, both experimental and theoretical, on the role of hydrogen in a-Si:H. For instance, Fehr et al. (2010) investigated the distribution of hydrogen
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  Optoelectronics - Materialsand Techniques 22 atoms around native dangling bonds in a-Si:H by electron-nuclear double resonance (ENDOR). The authors suggest that the hydrogen distribution is continuous and homogeneous and there is no indication for a short-range order between hydrogen atoms and dangling bonds. This is in contrast with current understanding that hydrogen is distributed as a succession of clustered and diluted phases (Gaspari et al., 2010; Tuttle Adams, 1997). Such controversies can only be addressed by using a rigorous, realistic model to simulate properties and dynamic processes. 6. Conclusions Hydrogenated Amorphous Silicon (a-Si:H) has been the subject of intensive investigation for over 30 years. The main role of hydrogen in amorphous silicon is the passivation of the Si dangling bonds (DBs) to restore a proper energy gap and the semiconducting properties, thus enabling extensive application of a-Si:H in the microelectronics and the photovoltaic industry. Due to the importance of hydrogen, many experimental methods have been used to characterize the DBs passivation, bonding chemistry and related mechanisms of degradation of the material. Among the numerous experimental techniques used to study a- Si:H and the role of hydrogen, the Fourier Transform Infrared Spectroscopy (FTIR) is used extensively to analyze vibrational spectra of a-Si:H. Although FTIR represents one of the most common and powerful techniques, no microscopic links between the observed vibrational features of the hydrogen and the microscopic properties of a-Si:H can be yet established by any experimental means. A number of other important fundamental issues remain unresolved for a-Si:H as well. Microscopic atom dynamics, for instance, influences atomic structure, chemical bonding, diffusion and vibrations, and are difficult to study both experimentally and theoretically. However, the microscopic details of disordering, hydrogen migration and bonding within the amorphous silicon network is crucial for the understanding of a-Si:H, and for the improvement of the overall quality of the material. The Staebler-Wronski effect epitomizes this need. It is generally accepted that a-Si:H light- soaking degradation, observed by Staebler and Wronski, is caused by Si-H bonds breaking during illumination. However, the microscopic details of the SW effect are still controversial and it is not clear how to experimentally predict the stability of a-Si:H films, grown at particular temperature and hydrogen concentration, with respect to light induced degradation. Furthermore, a number of alternative techniques have been used to create dangling bonds, and the same dynamics has been observed in the curing (annealing) phase. That is, no matter how the dangling bonds were formed, a similar curing process occurs during annealing. This might be due to diffusion of hydrogen atoms, structural readjustment, or a combination of the two. In this chapter I have briefly summarized how the optical and electronic properties of a-Si:H are dependent on the hydrogen content and pointed out that the challenge of uncovering the microscopic details of hydrogen bonding and distribution and their correlation with hydrogen dynamics cannot be answered by standard experimental techniques. On the other hand, with the continuous improvement of computational capacity and software quality, the simulation of realistic structures is becoming ever more feasible. In particular, Ab Initio Molecular Dynamics (AIMD) allows highly accurate simulation of the dynamical properties of various systems, including amorphous materials.
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  Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling 23 The goal of such simulations is to be able to reproduce dynamic processes and follow the diffusion of hydrogen, the bond breaking processes, and the structural reorganization of the material, following external perturbations. The DB creation process in tritiated amorphous silicon can provide a simple and convenient source of experimental data that can be used as a basis for such simulations, since the tritium decay process is well understood, and its effect on a-Si:H can be treated as the simple removal of an hydrogen atom from an existing Si—H bond. The main challenge is of course to make sure that the simulated structure is indeed a realistic one. The author of this chapter has shown that several models lack the necessary realism, since the validation of the model is based on the radial distribution function of the Si—Si bonds. The author has also shown that the reproduction of the vibrational modes of a- Si:H represents a much better validation test for a realistic structure. As the continuous advances in computational science will allow for the use of bigger simulated structures, the future direction of these studies should aim at reproducing other fundamental properties, such as the band-gap, the density of states, etc. By achieving this goal, it will be possible then to simulate dynamic processes too, such as the SW effect, and to shed light both on the formation phase of the dangling bonds and on the curing phase. 7. Acknowledgment The work by the author was supported by the Shared Hierarchical Academic Research Computing Network (SHARCNET) and Natural Sciences and Engineering Research Council of Canada (NSERC). The author would also like to thank Dr. A. Chkrebtii for his invaluable contribution and leadership in the development of the AIMD algorithm. Thanks go also to Dr. J.M. Perz, Dr. S. Zukotynski, and Dr. N. P. Kherani for their support and helpful discussions spread over 20 years. 8. References Abtew, T.A., Drabold, D.A. (2006) Phys. Rev. B, Vol. 74, 085201. Adler D., (1984) AIP Conference Proceedings, n 120, 70-77. Akkaya,A., Akta,G. (1995). Mater. Lett., Vol. 22, 271. Baranovski, S. (2006). Charge transport in disordered solids with applications in electronics. John Wiley $ Sons, ISBN: 9780470095041, New York. Biswas, R., Li, Q., Pan, B.C., Yoon, Y. (1998) Phys. Rev. B, Vol. 57, 2253. Biswas, R., Pan, B.C. (2003) Solar Energy Materials and Solar Cells, Vol. 78, 447. Branz, H.M., Asher, S.E., Nelson, B.P., (1993)Phys. Rev. B Vol. 47, 7061. Branz, H.M. (1999) Phys. Rev. B Vol. 59, 5498. Branz, H.M., Asher, S.E., Gleskova, H., Wagner S., (1999)Phys. Rev. B Vol. 59, 5513. Bruno, G., Capezzuto P., Madan A., (Eds.) (1995) Plasma Deposition of Amorphous Silicon- Based Materials. Academic Press, ISBN: 9780121379407, Burlington, MA. Cheong, H.M., Lee, S.H., Nelson, B.P., Mascarenhas, A., Deb, S.K., (2000) App. Phys. Lett. B Vol. 77, 2686. Cody, G.D., Tiedje, T., Abeles, B., Brooks B., Goldstein, Y. (1981) Phys. Rev. Lett. Vol. 47, 1480.
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  Optoelectronics - Materialsand Techniques 24 Costea, S., Gaspari, F., Kosteski, T., Zukotynski, S., Kherani, N. P., Shmayda, W.T. (2000) Mat. Res. Soc. Symp. Proc, Vol. 609, A27.4 (2000). Costea, S., Pisana, S., Kherani, N.P., Gaspari., F., Kosteski, T., Shmayda, W.T., Zukotynski, S. (2005) Fusion science and technology Vol. 48, 712. Danesh, P., Pantchev, B., Vlaikova, A. (2005) Nuclear Instruments and Methods in Physics Research B, Vol. 239, 370. Daouahi, M., Ben Othmane, A., Zellama, K., Zeinert, A., Essamet, M.,Bouchriha, H. (2001) Solid State Communications Vol. 120, 243. Darwich, R., Roca I. Cabarrocas, P., Vallon, S., Ossikovski, R., Morin, P., Zellama, K. (1995) Phil. Mag. B, Vol. 72, 363. Dubeau, J., Hamel, L.A., Pochet, T., (1996) Phys. Rev. B 53, 10 740 Fehr, M., Schnegg, A., Teutloff, C., Bittl, R., Astakhov, O., Finger, F., Rech, B., Lips, K. (2010) Physica Status Solidi A, Vol. 207, 552. Gaspari, F., O’Leary, S.K., Zukotynski, S., Perz, J. (1993) J. Non-Cryst. Solids Vol. 155, 149. Gaspari, F., Kosteski, T., Zukotynski, S., Kherani, N. P., Shmayda, W. (2000) Phil. Mag. B, Vol. 80, 561. Gaspari, F., Shkrebtii, A., Kupchak, I., Perz, J.M. (2009) Phys. Rev. B Vol 79, 224203. Gaspari, F., Shkrebtii, A., Kupchak, I., Teatro, T., Ibrahim, Z.A. (2010) 35th IEEE Photovoltaic Specialists Conference Proceedings, Honolulu Hawaii, June 20-25, 003671-75. Ishimaru, M. (2002) J. Appl. Phys. Vol. 91, 686. Izumi, S., Hara, S., Kumagai, T., Sakai, S. (2005) J. Cryst. Growth Vol. 274, 47. Jackson, W.B., Tsai, C.C. (1992) Phys. Rev. B, Vol. 45, 6564. Jackson, W.B., Santos, P.V., Tsai, C.C. (1993) Phys. Rev. B, Vol. 47, 9993. Jeffrey, F.R., Shanks, H.R., Danielson, G.C. (1979) Appl. Phys. Lett. Vol. 50, 7034. Kasap, S. (2005) Principles of Electronic Materials and Devices. McGraw-Hill, retrieved from http://Materials.Usask.Ca. Kato, S., Aoki, T. (1985) J. Non-Cryst. Solids Vols. 7778, 813. Kherani, N.P., Liu, B., Virk, K., Kosteski, T., Gaspari, F., Shmayda, W.T., Zukotynski, S., Chen, K.P. (2008) J. Appl. Phys. Vol. 103, 024906. Knights, J.C., Lujan, R.A. (1979) Appl. Phys. Lett. Vol. 35, 244. Kosteski, T., Gaspari, F., Hum, D., Costea, S., Zukotynski, S., Kherani, N.P., Shmayda, W.T. (2000) Mat. Res. Soc. Symp. Proc. Vol. 609, A30.1. Kosteski, T., Stradins, P., Kherani, N.P., Gaspari, F., Shmayda, W.T., Sidhu, L., Zukotynski, S. (2003) IEE Proc. Circuits, Devices and Syst., special issue on Amorphous and Microcrystalline Semiconductor Devices, Vol. 150 no. 4, 274. Kupchak, I. M., Gaspari, F., Shkrebtii, A. I., Perz, J. M. (2008) J. Appl. Phys. Vol. 104, 123525-1 Laaziri, K., Kycia, S., Roorda, S., Chicoine, M. Robertson, J. L., Wang, J., Moss, S. C. (1999) Phys. Rev. Lett. Vol. 82, 3460. Ley, L. (1983) “Photoemission and Optical properties”, in The Physics of Hydrogenated Amorphous Silicon, Vol II, Eds. J.D. Joannopoulos G. Lucovski, Springer-Vderlag, ISBN: 0387128077, New York. Longeaud, C., Roy, D., Teukam Hangouan, Z. (2000) App. Phys. Lett. Vol. 77, 3604. Lucovski, G., Davidson, B.N., Parsons, G.N., Wang, C. (1989) J. Non-Cryst. Solids Vol. 114, 154. Malik, S. M., O'Leary, S. K. (2004) J. Non Cryst. Solids, Vol. 336, 64.
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  Optoelectronic Properties ofAmorphous Silicon the Role of Hydrogen: From Experiment to Modeling 25 Morigaki, K., Hikita, H. (2007) Phys. Rev. B 76, 085201 Morigaki, K., Takeda, K., Hikita, H., Ogihara, C., Roca i Cabarrocas, P. (2008) J. Non-Cryst. Solids, Vol. 354, 2131. Mott, N. (1983) “Conductivity, Localization, and the Mobility Edge”, in The Physics of Hydrogenated Amorphous Silicon, Vol II, Eds. J.D. Joannopoulos G. Lucovski, Springer-Verlag, ISBN: 0387128077, New York. O'Leary, S.K., Sidhu, L.S., Zukotynski, S., Perz, J.M. (1996) Canadian Journal of Physics, Vol. 74, S256-9. Powell, M.J., Deane, S.C., (1996) Phys. Rev. B, Vol. 53, 10121. Rui, Y., Mei, J., Xu, J., Yang, L., Li, W., Chen, K. (2005) Proceedings of SPIE - The International Society for Optical Engineering, Vol. 5774, 279. Santos, P.V., Johnson, M.N., Street, R.A. (1991) Phys. Rev. Lett. Vol. 67, 2686. Santos, P.V., Johnson, N.M., Street, R.A., (1992) Mat. Res. Symp. Proc. Vol. 258, 353. Santos, P.V., Johnson, M.N., Street, R.A. (1993) J. Non-Cryst. Solids Vols. 164-166, Part I, 277. Schneider, U., Schröder, B. (1990) Photovoltaic Specialists Conference. Conference Record of the Twenty First IEEE, vol. 2, 1521. Searle, T. (Ed.) (1998) Amorphous Silicon and its Alloys, INSPEC, ISBN: 0852969228, London. Sholz, A., Schröder, B., Oechsner, H. (1994) Mat. Res. Symp. Proc. Vol. 336, 293. Sidhu, L. S., Kosteski, T., Zukotynski, S., Kherani, N. P. (1999) J. Appl. Phys. Vol. 85, 2574. Singh, R., Prakash, S., Shukla, N., Prasad, R. (2004) Phys. Rev. B Vol. 70, 115213. Smets, A.H.M., van de Sanden, M.C.M. (2007) Phys. Rev. B, Vol. 76, 073202. Staebler, D.L., Wronski, C.R. (1977) Appl. Phys. Lett. Vol. 31, 292. Street, R.A., Biegelsen, D., Stuke, J., (1979) Philos. Mag. B Vol. 40, 451. Street, R.A. (1980) Phys. Rev. B, Vol. 21, 5775. Street, R.A. (1991) Hydrogenated Amorphous Silicon, Cambridge University Press, ISBN: 0521371562, New York. Street, R.A. (Ed.) (2000) Technology and Applications of Amorphous Silicon, Springer Verlag, ISBN: 3540657142, New York. Street, R.A., Tsai, C.C. (1988) Philos. Mag. Vol. B57, 663. Stutzmann M., Jackson W.B., Tsai, C.C. (1985), Phys. Rev. B, Vol. 32, n 1, 23-47 Stutzmann M., (1991) in Amorphous and Microcrystalline Amorphous Devices, Vol. II, Ed. J. Kanicki, Atech House, Boston, p. 129. Tauc, J., Grigorovici, R., Vancu, A. (1966) Phys. Status Solidi, Vol. 15, 627. Thevaril, J.J., O’Leary, S.K. (2010) J. Appl. Phys., Vol. 107, 083105. Tuttle, B., Adams, J. B. (1997) Phys. Rev. B Vol. 56, 4565. Ukpong, A.M. ((2007) Turkish Journal of Physics, Vol. 31, 317. Van de Walle, C.G., Street, R.A. (1994) Phys. Rev. B, Vol. 49, n 20, 14766-9. Van de Walle, C.G., Street, R.A. (1995) Mat. Res. Soc. Symp. Proc., Vol. 377, 389. Yelon, A., Fritzsche, H, Branz, H.M., (2000) J. Non-Cryst. Sol. Vols. 266-268, 437. Ju, T., Whitaker, J., Zukotynski, S., Kherani, N., Taylor, P.C., Stradins, P. (2007) Mat. Res. Soc. Symp. Proc. Vol. 989, 9. Whitaker J., Viner, J., Zukotynski, S., Johnson, E., Taylor, P.C., Stradins, P. (2004) Mat. Res. Soc. Symp. Proc. Vol. 808, 153. Zanzucchi, P.J., Wronski, C.R., Carlson, D.E. (1977) J. Appl. Phys. Vol. 48, 5227.
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  Optoelectronics - Materialsand Techniques 26 Zeman, M. (2006) “Advanced Amorphous Silicon Solar Cell Technologies”, in Thin Film Solar Cells: Fabrication, Characterization and Applications, Eds. J. Poortmans V. Arkhipov, John Wiley Sons, New York. Zhang, S.B., Branz, H.M., (2001) Phys. Rev. Lett. Vol. 87, 105503 Zukotynski, S., Gaspari, F., Kherani, N., Kosteski, T., Law, K., Shmayda, W.T., Tan, C.M. (2002) J. Non-Cryst. Solids Vols. 299-302, 476.
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  2 Silicon–Rich Silicon OxideThin Films Fabricated by Electro-Chemical Method Pham Van Hoi, Do Thuy Chi, Bui Huy and Nguyen Thuy Van Vietnam Academy of Science and Technology, Vietnam 1. Introduction Porous silicon (PS) has attracted increasing research interest in basic physics as well as applications since 1990 when Canham reported on the efficient visible photoluminescence (PL) of porous silicon (Canham, 1990). Structurally, PS consists of many pores and silicon residuals and usually can be described as a homogeneous mixture of silicon, air and, even silicon dioxide. Based on porosity, PS can be classified into three types: nano, meso- and macro-pores. In the case of PS nano-pores, the size of both the silicon residuals and the air voids (pores) can be in the range of few nanometers. The exciton Bohr radius in Si is around 4.3 nm, so that quantum confinement can occur and change the electronic structure of those silicon nanocrystals. On the other hand, because the value of porosity is directly linked to the effective index of refraction of the PS layer, this layer appears as an effective medium, where the refractive index has a tunable value between the index of refraction of bulk Si and that of the air (pores). Those changes in the electronic structure and refractive index of PS when compared with bulk Si make it fascinating as both a low-dimensional material and an optical one. The considerable and controllable changes in the electronic structure and refractive index of PS fabricated by electrochemical anodization make it a promising material for photonics in comparison with bulk silicon and/ or pure silica. Using the oxidation process in O2 environment at high temperature, the PS samples become silicon- rich silicon oxides (SRSO), which has high chemical instability and avoids the aging of the PS that is important condition for optical devices such as planar optical waveguides, optical interference filters, micro-cavities, etc (Bettotti et al., 2002). During the last decade, Erbium (Er)-doped silicon-rich silicon oxide has attracted much interest due to its big potential application in Si-based optoelectronic devices for telecom and optical sensors. The Er-ions implanted in SRSO materials produce light emission at around wavelength range of 1540 nm, which corresponds to minimum light absorption in silica-based glass fibers. In this regard, a lot of studies have been carried out to improve the luminescence efficiency of this material. Such studies have revealed that co-implantation of Er and O2 induce a strong enhancement in the Er-ions related emission at range of 1540 nm. In first case, samples were prepared by co-implanting Si and Er into silica thin films or co-sputtering Si, Er2O3 and SiO2 on the silicon substrate (Shin et al., 1995). In second case, samples were prepared by implanting Er-ions into SiO2 films containing Si-nanocrystals (nc-Si) and/or by Er-ion electrochemical deposition on silicon-rich oxide (SRSO) layers. The room temperature luminescence emission at the range of 1540 nm from Er-electrochemically doped porous
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  Optoelectronics - Materialsand Techniques 28 silicon was first reported by Kimura T. et al in 1994 (Kimura et al., 1994) and then followed by some other authors. The strong luminescence emission around 1540nm-range of Er- doped SRSO layers at room temperature can be explained by energy transfer from excitons confined in the nc-Si to Er-ions and the evidence of energy transfer had been revealed in photo-luminescent excitation spectra in visible and infrared region when the exciting wavelength was not equalized to resonant absorption wavelength of Er-ions. Up to now, there are very few evidences of energy transfer given in the case of Er-electrochemically doped SRSO layers. In this book chapter, we will discuss the electrochemical method for preparing SRSO based on PS layers and Er-doped SRSO thin films for waveguide, optical filter and micro-cavity. In concentrating on the controllable changes in the refractive index of PS, we would like to use SRSO as a material for photonic devices such as optical interference filters, micro-cavities, etc. As an optical material, we present the fabrication method and properties of planar optical waveguides, active optical waveguides and optical interference filters operated in the range of infrared wavelengths. The advantage of optical waveguide amplifier based on Erbium-doped SRSO is the efficient energy transfer from electron-hole pairs generated in the Si nanocrystals to their neighbor erbium ions, which decay by emitting light at 1540nm (Bui Huy et al., 2008). The excitation cross-section of Er-ions in Er-doped SRSO is strongly increased in comparison of this one in the Er-doped silica glasses, so that the pump efficiency in Er-doped SRSO waveguides can be very high. The effect of energy transfer in elaborated Er-doped SRSO waveguides has also been explored. In order to design and predict the properties of the optical interference filters and micro-cavity based on SRSO multilayer, a simulation program based on the Transfer Matrix Method (TMM) was set up and the possible causes the difference in reflectivity spectra from this simulation and that from elaborated filters and/or cavity have been also given (Bui Huy et al., 2011). The structure and optical properties of SRSO layers are characterized by FE-SEM (Hitachi S- 4800), M-line spectroscopy (Metricon 2010/M) and luminescent measurement. The energy transfer effect between silicon nanocrystals and Er ions in the SRSO layers has been obtained from experiments. With the above-mentioned aim in mind, this chapter consists of the following sections: Section 2 presents the electrochemical method for preparing PS samples, Section 3 shows SRSO bi-layers based on PS annealed in oxygen environment at high temperature as a passive and active waveguides, Section 4 shows PS and/or SRSO multilayer with periodical refractive index change as an optical filter, Section 5 presents PS and/or SRSO multilayer with DFB configuration as micro-cavity, and Section 6 gives conclusions. 2. Electrochemical method for making SRSO thin films The porous silicon thin films were formed from silicon wafers by electrochemical etching in hydro-fluoric acid, without the necessity of any deposition process (Smith et al., 1992). During this anodization process a part of the silicon is dissolved and the remaining crystalline silicon forms a sponge-like structure with porosity between some tens percent up to more than 90%. The microstructure of the PS depends on the doping level of the silicon wafers: the use of low doped p-type substrates results in nanoporous silicon (with pore and crystallite size less than 2 nm) and the use of highly doped substrates in mesoporous silicon (size of 2-50 nm) (Herino et al., 1987). In the both cases the structures are much smaller than the wavelength of visible light and the materials appear as a homogenous, effective optical
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  Silicon–Rich Silicon OxideThin Films Fabricated by Electro-Chemical Method 29 medium. The effective refractive index of the porous silicon thin films is mainly determined by the porosity which can be varied by several anodization parameters. The most suitable way is changing the anodization current density, with high current densities resulting in high porosities and low refractive indices. A porous silicon thin film consisting of void spaces in silicon is created as a result of the electrochemical dissolution process in hydrofluoric acid, which can be expressed as in Equation (Valance, 1997): 2 6 2 Si 4HF 2F 2h SiF H 2H − + − − + + + → + + (1) The mass transport of positive charge carriers (h+) in the substrate and reactant fluorine ions (F-) in the electrolyte are key components in the dissolution process. As described in the model by Lehmann and Gösele (Lehman Gösele, 1991), dissolution begins when holes reach the silicon surface under anodic bias and enable a fluorine ion to replace a hydrogen atom bonded to silicon. Due to the polarizing influence of the bound fluorine ion, further reaction are initiated in which fluorine ions continue to bind to the silicon atom and hydrogen gas is formed. When all four silicon bonds are broken, the silicon atoms become soluble and leave behind an atomic size corrugation in the former atomically flat surface. Pore formation continues at the surface irregularity where the electric field is concentrated and holes are available. The interpore space is depleted of holes, inhibiting sidewall dissolution. In general, the preparation process of Er-doped silicon-rich silicon oxide layers can be divided into 3 steps: making a porous silicon (PS) layer by anodic etching of a Si-crystalline wafer in a HF solution; Er-ion deposition on the PS layer in Er content solution; and using thermal annealing at high temperature in oxygen and/or inert gases to obtain SRSO materials. The PS sample preparation is carried out in two approaches: keeping the current and/or the potential at a constant value during the electrochemical deposition (ECD) process. The difference between these two methods is that in the constant potential ECD, an n-type Si-crystalline wafer is usually used without annealing steps while in the constant current ECD, p-type Si-wafers are used and need thermal annealing. In our work we used both ECD methods for making PS layers on n- and p-type Si-crystalline wafers. 2.1 Experimental procedure In the electrochemical method for fabrication of porous silicon thin films, silicon wafer acts as the anode and is situated at the bottom of the Teflon cell. The silicon wafer was coated Au-thin film in back-side and contacted to HF-resistant metallic electrode in the form of the disk. This electrode disk enables a uniform contact on the whole area of silicon wafer. The electrolyte is a mixture of hydrofluoric acid and ethanol (C2H5OH) at different concentrations and poured into the Teflon cell. The platinum wire, which is also chemically resistant to HF, acts as the cathode. The shape of the cathode is critical to ensuring homogeneous samples, because it must promote a uniform electric field while allowing hydrogen bubbles formed during the anodization process to escape. The Teflon cylindrical tube with diameters of 10-15mm was placed between the upper and lower parts of the Teflon cell. Finally, a stainless steel ring and nuts are used to hold the cell together. We can use either current or voltage source for the anodization process. In our experiments, we used the electrochemical system Autolab PGS-30 as the electric current source, which can control the current with the nano-Amper range. Figure 1 presents the experimental setup for
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  Optoelectronics - Materialsand Techniques 30 making porous silicon thin films. The computer-controlled electric source used for the electrochemical process, so precise control over current density and etching time were achieved, and then it is resulting in a good control of the refractive index and thickness over the individual layers forming the multilayer. The program is a LabView virtual instrument realized to control the fabrication process of monolayer and multilayer of porous silicon with a friendly interface. The program controls the different parameters of the electrochemical process via GPIB. Those parameters include two current steps (to form layers with different refractive indices), duration time of each step (to determine the thickness of each layer), delay time (time between two consecutive electrochemical currents), and number of period (number of multilayer structure). Fig. 1. Electrochemical etching setup for fabricating PS layers 2.2 Silicon samples The initial Si-crystalline wafers, n-type with resistivities of 1-5 Ω.cm and p-type with resistivities of 0.01-1 Ω.cm, were used for constant potential and constant current ECD, respectively. For the case of n-type silicon substrates we need to illuminate the back side of silicon wafers. Resistivity of silicon wafer strongly affects on quality of porous silicon layers. High resistivity wafer often makes porous silicon layers with rough surface and easily peeled off from Si-substrate during fabrication or drying process, while the low resistivity sample have more flat surface of porous silicon layers. In order to form Ohmic contacts on the samples, we deposited pure gold (Au) and/or aluminum (Al) on the back faces of the n- and the p-type samples, respectively. The Si-crystalline wafers were anodic etched in a HF- ethanol solution with HF concentrations from 10% to 30% at a constant current density of 10-60 mA.cm-2 for time durations from some seconds to 15 minutes for controlling the refractive index of the PS layers. If the current density is modulated during the anodization, alternating layers of different porosities are formed as the silicon dissolution occurs primarily at the etched front PS/silicon substrates (Frohnhoff et al., 1995). Although the interface roughness between stacks is about 10-20nm, light scattering at these interfaces turned out to be very low. For this reason such layer stacks can act as optical waveguides and/or interference filters if the refractive indices are chosen properly (Krüger et al., 1998).
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  Silicon–Rich Silicon OxideThin Films Fabricated by Electro-Chemical Method 31 3. Active waveguide based on SRSO thin films Initially, Canham proposed that the up-shift of the luminescence spectrum into the visible was due to quantum confinement in the silicon crystalline wire structure and that the hydride passivation of the Si wire was the reason for the high efficiency of the observed photoluminescence (PL). For a short time after that, spectroscopic studies conducted particularly on the polarization of the PL (Kovalev et al., 1996) and on features observed under conditions of resonant excitation (Calcott et al., 1993) have provided strong positive confirmation of the quantum confinement model. However, there were a lot of spectroscopic phenomena that can not be explained by the simple quantum confinement model. As such, numerous models have been put forward as alternative explanations for the PL from PS such as hydrogenated amorphous silicon, surface hydrides, defects, molecules, surface states (Amato Rosenbauer, 1997). It is well known that in PS the surface to volume ratio is very large, so the surface effects are expected to have a significant influence on the material properties, especially optical ones (Kanemitsu et al., 1993). Because the Si atoms in Si nanocrystals are either at the surface or a few lattice sites away, the arrangement of interfacial atomic bonds, i.e. the passivation with Si-H or Si-O bonds, strongly affects the energy distribution of electronic states (Wolkin et al., 1999). In order to study PS as low-dimensional photonic materials, we elaborate on the effect of ageing on the spectral, intensity and lifetime of PL from the silicon nanocrystals in PS. Experimental results show that the effect of ageing on the spectral, intensity and PL lifetime of PS depends on the size of silicon nanocrystals. We focus our attention on strong emission properties and employ PS as a material for light emission sources, i.e. light emitting diodes and micro-cavity lasers operated in the visible region. In concentrating on the controllable changes in the refractive index of PS, we would like to use PS as a material for photonic devices such as planar optical waveguides, optical waveguide amplifier, optical interference filters, etc. As an optical material, we present the fabrication method for silicon rich silicon oxides (SRSO) thin films and properties of planar optical waveguides, active optical waveguides and optical interference filters operated in the range of infrared wavelengths. The advantage of optical waveguide amplifier based on Erbium doped SRSO is the efficient energy transfer from electron-hole pairs generated in the Si nanocrystals to their near erbium ions, which decay by emitting light at 1540nm. The excitation cross-section of Er-ions in Er-doped SRSO is increased more than two orders in comparison of this one in the Er-doped silica glasses (Friolo et al., 2001), so that the pump efficiency in Er- doped SRSO waveguides can be very high. 3.1 Porous silicon as a low-dimension photonic material In the first part of this section we explain the effect of surface states on the PL properties of PS based on the ageing process in air. In the last part, we present the reason for the intense and stable luminescence of blue region which has been of great interest in recent studies (Gorelkinskii et al., 2008). Previous studies on the interaction of oxygen in air on the as- prepared PS (Wolkin et al., 1999) show that: I) the as-prepared samples were well passivated by hydrogen and free of oxygen, ii) after exposure to air the samples were gradually passivated by oxygen, and the red-shift of PL spectral occurred as samples exposure to air and was nearly completed after ageing of 24 h. It was suggests that the ageing process can be divided into two periods: the first one in which the transition of the luminescence mechanism occurs after exposing the sample to air for a short time, and the second one in which the non- radiative center concentration is changed by oxygen passivation (Bui Huy et al, 2003).
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  Optoelectronics - Materialsand Techniques 32 Fig. 2. PL spectra of the as-prepared samples and after exposure to air for 1-month; samples, denoted as 1,2 and 3, were prepared by the anodic etching in 20%, 13% and 10% HF solution, respectively. (a) sample 1, (b) sample 2 and (c) sample 3 In order to investigate the effect of surface passivation on the size of Si nanocrystals, a series of PS samples denoted as 1, 2 and 3 were prepared by anodic etching in 20%, 13% and 10% HF solution respectively. As seen in figure 2, the PL peaks of the as-prepared samples 1, 2 and 3 have energy levels of 1.73, 1.84 and 2.00 eV respectively. This is related to a decrease of particle size in the considered samples. The figure also reveals that the ageing produces a pronounced increase in PL intensity in sample 1 and only a slightly increase in samples 2 and 3. As seen in figure 3, the decay rate of the as-prepared samples (the curves 1a, 2a and 3a) shows that the concentration of non-radiative centers in sample 1 is higher than those in samples 2 and 3. The pronounced increase in intensity (in figure 2) as well as the pronounced decrease in decay rate (in figure 3) of sample 1 could be caused by the oxygen passivation of non-radiative defects. In samples 2 and 3 containing smaller particles, the initial passivation degree is higher, therefore the ageing is expected to induce a small change both in intensity and decay rate. The data comparison from curves 2a and 2c in figure 3 reveals that the modification of emission mechanism has no effect on the decay rate as well as its energy dependence τ-1(E). This result seems to indicate that the replacement of Si-H bond by a Si-O one acting as a radiative center has no effect on the lifetime. Fig. 3. Evolution of decay rate as a funtion of emission energy from sampes after preparation, curves 1a, 2a, 3a and after exposure to air for 1-month, curves 1b, 2b. Curve 2c coresponds to sample 2 for 24 h (Bui Huy et al., 2003). Energy (eV) Intensity (a.u) After 1 month As-prepared
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  Silicon–Rich Silicon OxideThin Films Fabricated by Electro-Chemical Method 33 Figure 2 and 3 established the relation between the size of particle, intensity and decay rate during ageing. In the sample containing larger nanocrystals, the change in intensity and decay rate, i.e. the luminescence lifetime, is much larger compared with that of the smaller nanocrystals during ageing process. Figure 4 shows the evolution of PL spectra, measured at the end of an excited pulse after different exposure times. The figure reveals that the blue zone with the PL emission peaked at 470 nm is only observed after 72 hours of exposure to air. Furthermore, the figure also reveals that the PL intensity increases with increasing air exposure time. These observations differ from those reported by Volkin et al. (Wolkin et al., 1999) in which the intensity of blue emission from the as-prepared sample containing the small Si particles was shown to decrease as the exposure time increased. This result indicates that the blue-light emission observed in the present work does not originate from very small nanocrystals. Curve 4im shows the PL spectrum of a sample, which was exposed to air for 94 hours and then immersed in HF: ethanol solution. In comparing curves 4 and 4im, one can state that the blue zone in the PL spectrum observed for the sample after 94 hours of exposure to air is completely quenched. This quenching clearly relates to the fact that the silicon oxide layers in the exposed sample have been removed. The above results indicate that the intense and stable emission in the blue zone of the PL spectra observed in the considered samples relates to defects in silicon oxide layers. Wavelength (nm) Intensity Energy (eV) Fig. 4. Evolution of PL spectral measured at the end of excitation pulse from a PS sample after different exposure time (1): as-prepared, (2; 3; 4): after 26, 72 and 94 h of exposure to air, respectively, (4im): corresponding to sample exposed to air for 94 h. and then immersed in 5% HF: ethanol solution for 10 sec (Bui Huy et al., 2006). 3.2 Fabrication and characteristics of SRSO planar and active optical waveguides In this section, before elaborating on the fabrication method and properties of planar optical waveguide, active optical waveguides, and optical interference filters based on SRSO thin films we explain the method of production for the PS multilayer which forms the basis for these devices. The production of PS multilayer is possible because: (i) the etching process is self-limited (i.e. once porous layer is formed, the electrochemical etching of this layer stops); (ii) the
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  Optoelectronics - Materialsand Techniques 34 etching occurs mainly in correspondence between the pore tips; (iii) the porosity depends only upon the current density once the other etching parameters are kept fixed; and (iv) the refractive index of PS depends on its porosity (Mazzoleni Pavesi, 1995). Therefore, by varying the current density during the etching process, it is possible to vary porosity in the etching direction. In this way, the formation of a stack of PS layers of different porosities (and hence, different refractive indices) results in a dielectric multilayer. Our process for preparing an optical planar waveguide consists of 2 steps: making a PS film which contained a core layer and a cladding one, and stabilizing the waveguide structure by thermal annealing at high temperature in oxygen ambient for obtaining SRSO. In the process of fabricating an active optical waveguide, a step of deposition of Er ion into the PS film was carried out before thermal annealing. The high temperature treatment can cause an optical activation of Er ions in SRSO. The PS films were formed by electrochemical etching of 1Ωcm p-type Si wafers in 30%HF: ethanol solution. The top core layer was fabricated by applying current density of 15mA/cm2 for 7 min. The cladding was formed in the same way, with current density of 65mA/cm2 for 3 min. These conditions allowed the core and cladding to have a porosity of about 60% and 65%, respectively. For Er-deposition on the PS layer, the PS layers were immersed in an Er-content solution, and a negative bias, relative to a platinum electrode, was applied to the PS samples for a certain time so that Er-ions were drawn into the pores of the PS. In the constant current ECD method, an ErCl3–ethanol solution with an Er- concentration of 0.2 mol / liter was used, and the drift current was changed from 0.17mA.cm-2 to 0.45mA.cm-2 to obtain different Er-concentrations in the PS layers. To enable the optical activation of Er ions distributed in the pores of the PS layer, the sample was annealed at 4000C for 2 h. For making the SRSO layers, we thermally annealed the samples at 8000 - 9500C in oxygen ambient for a short time (5-30 min.) and kept samples for a longer time in nitrogen gas at 11000C. Figure 5 presents the FE-SEM image of a two-layer SRSO with different refractive indices (n core n clad), which had been controlled by using current density of 20 and 30 mA.cm-2 for the core and the cladding layers, respectively. Based on the contrast between the core and the cladding due to the difference in porosities, it is observed that the film consisted of two layers in which the core layer thickness is about 4.5 μm, and the cladding about 7 μm. The thickness of layers depended on the time duration of electrochemical process, and layers of up to tens of microns could be grown. Core Cladding Si-Substrate Fig. 5. FE-SEM image of bi-layer SRSO on a silicon substrate.
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  Silicon–Rich Silicon OxideThin Films Fabricated by Electro-Chemical Method 35 Figure 6 shows the HRSEM image of the surface of the core and cladding before (PS layer) and after (SRSO layer) of thermal annealing. As seen from Figures 6a and 6c, the difference in the density of the black area and the pores in the PS layer show that the porosity in the core layer is lower than that in the cladding. From this image we also observed the differences in density of the black area and the contrast between the black area and the white one from the PS layers (Figures 6a and 6c) and SRSO layers (Figures 6b and 6d). Those differences suggest that the treatment can cause a decrease in the size of pores and the porosity of SRSO layers. The prepared SRSO layers were dense and therefore the optical properties of the waveguides were stabilized. a clad b clad c core d core Fig. 6. FE-SEM image of surface of the core and cladding of PS layer (a, c), and of SRSO layer (b, d), which was obtained before (a, c) and after (b, d) thermal annealing in order to estimate their porosities. Samples Sample layers Refractive index Thickness (μm) Series No.1 Core layer 1.4512 5.54 Cladding layer 1.4275 3.35 Series No.2 Core layer 1.6088 1.908 Cladding layer 1.5402 6.239 Table 1. Parameters of SRSO waveguide samples The waveguide properties of the SRSO multi-layers were characterized by using M-line spectroscopy with the Prism-coupler method (Metricon 2010/M), which has the capacity to measure the thickness, the refractive index and the wave-guided modes in thin films with high accuracy (±0.0005 for index). The number of wave-guided modes in the SRSO waveguide strongly depended on the thickness of the core layer. Figure 7(a) shows that a single mode of 1310 nm in wavelength could be guided in a core layer with a thickness of 1.9 micron. The measured indices of this sample were 1.6088 and 1.5402 for the core and the cladding layers, respectively. Figure 7(b) demonstrated the measured indices and the two-mode waveguide at a 1310-nm wavelength for a core layer with a thickness of 5.54 microns. The measured indices were 1.4522 and 1.4275 for the core and the cladding, respectively. This result shows that, by changing the current density in the ECD process, we can obtain a planar layer with different indices that support the waveguide properties in the layer. The Er-ion distribution in the SRSO layer was characterized by using the EDX method with the SEM technique. The Er-ion concentration, which was doped into PS, could be controlled by using an Er-content solution
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  Optoelectronics - Materialsand Techniques 36 and by using the current density in the ECD method. For the purpose of obtaining high- concentration Er-doped SRO materials (more than 0.1 atomic % of Er) without Er clusters, which would be good candidates for planar-waveguide amplifiers, we carried out a very careful study of the distribution of Er ions along the depth of the SRSO layer. (a) (b) Fig. 7. Waveguide properties of the SRSO core/cladding layers. (a) Single-mode in the sample with core/cladding thickness of 1.90/6.24 μm and indices of 1.6088/1.5402 (b) Multi- mode in the sample with core/cladding thickness of 5.54/3.35 μm and indices of 1.4512/1.4275. Samples Type and resistivity HF concentration Er-drift current Annealing (%) (mA.cm-2) (0C) BH-10 p-type, 1 Ωcm 30 0.17 950 BH-11 p-type, 1 Ωcm 25 0.20 950 BH-12 p-type, 1 Ωcm 20 0.17 950 BH-13 p-type, 1 Ωcm 30 0.20 820 BH-14 p-type, 1 Ωcm 25 0.25 950 BH-15 p-type, 10 Ωcm 30 0.17 950 BH-16 p-type, 1 Ωcm 25 0.45 950 Table 2. Preparation conditions of Er-electrochemically-doped SRSO samples Figure 8a shows a SEM image of the Er-doped SRSO wave-guide layer prepared by using the ECD method with a drift current density of 0.2mA.cm-2. The measurement was carried out for an Er-doped SRSO thickness of 5 microns, and the Er-ion concentration was measured at points along the depth of the SRSO layer. The Er-ion concentration increased from 0.11 atom % at the top surface to 0.2 atom% at the depth of 3.5 micron from the top surface. The Er-ion concentration decreased with further increased in the depth inside the samples (see fig. 8b). For the characterization of optical properties of Er-doped SRSO layers, the Nitrogen gas laser (LN 1000, λ=337.1nm) an Argon laser (Coherent Inova 300, λ= 488nm) and a 1-W continuous laser diode (λ=976nm) were used as optical excitation sources. The
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  Silicon–Rich Silicon OxideThin Films Fabricated by Electro-Chemical Method 37 luminescent emission from the samples was collected by using two optical fibers located at 10 mm from the sample surface. The luminescence emission was analyzed by using a Jobin- Yvon spectrometer (HR460) with a multi-channel Charge-Coupled Device (CCD) detector and by Triax 320 spectrometer with a C7211 Hamamatsu CCD infrared detector for visible and infrared light, respectively. 5 4 3 2 1 0.10 0.12 0.14 0.16 0.18 0.20 0.22 Concentration Er3+ [Atom%] Thickness [ μ m] (a) (b) Fig. 8. FE-SEM image of Er-doped SRSO layers on a Si- substrate with the Er-concentration measured at points along the depth of the layer by using EDX method (a) and the Er-ion distribution inside the sample (b). The first criteria for the Er-doped SRSO samples were that they could be in both optically activated centers: the Si-nanocrystal induced visible light and the Er ion induced infrared light. Our experiment shows that the samples without thermal annealing did not emit IR light, but after thermal annealing, they strongly emitted in the 1540-nm range. This fact shows that thermal annealing at high temperatures for obtaining Er-doped SRSO layers is an important condition for optical activation of Er ions. In general, the intensity of luminescence emission at 1540 nm will be increase with increasing concentration of Er ions in the SRSO layer (Elhouichet Oueslati, 2007), when the Er-ion concentration reaches its saturation value, the luminescence intensity at 1540 nm will be decreased due to the quenching effect from Er-ion clusters (Kit Polman, 2000). Figure 9 presents the luminescence spectra at 1540 nm for samples with different drift currents from 0.17 to 0.45mA.cm-2 under excitation by 976-nm laser beam. The 1540-nm luminescence intensity of all the samples increased with increasing drift current density from 0.17 to 0.25mA.cm-2, but when the drift current density was more than 0.25mA.cm-2 the luminescence decreased slightly with increasing drift current. The evidence of energy transfer can be obtained by changing the wavelength of the excitation source. The pump at 976 nm only caused a direct excitement of Er ions (from 4I15/2 to 4I11/2 level), whereas the pump at 488 nm caused both a direct excitement of Er ions (from 4I15/2 to 4F7/2 level) and an indirect one related to the energy transfers from Si-nc to Er ions. As the effective Er excitation cross-section in the Er-doped SRSO layer is more than two orders of magnitude higher with respect to the Er resonant absorption of a photon, the pump at 976 nm causes a linear dependence of intensity on excited power (Najar et al., 2006) and the pump at 488 nm causes the non-linear one as seen in Figure 10. The photoluminescence intensity of samples irradiated by a 976-nm wavelength increased linearly with increasing excitation power when the PL emission of the sample pumped at a 488-nm wavelength has reached saturation at high power.
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  Optoelectronics - Materialsand Techniques 38 Intensity (a.u.) Wavelength (nm) 1400 1500 1600 1700 2 4 8 12 16 20 λexc = 976 nm 2 3 1 T=300K Fig. 9. Luminescence spectra from samples 1, 2 and 3 under drift current density of 0.17, 0.25, and 0.45mA cm-2, respectively. 0 200 400 600 800 1000 50 100 150 200 λExc = 488 nm λExc = 976 nm λDte = 1534 nm 100 200 300 400 500 600 Intensity [arb.units] Excitation Power [mW] Fig. 10. Dependence of luminescence intensity of Er-ions at 1534 nm on the power of the excitation laser at wavelengths of 488 nm and 976 nm. 4. Interference filters based on porous silicon and silicon-rich silicon oxide layers Interference filters based on PS were realized for the first time in the last decade (Vincent, 2004). They are formed from silicon wafers by electrochemical etching in HF solution. Compared with other methods, the electrochemical etching method avoids the difficulty associated with the stacking and assembly of dielectric layers, eliminates the need for the lengthy deposition of thick films, and permits a wide range of refractive indices to be fabricated from a single silicon substrate. PS interference filters usually formed from different dept profiles of the refractive index of PS multi-layers which act as Bragg reflectors. The optical thickness of the high- (nH) and low-refractive index (nL) layers are 1/4 of the filter wavelength, so that these structures are usually called quarter-wave-stacks (Kruger et al., 1998). The effective refractive index of PS layer is mainly determined by the porosity which can be varied by several anodization parameters. The most suitable way is changing
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  imprisonment of, iii.318-320; engagement to Margery Paston, v. 21, 39. Calle, Richard and Wm., sons of the preceding, vi. 52. Calman, Nich., iv. 197. Calthorpe, ——, i. 324; iv. 60; v. 93, 227. Calthorpe, Alianora, ii. 46. Calthorpe, Dame Elizabeth, v. 66, 175; vi. 121. Calthorpe, John, iv. 105. Calthorpe, Lady, i. 264; v. 108; vi. 57, 59. Calthorpe, Master, ii. 180, 224, 225, 240; iii. 140; iv. 107, 108; vi. 119. Calthorpe, Philip, vi. 155. Calthorpe, Sir Wm., i. 31, 285; ii. 46; iii. 266, 286; iv. 3, 266; v. 17, 90, 302. Calthorpe, William, i. 216-217, 338; ii. 308; iii. 128, 227, 276; v. 32, 179. Calthorpe, Young, iii. 6, 110. Caly, ——, ii. 157, 278. Calybut, ——, ii. 159. Camber, the, near Rye, vi. 155. Cambridge, i. 27, 116, 241, 319; ii. 40, 100, 103, 106, 107; iii. 123, 135, 215, 220; iv. 9, 21, 33, 36, 102, 241, 293; v. 72 (bis), 137, 174, 314; vi. 129, 130, 157, 206. Peter House in, ii. 49. Trinity Hall, ii. 40. Provost of St. Mary’s, iv. 87, 140. Scholar of, ii. 106. Cambridge, ... vi. 129, 130, 157, 206 Text has “vi. 129, 130, 157; vi. 206” Cambridgeshire, iv. 95, 116, 125; v. 95. [265b]
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  Camplyon, Robt., ii.313. Candle. ‘Set a candle before the devil,’ iv. 12. Cane, Dan Thomas of, ii. 21. Canon, Thos., ii. 141. Canon, Thos., the elder, iii. 23. Canterbury, i. 156, 249, 266; iii. 13, 118; iv. 90, 187, 256, 300; v. 40, 74, 75, 112, 185, 245; vi. 153, 206. Canterbury, Archbishops of. See Stafford, John (1443-52); Kemp, John (1452-4); Bourchier, Thos. (1454-86); Morton, John (1486- 1500). Canterbury, Auditor of Archbishop of, iv. 286. Court of Audience of, i. 216, 224; iv. 182. Monastery of St. Augustine’s, iv. 90. Canterbury (Caunterbury), Thos., iv. 279. Canwike, Laurence, vi. 207. Canyngs, Wm., iii. 126. Capell. See Cabell. Capgrave, ——, i. 109. Cappe, John, iii. 237. Capron, ——, ii. 102. Capron, R., v. 149. Carbalde, John, v. 184. Carbonell, Sir John, ii. 11, 46. Carbonell, Margery, ii. 11. Carbonel, Sir Richard, i. 30. Carbrooke, Norfolk, v. 168. Carbrooke, Master of, iii. 261; iv. 41. Cards, Game of, vi. 79. Carentan (Karentine), Normandy, ii. 6. Carenton, John, Bailiff of Maldon, v. 148.
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  Carew, Sir William,vi. 102, 187. Carleton, Norfolk, v. 299; vi. 75, 191, 192, 195. Carleton, by Snaith, Yorks, ii. 19. Carlisle, iv. 50. Siege of, i. 205; iii. 276. Carlisle, Bishop of, i. 336. Carpentarius, Robertus, ii. 3. Carr, John, vi. 173. Carre, Sir Wm., v. 105. Carre, Wm., vi. 155. Carrow, Prioress of, iv. 227. Carsell, France, ii. 7. Carvel, a kind of ship, iii. 306. Carvil, Edmund, iv. 196. Caryour, Roger, ii. 293. Case, John, iv. 178. Casewick, Norfolk, vi. 190. Cassell, v. 297. Cassibelaun, vi. 65. Castile, Henry IV., King of, iv. 35. See also Philip. Castile, Queens of. See Isabella and Joan. See Isabella and Joan. Isabella listed under Spain and as “Ferdinand and Isabella”. Castillon, in Perigord, i. 122, 125; ii. 123 note 1. Castleacre, Norfolk, ii. 64, 215. Castlecombe, Wilts., i. 152; iii. 126. Parson of. See Howes, Thos. 266
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  Entry spelled “Howys”. CastleRising, i. 27. Caston, Norfolk, ii. 11. Caston, Wm., Esq., ii. 46. Castre. See Caister. Catelyn, ——, iv. 256. Cates, John, ii. 60. Catesby, John, Sergeant-at-law, iv. 88, 105, 218. Catherine (Kateryne), Mrs., ii. 131; v. 161; vi. 45. Catherine of Arragon, i. 317; vi. 161. Cattes, Norfolk, iv. 292, 295, 296; v. 76, 94. Caudebec (Galdebek), France, i. 56, 106; ii. 6. Caulx, i. 106. Caus, Richard, vi. 94. Causton. See Cawston. Causton, Richard de, ii. 10, 31. Causton, Richard, Vicar of Paston, vi. 189 (bis). Cavendish, Suffolk, ii. 315. Cavendish, ——, ii. 315. Cawston (Causton), Norfolk, ii. 93; iv. 161, 207, 256. Caxton, Wm., i. 153, 323. Celot, M. J., ii. 218. Chaderton, Edmund, vi. 93. Chafeveleyns, javelins, iv. 266. Chalais in Perigord, i. 122. Chalers, Sir John, ii. 307. Chamberlain, My Lord, iv. 275; v. 62, 148, 149, 166, 178, 247, 251, 252, 290; vi. 20.
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  Chamberlain, Sir Robt.,iv. 60; vi. 2. Chamberlain, Lady, iv. 82. Chamberlain, Sir Roger, i. 74, 164; ii. 75, 185 note 1; iii. 34, 36 note 1, 38, 76; iv. 73; v. 309. Chamberlain, Sir William, i. 338; ii. 185; iii. 278; iv. 2, 18. Chamberlain, W., iii. 196, 227. Chambers (or Chambre), Eleanor, ii. 51, 52. Chambre, John, ii. 52. Chambre, Thos., ii. 52, 84, 311; iv. 279. Chambrois (Chamberexs) Castle, France, i. 56; ii. 6. Chamer, George of, ii. 312. Champagne, France, i. 13. Chancehose, Nic., ii. 3. Chancellor, My Lord, i. 60, 62, 63; ii. 20, 205, 207, 208, 210; iv. 272, 273, 274, 310; vi. 20. Chancellors, Lord. See Beaufort, Henry (1413-17 and 1424-6); Kemp, John (1426-32 and 1450-4); Stafford, John (1432-50); Salisbury, Rich. Nevill, Earl of (1454-5); Bourchier, Thos. (1455- 6); Waynflete, Wm. (1456-60); Nevill, George (1460-7 and 1470-1); Stillington, Robt. (1467-73); Rotherham, Thos. (1474- 83); Alcock, John (1485-6); Morton, John (1487-1500). Chancery, ii. 227; v. 314. Chapell, John, ii. 65. Chapelleyn, Robert, i. 331. Chapman, Sir John, Rector of St. Mary’s, Melton, iv. 306; vi. 5. Chapman, John, v. 56. Chapman, Thos., iv. 47-48. Chardequeyns. Preserve of quinces, ii. 276. Charles II., i. 27. Charles VII., King of France, i. 48-57, 95, 109, 120-125; ii. 121-127; iii. 125, 234, 306. [266b]
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  Charles VIII., vi.82, 83. Charles the Bold, Duke of Burgundy (1467-77), i. 107, 239, 247, 259, 277, 288, 289, 299; iv. 277, 296, 299; v. 175, 184, 219, 225, 258, 270, 287. Charles, Ric., ii. 95, 282; iv. 154, 155, 228, 229; vi. 180. Charles, Thos., vi. 143. Charles and his fellowship, ii. 248. Charleton, Sir Thos., i. 137. Charolois, Chas., Count of, i. 247. Chatterley, ——, i. 336. Chaucer, Geoffrey, i. 323. Chaucer (or Chawsers), Thos., Chief Butler to Henry V., and son of the Poet, i. 39, 41, 42; ii. 29, 31. Chaumpeneys, Ric., ii. 3. Chawsers. See Chaucer. Chawson, Brittany, vi. 126. Checkroll, ii. 181. Chedworth, John, Bishop of Lincoln (1451-71), i. 127, 336; iii. 47-48. Cheke, Sir Wm., vi. 152. Chelmsford, Essex, i. 171; vi. 121, 122, 181, 182. Chelsfeld, Kent, iv. 310. Chelsea (Chelchyche), vi. 108. Cheney, Sir John, vi. 118, 187. Chenne. See Sheen. Cherbourg, i. 57, 67, 175; ii. 6, 162, 291. Cherche. See Chirch. Cheseman, ——, iv. 18 (bis). Chesham, Thos., ii. 69. Cheshire, iv. 96. Chess, Game of, vi. 66. 267
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  Chesten, Henry, iv.150. Chester, i. 81, 171, 173, 190; iii. 92, 108, 233. Chester, Bishop of. See Coventry and Lichfield. Chesterfield, Yorkshire, v. 70. Chesteyne, Friar John, iv. 257. ‘Chevers Mater,’ ii. 134. Chevissche. To procure, or provide, v. 48. Chevyssance. Borrowing at interest, iv. 267. Cheynour, John, iv. 280. Cheyny, ——, v. 32. Cheyny, Sir Robt., vi. 102, 187. Chicheley, Henry, Archbishop of Canterbury (1414-43), ii. 27. Chichester, i. 75. Chief Justice, ii. 20. Child (Chylde), ——, iv. 282. Child, Richard, ii. 293. Child, Thos., vi. 15, 16. Child, Wm., ii. 4. Chinon, France, i. 54. Chippenham, Master, iv. 268. Chirch, Nicholas, ii. 271, 311; iv. 237, 242, 243, 284. Chirch, Roger, i. 113, 114, 115; ii. 263-264, 266, 267, 270, 271 (bis), 272, 274, 308, 312. Chirch, Robert, ii. 311. Chitock’s Son, iv. 181. Choke, Richard, Serjeant-at-law, afterwards Judge of the Common Pleas, iii. 84. Chrewys, ——, iv. 12. ‘Christopher, The,’ ship, v. 190. Christopher, St., Life of, vi. 66.
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  Christopher, iii. 106. Church.See Chirch. Church Close in Paston, v. 209. Chyllins, Richard, iv. 158. Chynne. To split, crack, ii. 19. Cicero (Tully), de Amicitia and de Senectute, i. 153; vi. 66. de Amicitia and de Senectute Text has “de Amicitia et de Senectute” but works are named separately in primary text. Cinnamon, v. 119. Cinque Ports, i. 75. Clarence, George, Duke of, brother of Edward IV., i. 212, 248-249, 252, 253, 257, 271-272, 281, 313; iii. 233; iv. 25; v. 35, 43, 44, 46, 47, 55, 63, 70, 71, 80, 96, 135, 188, 189, 195, 199, 270. See also George. Clarence, Lionel, Duke of, i. 191. Clarence, Thos., Duke of, brother of Henry V., ii. 8; iii. 61, 63, 148, 149. Clarendom, i. 130. Clark, Thos., vi. 157. Claryngton, John, vi. 19. Claverying, John de, ii. 3. Claxton, Norfolk, i. 112; ii. 260; iv. 161, 163. Clay, John, iii. 233. Clay, Thos., iii. 199. Clay. See Cley. Claymond, Thos., i. 338. Clederro, ——, vi. 34-35. Clement, Mr., iv. 123.
• 64.
  Clement, Master (Felmingham?), iii. 211. Clement, John, vi. 28. Clement, Robt., ii. 305. ‘Clenefyed, strek met, and non inmet,’ iv. 130. Clerc, John, ii. 188, 191. Clere, ——, ii. 190; vi. 64. Clere, Anne, Wife of Robt., vi. 149, 150. Clere, Edmund or Esmond, i. 157; ii. 30, 51, 57, 67, 68, 307; iii. 13- 14, 126, 136, 246; iv. 105, 116, 126; vi. 33, 192, 193, 198. Clere, Eliz., widow of Robt., i. 155, 291, 324; ii. 99, 108, 109-111, 140 (bis), 141, 223, 224, 278, 280, 284-285, 286, 330; iii. 3, 264; iv. 13, 27, 61, 66, 83, 105, 180, 258; v. 93, 118, 125, 208, 304-305; vi. 108, 150. Clere, John, of Ormesby, ii. 30. Clere, Margaret, iv. 61. Clere, Mistress, ii. 83. Clere, Robt., of Ormesby, ii. 30, 39, 51, 57, 64-65, 68; vi. 192, 193, 198. Clere, Robt., Son of preceding (knighted 1494), v. 220, 221, 222, 227, 232-233, 324; vi. 108, 149, 152, 155, 167. Clere, Wm., ii. 85. Clerk, ——, vi. 166. Clerk, Geoffrey, iv. 9. Clerk, Henry, iv. 237, 242, 284. Clerk, John, iii. 141; iv. 237, 244, 269. Clerk, Margaret, iv. 20. Clerk, Roger, iii. 307. Clerke, Robt., iv. 279; v. 105. Clerke, Wm., iv. 280. Cleton, Wm., i. 148. Clewer (Cleworth), Windsor, i. 317; vi. 172.
• 65.
  Cley, Norfolk, vi.132, 200, 201. Cley, Sir John, iii. 312; iv. 83; v. 105. Cleyhall, iii. 263. Clifford, ——, heirs of, ii. 253, 254. Clifford, Henry, fourteenth Lord (1485-1523), vi. 151, 187. Clifford, John, twelfth Lord (1422-55), i. 110, 134, 161; ii. 296; iii. 25, 27, 28, 29, 30, 31. Clifford, John, thirteenth Lord (1455-61), i. 177, 178, 193-194; iii. 126, 267, 268, 276. Clifford, Sir Robt., vi. 118, 187. Clifton, Sir Jerveys, of York, v. 104; vi. 151. Clifton, Lady, ii. 275. Clifton, Sir John, iii. 93, 94. Clifton, Robt., ii. 217. Clifton, Sir Robt., Kt., ii. 14; iii. 12; vi. 102. Clinton, John, fifth Lord, i. 161; iii. 30, 199; iv. 57. Clippesby, Catherine, i. 307, 310; vi. 23, 25. Clippesby, John, ii. 308; vi. 24-25. Clippesby, John, of Oby, vi. 24, 164. Clippesby, Will., ii. 271. Clippesby’s Wife, ii. 95. Clist Heath, near Exeter, i. 165. Clopton, Mr., vi. 129, 130. Clopton, ——, v. 89, 274. Clopton’s Manor, Suffolk, vi. 80. Clopton, John, Sheriff of Norfolk, i. 337; ii. 258, 260, 261, 263-264, 314-315 (bis), 316; iii. 106; iv. 105. Clopton, Wm., i. 87; ii. 315-316. Clopton, Sir Wm., vi. 190, 191. Clot shon. Clouted or patched shoes, iv. 67. 268
• 66.
  Cloth of gold,v. 245, 246. Cloves, v. 119. Clowt up. Patch up, v. 9. Clugny, France, ii. 20. Clugny, Abbot of, i. 35, 36; ii. 32. Cluniac Order, ii. 32. Clyf, Wm., iii. 109. Clyfton. See Clifton. Cobbes place, ii. 91. Cobham, v. 316. Cobham, Eleanor, Wife of Humphrey, Duke of Gloucester, ii. 114 note 3. Cobham, Lord, i. 31, 98-99, 335; v. 105. Cobham, Sir Reynold, ii. 114, 115, 253. Cockermouth, iii. 269. Cockley Cley, Norfolk, ii. 215; iii. 261. Cocus, Thos., vi. 17. Coëtivy, Oliver de, i. 122. Cofe. See Cove. Cognac, i. 95. Cok, Alex., v. 98. Coke (Cook), John, ii. 52, 269, 310; vi. 19. Coke, Lincoln, iii. 133. Coke, Wm., ii. 137; v. 94. Coke (Koke), Sanders. See Kok, Sandre. Cokett (Kokett), ——, i. 283, 293; ii. 278; iv. 74; v. 197, 257, 293, 294 (bis). Cokett, John, ii. 308. Cokfeldes, Norfolk, v. 299. Cokkow, John, ii. 267. [268b]
• 67.
  Cokkow, Thos., ii.308, 312. Colby, ——, ii. 60. Colby, Henry de, ii. 3. Colbyes, Walsham, ii. 63. Colbyne, ——, iii. 234. Colchester, i. 75, 260; ii. 226, 322, 323; iii. 248; v. 84; vi. 121, 182. Cole, Jas., ii. 27. Cole, John, ii. 194, 195; iv. 57. Cole, Wm., ii. 134, 189. Cole, Wm., ‘Cole and Watkin’s (Shipdam’s) Chamber,’ iii. 182. Coleby (Colby), Norfolk, iv. 215. Colens, Wm., iv. 228. Coles, John, iv. 279. Coleshill, i. 173. Colet, Henry, v. 74, 289; vi. 77, 149, 150. Colet, ——, vi. 36. Colin, servant of Sir J. Fastolf, iii. 93. Colins, ——, ii. 228. Colinson, Robt., a priest, iii. 41. ‘Colinus Gallicus,’ i. 152 note 6; iii. 97, 213, 230, 231. Colle, v. 259. Colley Weston, Northamptonshire, vi. 158. Collys, Harry, ii. 93, 94. Colman or Coleman, John, iii. 315; iv. 93, 156; v. 97. Colman, Nich., iii. 93, 238, 246; iv. 94. Colneise, Bailiff of, iii. 90. Colneys, Carleton, vi. 75. Colson, ——, iv. 279. Colte, ——, of Newcastle, iv. 140.
• 68.
  Colte, Thos., iii.199, 249. Colton, John, i. 332. Colvyle, John, ii. 299; v. 242, 251, 260. Colyet, iv. 148, 226. Colyns, ——, iii. 214. Combesdon, Kent, iv. 309, 311. Comenaunte. A covenant, ii. 283. Comet, ‘A greet gleymyng sterre’ seen, iii. 75. Compiègne, France, iii. 58, 64. Compostella, St. James of, i. 287; v. 186, 190. Concarnewe, Brittany, vi. 126. Conches, Normandy. See Counsheux. Coningsby, Humph., vi. 205. Conorhithe, ii. 83. Constable, ——, iii. 114. Conway Castle, i. 81. Conyers, ——, vi. 2. Conyers, John, Son of Sir Robt., iii. 210. Conyers, Sir John, iii. 199. Conyers, Sir Robt., ii. 240, 275-276; iii. 128; iv. 66. Conyers, Sir Wm., i. 249. Coo, John, ii. 159. Cook. See Coke and Kook. Cook, Henry, iv. 280; v. 304, 307. Cook, Jas., ii. 317. Cook (Kook), John, iii. 90, 91; v. 6, 154, 177, 272. Cook, Robert, iv. 279. Cooke, Alderman, i. 71. Cookesay, Sir Thos., vi. 187. 269
• 69.
  Cookeson, Thurstan, v.97. Coole, Robt., ii. 187. Coomber, Ric., iii. 309. Cooting, Sir Wm. See Coting. Coppes, iv. 71. Coppini, Francesco, Bishop of Terni, i. 186. Coppyng, John, iii. 282. Copschotyn. High crested, iv. 306. Corbet, Sir Robt., v. 105. Corbett, ——, vi. 153. Corbrygg, Simon, iii. 263. Corby, or Courby, the carrier, v. 136, 212. Corbyn, Wm., iii. 28. Cork, John Paston (alias Wortes), made Bishop of, i. 33; ii. 22, 26, 27. Cornburght, Averey, iv. 47-48. Cornewaleys, ——, vi. 130. Cornewaleys, Thos., iv. 105. Cornwall, i. 281; iii. 269. Cornwall, Sir John, ii. 8. Cornwall, a gentleman of, iv. 194. Coronation of Edward IV., iii. 276, 278, 280, 286. Coroumbr (?), Yorkshire, iii. 269. Cosard, Philip, v. 91. Cossey (Costessey), Norfolk, i. 219, 221, 223, 227, 231; ii. 170, 285; iv. 137, 147, 148, 155, 158, 159, 163, 164, 170, 171, 179, 204, 205, 207. Cossey, Bailiff of, i. 218, 221, 227, 231; iv. 137, 142, 143, 148, 155, 160, 170, 171, 186, 203, 205; v. 321. Costaunce. See Coutances.
• 70.
  Costard, ——, ii.156, 158, 249. Costard, Francis, ii. 113, 172, 228; iii. 236, 282, 296. Costers, vi. 50 note 2. Costessey. See Cossey. Costid (?), ——, ii. 66. Coteler, Edw., ii. 288. Coteler, Robt. See Cutler. Coting (Cooting or Cutting), Wm., ii. 29, 102-103; iii. 263; iv. 108; v. 6, 86. Cotman, Sir Richard, v. 320. Cotton, Suffolk, i. 209, 225, 227-228, 229-231, 234, 240, 340; ii. 170; iii. 110, 308, 309, 315, 316, 317, 319; iv. 9, 52, 62, 63, 71, 72, 157, 168, 185, 191, 193, 197, 198, 199, 200, 235, 274, 292, 303; v. 5; vi. 106, 107. Cotton, Parson of, ii. 195. Cotton (Quotton), Wm., of Cambridgeshire, Receiver of the Duchy of Lancaster, ii. 209; iii. 28, 30, 31. Cotton, Mr., vi. 157. Coubitt. See Retford. Couche, John, ii. 67. Coule, Wm., ii. 31. Councils, Great, iii. 92. Counsheux (Conches, Normandy ?), i. 55; ii. 5. Courcy (Cursy), France, ii. 6. Courtenay, Sir Hugh, v. 105. Courtenay, John, iii. 307. Courtenay, Sir Philip, v. 105. Courtenay, Piers, Bishop of Exeter, vi. 81. Courtenay, Rich., Bishop of Norwich, i. 30. Courtenay. See Devonshire, Earl of. [269b]
• 71.
  Courtenays, v. 80,113. Courtonne (Curton), France, ii. 6. Cousinot, Wm., French Ambassador, ii. 121, 124. Coutances (Normandy), ii. 6. Couth. Publicly known, iii. 41. Cove (Cofe), Wm., of North Cove, Suffolk, iii. 264. Coventry, i. 75, 152, 172, 173, 182, 189, 244, 249, 250, 313; ii. 89, 90; iii. 103, 108, 198, 233, 242; Parliament at, iii. 226; Recorder of, ii. 176. Coventry and Lichfield, Bishop of, formerly called Bishop of Chester. See Booth, Wm., and Hales, John. Covine. Intrigue, ii. 271, 312. Cowhaw, Nacton, Suffolk, iii. 5, 20, 85, 90, 91. Cowling, Suffolk, iii. 155, 157. Cowntewery. Counter or desk, ii. 281. Cowper, ——, vi. 120-121. Cowper, Aleyn, iv. 280. Cragg. Stoneware, vi. 50. Craksheld, ——, v. 244, 245. Crane, ——, ii. 315; iii. 92. Crane, Alice or Alson, ii. 135, 301; iii. 40, 115; iv. 223. Crane, Eliz., vi. 206. Crane, John, ii. 148-149; iii. 31, 40 note 1. Crane, Margaret, iii. 31. Crawethorn, ——, v. 208. Cray, North Kent, iv. 309. Crayford, i. 335. Cressell, Monsieur de, iii. 306. Cressener, ——, iii. 2. 270
• 72.
  Cressingham, Norfolk, i.306; ii. 31, 159, 220; iv. 256; vi. 17, 27-28, 96, 97. Cressingham, Parson of, ii. 31, 219. Crevecœur, Ph. de, v. 297. Croft, Sir Ric., v. 105; vi. 69, 101, 187. Croile, France, ii. 7. Croker, Sir John, v. 105, 113. Crome, N., iii. 50; iv. 76, 77, 155, 156. Cromer (Crowmer), Norfolk, i. 25, 177, 254, 308, 309; ii. 19, 64, 136, 187, 244, 292-293; iii. 126, 128, 196; v. 298; vi. 24, 29. Cromer, Pers of, ii. 68. Cromes, ii. 128. Cromwell, Ralph, Lord Treasurer of England, i. 31, 59, 131, 163, 336; ii. 38, 47, 120, 190, 197, 198, 200, 275, 324; iii. 9, 30, 41, 44, 50, 52, 66, 69, 73, 156. Cromwell, Humphrey Bourchier, Lord (1461-71), iv. 8; v. 100. Cromwell, Sir Thos., v. 105. Crosseby, ——, iv. 270. Crostwight (Crostweyt), Norfolk, Parson of, iv. 229. Crowland Abbey, i. 248. Crowmer. See Cromer. Crowmer, ——, i. 72, 74; iv. 61. Crowmer, Sir Jas., v. 105. Crowmer, Robt., vi. 154-155. Croyland, Prior of, i. 66. Croxton, ——, ii. 158. Croxton (Croxeston), Audrey, iv. 224. Croxton, Thos., iv. 117. Cryne, T., vi. 55-56. Cubitt. See Kubyte.
• 73.
  Cullynge, ——, iv.10. Cumberton, Richard, iv. 279. Curde, John, iv. 22. Currants. See Raisins. Curson, Sir John, ii. 216. Curson, Philip, iv. 231. Cursy. See Courcy. Curteys, John, v. 306. Curton. See Courtonne. Curt baron, vi. 175. Custans, ——, ii. 130. Cutler (Coteler), Robt., Vicar of Caister, iv. 102, 103, 104, 124 note 3, 228, 229, 237, 238, 241, 242, 243, 285; v. 66, 67; vi. 50. Cutting. See Coting. Cuttyng, Gees, ii. 23 note 1. Cyselond. See Sisland. Da De Di Do Dr Du Dy Dacre of the North, Ralph, first Lord (1459-62), iii. 268; iv. 51. Dacre (Dakarys) of the South, Richard Fynes, Lord (1459-84), i. 193, 215; iv. 57, 61; v. 62; Thos. Fynes, Lord (1484-1534), i. 69; vi. 151. Delalaunde, Sir Thos., v. 70. Dalby, Constantine, iv. 105-106, 124. Dale, Edward, iii. 308, 316, 317, 319; iv. 303. D’Alençon, Duc, i. 106. Dalling, Norfolk, iii. 287. Dalling, Vicar of, iv. 230, 231. Dalling, ——, Under Escheator, ii. 200, 201, 212. Dalling, Robt., i. 114; ii. 267, 271, 274, 308, 311.
• 74.
  Dalling, Sir Roger,iv. 8. Dalling, Thos., ii. 271, 311. Dalling, Wm., ii. 38-39, 214; iii. 263. Dam (Damme), John, i. 44, 90; ii. 85, 95, 111-112, 128, 133, 148, 176, 179, 180, 186, 217, 223, 225, 228, 279, 315; iii. 227, 239, 266, 270, 298; iv. 15, 16, 17, 27, 48, 75, 250, 251, 252, 254; v. 308; vi. 189, 192, 193, 198, 199. Dam, Simon, iv. 280, 281; v. 196, 252. Dam, Wm., vi. 4. Damask, price of, iv. 287. Danby, Robt., iii. 122; iv. 88. Dane, John, iv. 310. Danger. ‘In his danger,’ i.e. in his debt, or responsible to him, iii. 38. Daniel, ——, vi. 35. Daniel, ——, of Gray’s Inn, iii. 296. Daniel, John, vi. 117, 119. Daniel, Thos., of Rising, Norfolk, i. 68, 94-95, 113, 114; ii. 80, 82, 86, 87, 103, 137-138, 145, 150, 172, 176, 202, 214, 216, 217- 218, 254 note 1, 255, 256, 265, 274, 296, 312; iii. 53, 60, 63, 100, 132, 226; iv. 96. Daniel, Walter, vi. 75. Danvers, Robt., Justice of the Common Pleas, ii. 162. Danvers, Master T., v. 314. Darby. See Derby. Darcort, ——, iii. 266. Darcy (Darsy), Roger, iii. 98; vi. 172. Darell, Sir Edward, vi. 102, 187. Dartford, Kent, i. 74, 99, 335-6. Dartmouth, vi. 126. Dates, v. 313. Daubeney, ——, ii. 54.
• 75.
  Daubeney, Giles, Lord(1486-1507), vi. 145, 207. Daubeney, John, i. 221, 240, 252, 253, 254, 340; iii. 297; iv. 13, 27, 41 (bis), 42-44, 64, 66, 70, 73, 79, 84, 85, 121, 123, 127, 136 (bis), 137, 139, 140, 142, 150, 152, 156, 160, 161, 162, 192, 227, 230, 266, 271, 277, 283, 299, 306, 307; v. 14, 22, 24, 25, 45, 47, 57 (bis), 61, 63, 66, 67, 97, 98, 140, 185, 302; vi. 50, 107-108. Daubeney, Wm., iii. 237, 266. Daune, Cecily, iv. 291-292. Daungell, France, ii. 7. Daunson, John, iii. 142, 284. Dauphin, The, i. 290. Dauson. See Dawson. Davers, Thos., iv. 267-268. Daverse, ——, i. 240. Davington, ii. 188. Davy, ——, ii. 93, 94; v. 78. Davy, Alice, vi. 14. Davy, Elyse, iv. 134. Davy, John, ii. 112, 209; iii. 253-254; iv. 237, 238, 239; vi. 14. Davy, John, Vicar of E. Tuddenham, iv. 102, 103, 104, 239; vi. 75. Davy, Richard, ii. 319. Davy, Thos., vi. 19, 57. Dawbeney. See Daubeney. Dawes, Thos., iv. 178. Dawnson, ——, v. 214. Dawson (Dauson), ——, iii. 307; v. 214, 215. Dawson, John, iv. 236, 241. Daye, John, v. 174. Deal, Kent, i. 316. 271
• 76.
  Dean, Forest of,i. 250. ‘Death of Arthur, The,’ i. 323. Debenham, Suffolk, ii. 322. Debenham, Gilbert, ii. 235-236, 243, 273, 275, 276, 307; iii. 35, 309, 317, 319; iv. 9, 42-43, 62, 63, 68, 69-70, 71-72, 151, 157, 192, 193, 198. Debenham, Gilbert, the younger, knighted before 1465, father of Elizabeth Brews, i. 210, 212, 214, 230, 337; iv. 43, 45, 46, 88, 94, 192, 198, 199, 274; v. 50. Debenham, Sir Gilbert, son of the preceding and uncle of Margery Paston, i. 210, 229, 230, 340; vi. 106, 118. Debretayll, Lewis, v. 32. Decies tantum, writ of, ii. 17. De Coussy, i. 123. Dedham, Essex, ii. 149, 170, 193, 194, 195; iii. 37, 56, 65, 67, 68, 94, 96, 109, 142, 301, 302; iv. 158; v. 117. Dedham Netherhall, Essex, iii. 154; iv. 236. Dedham Overhall, Essex, iv. 236. Delaber, Sir Richard, vi. 187. Delevere, Sir Richard, vi. 101. Delgate, Thos., ii. 53. Delvys, Sir John, v. 104. Denbigh, i. 192. Dene, Robt., iii. 68. Dengayn (Dyngane), Joan, iv. 86. Dengayns, Yarmouth, v. 94. Denham (Dynham), John, afterwards Lord, i. 182, 184, 185, 281; iii. 199, 203; v. 188; vi. 107, 120. Denmark, King of, vi. 136, 138. Denney (Jenney ?), ——, ii. 238. Denny, William, iv. 279. [271b]
• 77.
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