Electroluminescence I, Volume 64 (Semiconductors and Semimetals)

Electroluminescence I SEMICONDUCTORS AND SEMIMETALS Volume 64 Semiconductors and Semimetals A Treatise Edited by R. ...

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Electroluminescence I SEMICONDUCTORS AND SEMIMETALS Volume 64

Semiconductors and Semimetals A Treatise

Edited by R. K Willardson

Eicke R. Weber

CONSULTING PHYSICISTDEPARTMENT OF MATERIALS SCIENCE SPOKANE, WASHINGTON AND MINERAL ENGINEERING UNIVERSITY OF CALIFORNIA AT BERKELEY

Electroluminescence I SEMICONDUCTORS AND SEMIMETALS Volume 64

Volume Editor GERD MUELLER COMMUNICATIONS AND OPTICS RESEARCH LABORATORY

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Contents PREFACE . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . LISTOF CONTRIBUTORS . . . . . . . . . . . . . . . . . . . . . . . . . . . .

Chapter 1 Visible Light-Emitting Diodes . . . . . . . . . . . . M . G. Craford. S . A . Stockman. M . J. Peanasky. and F. A . Kish . . . . . 1. Crystal Growth . . .

. . . . . . . . . . . . 2 . Epitaxial Film Growth . .

1

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3. Wafer Fabrication Processes . . . . . . . . . . . . . . . . . . . . . 4. Packaging Issues for LEDs . . . . . . . . . . . . . . . . . . . . . . IV . LED PERFORMANCE AND RELATED TECHNOLOGIES . . . . . . . . . . . . . V . APPLICATIONS OF VISIBLELEDs . . . . . . . . . . . . . . . . . . . . . 1 . Power Signaling . . . . . . . . . . . . . . . . . . . . . . . . . . . 2. Lighting and Nlumination . . . . . . . . . . . . . . . . . . . . . . . VI. FUTURETRENDS . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1. Future LED Performance and Cost . . . . . . . . . . . . . . . . . . 2. White Light . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3. Alternative Technologies . . . . . . . . . . . . . . . . . . . . . . . 4. Market Growth . . . . . . . . . . . . . . . . . . . . . . . . . . . REFERENCES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1 4 5 12 19 19 22 30 32 32

Chapter 2 High-Eficiency AlGaInP Light-Emitting Diodes. . . . H . Chui. N . F. Gardner. P . N . Grillot. J. W. Huang. M . R. Krames. and S. A . Maranowski

49

1. HISTORICAL OVERVIEW I1. LEDBASICS . . . . 1. Material Properties 2 . LED Structures . . 111. LED TECHNOLOGY . .

. . . . . . .

ix xvii

I . INTRODUCTION . . . . . . . . . . . . . . . . . . . I1. ALGAINPMATERIAL PROPERTIES . . . . . . . . . . . . 1 . Band Structure and Lattice Constant . . . . . . . . 2. Optical Properties . . . . . . . . . . . . . . . . . 3 . Mobility and Carrier Transport . . . . . . . . . . . V

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35 36 39 40 40

42 43 44 45

50 51 51 54

57

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CONTENTS

111. ALGAINPLED DEVICE DESIGN.

. . . . . . . . . . . . . . . . . . . .

1. Introduction and Overview of Device Design . . . . . . . . . . . . . . 2. Double Heterostructure Devices . . . . . . . . . . . . . . . . . . . . 3. Active Layer Design: Multiquantum Wells . . . . . . . . . . . . . . 4. Current Spreading . . . . . . . . . . . . . . . . . . . . . . . . . . 5 . Light Extraction . . . . . . . . . . . . . . . . . . . . . . . . . . IV . EPITAXIAL GROWTH. . . . . . . . . . . . . . . . . . . . . . . . . . 1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2. Growth of AlGaInP Materials . . . . . . . . . . . . . . . . . . . . . 3. Growth of Device Structures . . . . . . . . . . . . . . . . . . . . . 4. Manufacturing Issues . . . . . . . . . . . . . . . . . . . . . . . . V . DEVICE FABRICATION AND PACKAGING . . . . . . . . . . . . . . . . . 1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2. Die Definition and Contacts . . . . . . . . . . . . . . . . . . . . . . 3. Wafer Bonding . . . . . . . . . . . . . . . . . . . . . . . . . . . 4. Dicing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5. Packaging Issues . . . . . . . . . . . . . . . . . . . . . . . . . . VI . ALGAINPDEVICE PERFORMANCE . . . . . . . . . . . . . . . . . . . . . 1. LED Eflciency and Current- Voltage Characteristics . . . . . . . . . . 2. Eficiency vs Wavelength . . . . . . . . . . . . . . . . . . . . . . . 3. Luminous Eficiency . . . . . . . . . . . . . . . . . . . . . . . . . 4. Thermal Management . . . . . . . . . . . . . . . . . . . . . . . . 5. High-Power AlGalnP LEDs . . . . . . . . . . . . . . . . . . . . . 6 . AlGalnP Degradation . . . . . . . . . . . . . . . . . . . . . . . VII . CONCLUSIONS . . . . . . . . . . . . . . . . . . . . . . . . . . . . REFERENCES. . . . . . . . . . . . . . . . . . . . . . . . . . . . .

. .

.

.

51 51 58 63 64 66 12 72 13 88 92 99 99 99 100 100 1-01 102 102 108 109 112 116 118 123 123

Chapter 3 High-Brightness Nitride-Based Visible-Light-Emitting Diodes . . . . . . . . . . . . . . . . . R. S . Kern. W. Gotz. C. H . Chen. H . Liu. R . M . Fletcher. and C. P. Kuo I . INTRODUCTION . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1. Historical Background . . . . . . . . . . . . . . . . . . . . . . . . 2. Applications and Markets f o r Gulliitm Indium Nitride Light-Emitting Diodes . . . . . . . . . . . . . . . . . . . . . . . . 11. 111-V NITRIDE CRYSTAL GROWTH . . . . . . . . . . . . . . . . . . . . 1. Nitride OMVPE Sources . . . . . . . . . . . . . . . . . . . . . . . 2. Nitride OMVPE Reactors . . . . . . . . . . . . . . . . . . . . . . 3. Substrates for Nitrides . . . . . . . . . . . . . . . . . . . . . . . . 4. Gallium Nitride Growth on Mismatched Substrates . . . . . . . . . . . . 5. Gallium Nitride p - and n-Type Doping . . . . . . . . . . . . . . . . . 6. Alloy Growih . . . . . . . . . . . . . . . . . . . . . . . . . . . . I. Selective Area Epitaxy . . . . . . . . . . . . . . . . . . . . . . . . I11. ELECTRONIC PROPERTIES AND CHARACTERIZATION OF 111-V NITRIDEFILMS . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1. Structural Properties . . . . . . . . . . . . . . . . . . . . . . . . . 2. Luminescence Properties . . . . . . . . . . . . . . . . . . . . . . . 3. Properties and Hall-Efect Characterization of Doped III- V Nitrides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

130 131 132 133 134 135 138 139 145 148 1so 152 154 157 162

vii

CONTENTS 4. Gallium Indium Nitride . . . . . . . . . . . . . . . . . . . . . . . . 5. Aluminum Gallium Nitride . . . . . . . . . . . . . . . . . . . . . . IV. NITRIDE LED DEVICE STRUCTURES AND PERFORMANCE . . . . . . . . . . 1. LED Structures . . . . . . . . . . . . . . . . . . . . . . . . . . . 2. Substrate Selection . . . . . . . . . . . . . . . . . . . . . . . . . 3. Device Design Considerations for n-Type Layers . . . . . . . . . . . . 4. Device Design Considerations for p-Type Layers . . . . . . . . . . . . 5. Design of Active Layers . . . . . . . . . . . . . . . . . . . . . . . 6. Electrostatic Discharge Sensitivity and Reliability Characteristics . . . . . 7. WhiteLEDs . . . . . . . . . . . . . . . . . . . . . . . . . . . . V . SUMMARY AND CONCLUSIONS . . . . . . . . . . . . . . . . . . . . . . REFERENCES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

. . .

.

Chapter 4 Organic LED System Considerations . . . . . . . . .

174 175 177 177 179 182 184 185 190 193 193 196

209

Yoshiharu Sat0 I. INTRODUCTION . . . . . . . . . . . . . . . . . 11. RELIABILITY ISSUES FOR PRACTICAL APPLICATION. . 1. Degradation Mechanism . . . . . . . . . . . . 2. Operation Stability . . . . . . . . . . . . . . 3. Temperature Dependence . . . . . . . . . . . 111. DRIVING SCHEME . . . . . . . . . . . . . . . . 1. Driving Method . . . . . . . . . . . . . . . . 2. Passive us Active Matrix . . . . . . . . . . . . IV . OLED PANELS . . . . . . . . . . . . . . . . . 1. Material Issues . . . . . . . . . . . . . . . . 2. Process Considerations . . . . . . . . . . . . . 3 . Monochrome Panels . . . . . . . . . . . . . . 4. Multicolor and Full-Color Panels . . . . . . . V . SUMMARY . . . . . . . . . . . . . . . . . . . . REFERENCES. . . . . . . . . . . . . . . . . . .

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. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

209 214 214 217 228 231 231 236 239 239 244 246 247 248 252

Chapter 5 Molecular Organic Light-Emitting Devices . . . . . . 255 V. BuloviC. P . E. Burrows. and S. R . Forrest I. OVERVIEW. . . . . . . . . . . . . . . . . . . 1. Historical Perspective . . . . . . . . . . . . . . . . 2. Why Make OLEDs? . . . . . . . . . . . . . . . . . 11. THEORY OF OPERATION OF ORGANIC LIGHT-EMITTING DEVICES. 1. Conventional OLED Structures . . . . . . . . . . . . . 2. Transport Properties of OLEDs: General Considerations . . . 3. Characteristics of Single-Organic-Layer OLEDs . . . . . . 4. Characteristics of Heterojunction OLEDs . . . . . . . . . 111. IMPROVED OPERATION OF ORGANIC LIGHT-EMITTING DEVICES . . 1. Modifying Emission Color and Eficiency by Doping . . . . . 2. Degradation Processes . . . . . . . . . . . . . . . . . . 3. Increasing Eflciency Using Shaped Substrates . . . . . . . 4. Flexible OLEDs . . . . . . . . . . . . . . . . . . . . . 5. Transparent OLEDs . . . . . . . . . . . . . . . . . . . 6. Inverted OLEDs . . . . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

. . . . . . . .

255 255 256 257 257 261 263 273 279 279 283 285 286 287 29 1

...

Vlll

CONTENTS

IV . COLOR-TUNABLB OLEDs I N FULL-COLOR DISPLAYS . . . 1. Side-by-Side Patterning of Red. Green. and Blue OLEDs 2. Color Pusshand Filtering of White 0LED.y . . . . . . 3 . Wavelength Down-Conversion . . . . . . . . . . . . 4 . Microcavity-Filtered White OLEDs . . . . . . . . 5 . Stacked Color-Tunable OLEDs . . . . . . . . . . . V . SUMMARY . . . . . . . . . . . . . . . . . . . . . . REFERENCES. . . . . . . . . . . . . . . . . . . . .

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INDEX . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . o h VOLUMES IN THISSERIES . . . . . . . . . . . . . . . . . . CONTENTS

295 296 291 298 298 299 300 302

301 313

Preface Electroluminescence is the direct conversion of electrical energy into radiation. This broadest definition of the term makes a sharp distinction from thermal radiation, but it leaves open the question of the material to be used. What we are discussing here is electroluminescence of solids as distinct from gases, plasmas, and, if possible at all, fluids. We will avoid the narrowing “of solids,” but take it as implicit for the purpose of this volume. The idea for these volumes was to have experts of the various subfields or technologies of electroluminescence present an account of the present status, covering the most important materials for electroluminescenceinorganic and organic. The expression “most important” always brings up the question of “for whom,” “by what,” or “when.” Admitting from the beginning that the selection is highly subjective, nevertheless, some discussion is appropriate before leaving the selection as well as the whole book to the judgment of the reader. Electroluminescence (of solids) was first described in the scientific literature by G. Destriau in 1936 ( J . Chim. Physique 33, 587). In fact, the original paper concerns photoelectroluminescence from ZnS powder embedded in a dielectric matrix under the action of a high electric field, and for many years thereafter electroluminescence (EL) was mainly attempted and executed on (inorganic) powders. In the 1960s and 1970s, a rather complicated mechanism of electric excitation of suitable dopants (centers) gained some industrial use in EL powder lamps, and in the late 1980s in displays. Even today inorganic powder EL is used in emergency guide lights in airplanes and a variety of niche applications. Thus, it was initially on the list of possible topics for this volume. However, because it has not been documented as “in strong progress” by papers at recent international conferences or product release news, it was considered as “covered” in the prior literature. As far as I understand its status, powder EL-DC as well as AC-suffers from short-term, drive-dependent degradation, and its mechanism as the proix

X

PREFACE

cesses in much more homogeneous systems such as thin films or even single crystals is not well understood.* Implicitly, the preceding paragraph identifies some of the selection criteria: Applications or new potential applications New physical concepts or improved understanding Examples of a variety of different mechanisms and materials Of course, the five different examples covered in these volumes are in different stages of development, but all are of present interest from an applications point of view. High sophistication in understanding and manufacture of inorganic light-emitting diodes (LEDs) has been reached. This is true in slightly different degrees for the classic GaHP-based red- and amberemitting phosphide devices and the relatively new GaN-based blue and green ones. One chapter is devoted to each of them, and one chapter points out the system aspects and the differences and similarities. For organic LEDs, three chapters address the two groups of materials - small molecules and polymers -implementing much the same mechanism, and one chapter describes system aspects and applications. Thin film electroluminescence has received much less attention in the literature. What has been written primarily aims at high-information-content displays, and with some $100 million/year total revenue is dwarfed compared to LEDs. However, from the point of view that a rather limited selection of materials (treated in one chapter) is able to realize the special mechanism (a second chapter), it appears interesting enough to be included. All EL devices suffer substantial difficulty in bringing the light generated in a high-refractive-indexmaterial into the outer world. Optical interference always accompanies this process, and microcavities can improve or at least influence it, as outlined in Volume 65. The extreme case of a microcavity is a laser resonator. Semiconductor lasers are applications of the electroluminescence processes, usually very much the same as in corresponding LEDs, however strongly coupled to the radiation field in the resonator. Recently, blue H and UV lasers were introduced based on AlInGaN, which by far outdated the performance of blue 11-VI compound lasers. Infrared lasers, GaAs- and InP-based ternary or quaternary compounds of the (Ga,In)(As,P) system, are now and will be even more in the near future, the backbones of digital communication. Data storage (CD, DVD) and laser printing are big markets for electroluminescent devices. However, as every volume of Semiconductors and Semimetals is *Shortly after writing this preface, 1 was confronted at the exhibition of the Meeting of the Society of Information Displays in San Jose with new powder EL products, which were developed without much publicity outside the patent literature.

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limited to a maximum number of pages, lasers have been excluded from the scope of these volumes as an explicit subject. The nine chapters in Volumes 64 and 65 review a substantial fraction of the present research and development activities undertaken worldwide and aimed at signaling, imaging, and display, but recently in an increasing proportion also on solid state lighting. This part of the palette of possible applications seemed out of reach for quite some time after the dream of “cold light” from powder electroluminescence failed in the 1970s. The advent of fairly efficient blue and green LEDs, and the dramatic increase in the efficiency of the red LEDs has opened up a completely new avenue to illumination by mixing white from the three primary colors. Even blue diodes alone can be used to excite a broadband-emitting phosphor to produce white light by complementing its own only partially absorbed emission. Readers who are especially interested in one technology or topic should stop reading here and go to the appropriate chapter. But then perhaps one should come back to this point and become interested in other topics by the short, subjective outline of the highlights provided next. These topics are certainly contained in the respective chapters, but are not necessarily immediately obvious. And again, of course, they are subjectively chosen. Light-emitting diodes have been industrially produced over the last three decades. The overall consumption since their inception has been about 50 LEDs per human. Thus, every reader can make use of about 100 LEDs, as there are certainly others using fewer. In addition, many LEDs have probably been scrapped and replaced by new and better ones. The first chapter provides some information about the tremendous improvements in LEDs achieved over the years, but one should keep in mind that a substantial fraction of today’s production of amber and red diodes is not highly efficient TS-AlGaInP. The reason is twofold-lower price and no need for high output or efficiency. Many signaling applications can function using the cheapest diode available. Efficiency in the red-amber LED is a matter of avoiding radiationless transitions, and this is essentially a matter of crystallinity. The new, quaternary 111-V compound became interesting after the discovery of the almost perfect lattice match to (available) good GaAs substrates of (Al,Ga)o,51no,5Pand the possibility to tune its bandgap-emission by the A1:In ratio. The second achievement contributes to the high efficiency and comes from a completely different direction. While the first one ramps up the internal efficiency, the transparent substrate increases the light extraction dramatically. And this transparent substrate (TS) is a replacement of the original GaAs substrate by a GaP “carrier,” which no longer must be lattice-matched, as now nothing is “growing.” Thus, the newer history of the red-amber LED is a fascinating example of improvement of both factors

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governing the overall efficiency -internal quantum efficiency and light extraction. To appreciate the achievement represented by an external power efficiency of 28% [W/W], one should keep in mind that the light extraction from a material with a refractive index > 3.0 is < 5%. The story repeats to some extent or at least in principle in the (Al,Ga,In)N case. The outcoupling is not as low and the substrate is not absorbing from the beginning, as no lattice-matching substrate is available. But there is a gift of nature: (Ga,In)N grows on (transparent) sapphire (Al,O,) substrates, in spite of a 13% mismatch, and consequently with a large dislocation density, but with a reasonable internal quantum efficiency. In addition, at least some papers claim that a drastic reduction of imperfections does not increase EL light output -one of the points that justifies hopes for tangible increases of efficiency. Thus, comparing AlGaInP and AlGaInN, the phosphide exhibits extremely high internal quantum efficiencies- up to 80%-ccombined with very low outcoupling of the light (if no special tricks are played), but the opposite is true in the nitride devices. Another difference goes relatively unnoticed: in the phosphides the dependence of efficiency upon drive is “normal,” nonradiative channels must saturate before reaching the ultimate values. In the nitrides, the efficiency drops with increasing drive current, possibly because of overflow of the quantum wells or “dots.” Rather unobserved by a large fraction of even the engineering community, the red LEDs have matured to an extent, which has enabled them to enter the traffic light market. In the San Francisco Bay area, a large-scale replacement of the maintenance-intensive incandescent bulbs, which also consume much more energy, has begun. Brake lights on automobiles and many other items will also fall to LEDs quite soon. Summing up on inorganic LEDs, the physics is “simple”: forward current driven over a p - n junction and electron-hole recombination through shallow states or in quantum wells. Suitable materials have been found that have the bandgaps in the right place for visible emission (or infrared emission in the communications area). Only single-crystal material is capable of exercising this kind of recombination radiation efficiently enough for practical use. To be manufacturable at reasonable costs, epitaxy on affordable single-crystal substrates had to be developed. But even then, only since 1996 have LEDs been considered for applications beyond the traditional indicator functions and assembled displays with a limited number of pixels. Because of this limitation of LED use, researchers looked for other forms of electroluminescence and materials that could work in amorphous or polycrystalline films. Large-area displays were the primary target of those activities. Thin film electroluminescence was a consequent extension of the

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xiii

research done on powders. Direct current and alternating current versions were tried. What survived and is a stagnating niche market of some $100 million/year over the last years is AC. Thin film EL (TFEL) is addressed in Chapters 2 (mechanisms) and 3 (materials) of Volume 65. It differs in almost all of its essential features from LEDs, inorganic as well as organic. The only similarity is the difficulty of light extraction, a negative aspect. While (radiative) recombination of free electrons with free holes is the luminescence process in an LED, and presence and transport of electrons and holes are equally important, TFEL relies on majority carriers (electrons) only, and the radiative transition is an innershell transition in a metal or a rare-earth ion. It works only in materials that allow electrons to achieve kinetic energies of some 3eV in an electric field of about 2MV/cm. The energetic electrons impact-excite the innershell electrons of the dopant, without removing them from the individual ion. The removal of the excited electron (exciton) from its ion is usually a loss process, as is the generation of electron-hole pairs by impact-ionization of valence band electrons. Internal quantum efficiencies up to 5% have been reported, which convert into external power efficiency values around 1%. This does not appear very attractive from a light generation point of view, and is not a potential method for solid state illumination. However, there are two properties that could make the mechanism a winner in the display world. One is multiplexability, which is so poor in LCDs, is no problem even at 1000 lines or HDTV. The other is the simplicity of design and manufacture of five thin films on glass, of which only the electrode stripes must be patterned. This is beatable only (eventually) by polymer organic LEDs (OLEDs). Monochrome, 19-in. diagonal XGA displays have been manufactured on a large scale. However, two problems persisted for many years: blue was difficult to generate with adequate efficiency, and the driver cost is high, as about 200 V must be used. The problem of blue deficiency is about to be solved, and by the time these volumes are published the problem may have disappeared. The driver costs might be driven down by the plasma displays’ consumption of increasing numbers of (almost) the same chips. From a scientific point of view, some unsolved questions about the special high-field transport mechanism remain, as well as interesting problems with interface-related degradation. The latter could turn out to be a model case for other EL phenomena. As easy to manufacture, but possibly more difficult to control, the advances in the EL of polymeric light emitting devices (PLEDs) have drawn even more attention to OLEDs in general. In Chapter 4 of Volume 64 the system aspects and the applications issues of small molecule and polymeric organic LEDs are addressed. Questions such as driving schemes of large arrays of organic LEDs, and their reliability, as well as the various possible schemes for multi- or full-color displays are discussed. Chapter 5 of Volume

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64 provides a rather detailed account of the status of OLEDs based on small molecules, and Chapter 1 of Volume 65 compares polymeric and molecular LEDs. The big run for organic electroluminescence started after almost two decades of stagnation, when a Kodak team published breakthrough results in 1987, based on a novel thin film structure. It was only 3 years later that a group from the University of Cambridge reported the first EL results on polymer films. After that, industrial interest centered for some time on small molecules, but now seems to lean toward polymers, and “best values” are being reported in short time sequence by both “parties” with changing top positions. The term LED used for the device structure is well justified in that the basic mechanism is an injection of electrons and holes into a heterojunction, and a radiative recombination of excitons formed from the electrons and holes. There are, however, many properties differentiating the classic LED, which some people now like to call ILED (for inorganic LED), from the OLED. The main point is the extremely low mobilities of the carriers, about six orders of magnitude lower than in the 111-V compounds. This is caused not only by the fact that the small-molecule OLEDs use amorphous films, and the polymer films cannot be regarded as crystals either; it is also influenced by the nature of the electronic transport -mainly hopping -that appreciable voltage drops over film thicknesses of some 100nm occur. There are some distinctions between the two classes of materials, and polymers are preferable from this point of view. A big advantage of the organic materials is their low refractive index, smaller than 1.7, often as low as 1.5. Thus, light extraction is much easier. Quantum efficiencies are constant over many orders of magnitude, so that efficiency is often given in candelas per ampere, which tells only part of the story. The low mobility causes appreciable increases in voltage drop with increasing current and reduces the power efficiency with drive drastically. Under these circumstances multiplexing is accompanied by a substantial efficiency penalty. This is one of the reasons why a strong tendency toward active matrix addressing developed in the community. Going to thinner films could in principle reduce the voltage drop, However, early experiments have shown that a certain distance of the recombining electron-hole pairs or excitons from the metal electrodes is necessary to maintain high radiative probability. This given, the internal quantum efficiency can approach 25%, the limit set by the fraction of singlet excitons, and triplet excitons annihilating radiationless. Attempts to increase this fraction by phosphorescence based on triplet excitons were not very successful yet. With 25% internal quantum efficiency and 25% light extraction from a material with an index of 1.5, we can expect 6.2% external quantum

PREFACE

xv

efficiency, which is much more than in the TFEL case, but much less than in many over-the-counter LED lamps. This best value has been approached mainly in the green part of the visible spectrum. Every technology seems to have a sweet spot, or better, a “sweet color”: while red is the best for LEDs, yellow is best for TFELs, and OLEDs prefer green. Blue has been critical for all of them, but very recently power efficiency values of 22% have been reached in blue LEDs. PLEDs have had some very recent progress in blue and seem now to be slightly better than OLEDs there. The differences in spectral widths are also interesting: LEDs are the narrowest emitters, with linewidths that range from 25 nm in the green-blue to lOnm in the red; OLEDs and PLEDs have the widest emission bands, seldom narrower than 100nm. TFEL devices can span the whole range from narrow-from lOnm in the case of 4f-4f transitions of rare-earth dopants-to as wide as 70nm in allowed 5d-4f transitions of Ce3+.This question of linewidth is not only interesting from a physics point of view, but is of some practical importance. Color saturation in displays depends very much on it, and sometimes as in lighting one would prefer broader emission. When talking about colors and emission bandwidth, one must keep in mind that all the structures of EL devices resemble an interference filterfilms with different refractive indices and thicknesses of the order of one wavelength of visible light are stacked above each other. This is exactly the structure now termed microcavity. Chapter 4 of Volume 65 provides a systematic and very subject oriented treatise of light emission from microcavities. What should be stressed once more and over all is the angular dependence of radiation from microcavities, which very often has not found adequate attention in the measurement of brightness or efficiency of test structures and devices. It is standard practice to measure in or into integrating spheres in LED labs. This is not as much the case in TFEL and OLED labs. And if for some OLEDs external power efficiency values above 7% are reported, the packaging and/or measuring conditions should be included in the paper. Microcavities can greatly enhance the usage of light generated inside a dielectric stack, if only a certain aperture accepts the light--and if the linewidth is not too large. Chapter 4 of Volume 65 gives some very useful examples for any of the three technologies, which can be extended easily on the basis provided. Even some of the authors have asked which is the most promising technology. So one can expect some of the readers to also ask this. The answer is inevitably not a simple one and it mainly depends on the application. It appears very likely that LEDs will go for solid state lighting (SSL). Recently formed ventures of LED manufacturers with lighting companies -Emtron and General Electric, Siemens and Osram, Hewlett-

xvi

PREFACE

Packard and Philips -underline the seriousness of the issue. And a market size well above $10 billion/year does no1 play it down. All colors are now reasonably well developed to generate “white.” Very soon the efficiency of white solid state lamps will surpass that of incandescent bulbs, but it will be quite some time- if ever before retail prices can come close. The availability of all basic colors from LEDs has enabled the production of huge displays for sports stadiums and general advertisements. As soon as the cost of these mostly hand-assembled giants comes down, we will see them more, also outside Japan. O n the other hand, the use of LEDs in small displays such as, for instance, in cellular phones might be cut back by the advent of cheaper direct multiplexed PLED or OLED devices. However, any forecast of the development of electroluminescent devices in the display sector is extremely difficult as factors other than performance criteria will play an important role. There is little doubt that PLEDs and OLEDs will conquer certain niche markets, as TFEL has done and might continue to do. The costs of large-scale manufacture of the displays have not yet been established, and they might depend not so much on the specific material. It appears certain that all technologies discussed in these volumes will obtain continued attention and funding, and especially in the arena of organic LEDs, both molecular and polymeric, the materials list is in its infancy. Finally, thanks are expressed to all the contributors for their great efforts and all the wonderful accomplishments they have brought to these volumes. Thanks are also due to the series editor, Eicke Weber, to the mentor at Academic Press, Zvi Ruder, and his staff, and to my wife, Regina, who tolerated the long hours I have worked on this volume. ~

GERDMUELLER

List of Contributors

Numbers in parentheses indicate the pages on which the authors’ contribution begins.

V. B U L O V I(255), ~ Center for Photonics and Optoelectronic Materials (POEM), Department of Electrical Engineering and the Princeton Materials Institute, Princeton University, Princeton, New Jersey P. E. BURROWS (255), Center for Photonics and Optoelectronic Materials (POEM), Department of Electrical Engineering and the Princeton Materials Institute, Princeton University, Princeton, New Jersey C. H. CHEN,(129), Optoelectronics Division, Hewlett-Packard Company, San Jose, California H. CHUI (49), Optoelectronics Division, Hewlett-Packard Company, San Jose, California M. G. CRAFORD (l), Optoelectronics Division, Hewlett-Packard Company, San Jose, California R. M, FLETCHER (129), Optoelectronics Division, Hewlett-Packard Company, San Jose, California S. R. FORREST (255), Center for Photonics and Optoelectronic Materials (POEM), Department of Electrical Engineering and the Princeton Materials Institute, Princeton University, Princeton, New Jersey N. F. GARDNER (49), Optoelectronics Division, Hewlett-Packard Company, San Jose, California W. GOTZ(1291, Optoelectronics Division, Hewlett-Packard Company, San Jose, California P. N. GRILLOT (49), Optoelectronics Division, Hewlett-Packard Company, San Jose, California J. W. HUANC(49), Optoelectronics Division, Hewlett-Packard Company, San Jose, California xvii

xviii

LISTOF CONTKIBUTORS

R. S. KERN(129), Optoelectronics Division, Hewlett-Packard Company, San Jose, California F. A. KISH (l), Optoelectronics Division, Hewlett-Packard Company, San Jose, California M. R. KRAMES(49), Optoelectronics Division, Hewlett-Packard Company, San Jose, California C. P. Kuo (1 29), Optoelectronics Division, Hewlett-Packard Company, San Jose, California H. LIU (129), Optoelectronics Division, Hewlett-Packard Company, San Jose, California S. A. MARANOWSKI (49), Optoelectronics Division, Hewlett-Packard Company, San Jose, California M. J. PEANASKY (l), Optoelectronics Division, Hewlett-Packard Company, San Jose, California YOSHIHARUSATO (209), Yokohama Research Center, Mitsubishi Chemical Corporation, Yokohama, Japan S. A. STOCKMAN (l), Optoelectronics Division, Hewlett-Packard Company, Sun Jose, California

SEMICONDUCTORS AND SEMIMETALS, VOL. 64

CHAPTER 1

Visible Light-Emitting Diodes M. G. Craford, S. A . Stockman, M. J. Peanasky, and F. A. Kish OPTOELECTRONICS DIVISION HEWLETT-PACKARD COMPANY SANJOSE,CALIFORNIA

I. HISTORICAL OVERVIEW . .

. .

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . REFERENCES.. . . . . . . . . . . . . . . . . . . . .

11. LEDBASICS.. . . . . . . . . . . . . . . . . . 1. Material Properties . . . . . . . . . . . . . . 2. LED Structures . . . . . . . . . . . . . . . . 111. LED TECHNOLOGY . . . . . . . . . . . . . . . . 1. Crystal Growth . . . . . . . . . . . . . . . . 2. Epitaxial Film Growth . . . . . . . . . . . . . 3. Wafer Fabrication Processes . . . . . . . . . . . 4. Packaging Issuesfor LEDs . . . . . . , . , . . IV. LED PERFORMANCE AND RELATED TECHNOLOGIES . . . V. APPLICATIONSOF VISIBLELEDs . . . . . . . . . . 1. Power Signaling . . . . . . . . . . . . . . . . 2. Lighting and Illumination . . . . . . . . . . . , VI. FUTURE TRENDS. . . . . . . . . . . . . . . . . 1. Future LED Performance and Cost . . . . . . . . 2. WhiteLight. . . . . . . . . . . . . . . . . . 3. Alternative Technologies . . . . . . . . . . . . 4. Market Growth . . . . . . . . . . . . . . . .

. . . . . . . . .

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . , . . . . . . . . . . . . . . . . . . . . . .

. .

. . . . . . . . . . . . . . . . .

1 4 5 12 19 19 22 30 32 32 35 36 39 40 40 42 43 44 45

I. Historical Overview

Visible light-emitting diodes (LEDs) have been available commercially since the 1960s. The first products were sold by General Electric, based on the pioneering work by Holonyak and Bevacqua (1962). In the late 1960s, LEDs became available in high volume produced by Monsanto and Hewlett-Packard, and soon followed after by other companies. The early LEDs utilized the GaAsP materials system. They were direct energy gap 1 Copyright 0 2000 by Academic Press All rights of reproduction in any form reserved ISBN 0-12-752173-9 ISSN 0080-8784/00 $3000

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M. G. CRAFORD et al.

devices, with near-band edge emission. This limited them to the red spectral region with an emission wavelength of 650 nm due to the direct-indirect energy gap transition at higher energies. Figure 1 shows the evolution of LED performance, from the early GaAsP devices, to the devices that are available today that have more than 1OOx improvement in performance with colors throughout the visible spectrum. Visible LEDs have been reviewed frequently over the past several years (Craford and Steranka, 1994; Kish, 1995; Haitz et al., 1995; Stringfellow and Craford, 1997) and the reader is referred to these references as well as other chapters in this book for additional information. In this chapter, we cover only the key issues aimed at giving a brief but self-contained overview. The first substantial improvement in performance after the work of Holonyak was the addition of the isoelectronic impurity nitrogen to GaP (Logan et al., 1968) and GaAsP (Groves et al., 1971), which enabled the fabrication of LEDs with about l o x higher performance and colors from red through green. These are indirect energy gap devices, but the presence of the nitrogen trap, which is an efficient recombination center, enables them to achieve a surprisingly high level of luminous performance. These GaP:N and GaAsP:N technologies become widely used for many applications and continue to dominate the LED business in terms of the volume of LEDs manufactured. They are the least expensive LEDs, many types selling for less than five cents in high volume, and the performance is adequate for

1960

1970

I975

1980

1%

I990

199s

2ooo

FIG. 1. Evolution of LED performance with time as new technologies are introduced. The LED performance can be seen to have increased at the rate of approximately lox per decade. The performance (in lumens per watt) is calculated by multiplying the radiometric (wallplug) efficiency of the LEDs by the luminous efficacy as determined from the CIE curve.

1 VISIBLELIGHT-EMITTING DIODES

3

indicator applications, such as on-off status, numeric readouts, and indoor message signs. Red-emitting AlGaAs LEDs with performance about 10x higher than GaP:N and GaAsP:N were developed in the 1970s and became commercially important in the 1980s (Alferov et al., 1975; Nishizawa and Suto, 1977). These devices exceed the performance of red-filtered incandescent lamps. White incandescent lamps typically have performance of 10-20 Im/W, but to achieve red light a filter is used that wastes much of the power resulting in a red source with a typical performance of only 3-51m/W. As a consequence, for the first time LEDs were competitive in outdoor applications, such as automotive stoplights, and in applications where energy savings are important, such as red traffic signals. These high-performance AlGaAs devices are direct energy gap devices, like the early GaAsP devices, and like the GaAsP devices are limited to red emission due to the direct-indirect transition. However, the AlGaAs devices are much higher in performances since AlGaAs is lattice matched to the GaAs substrate. As a result, materials with much lower defect densities can be grown, as can heterostructures that enable better carrier injection and confinement in the active region, as well as the lower absorption losses in the device structure. In the early 1990s, AlInGaP devices were introduced (Kuo et al., 1990; Sugawara et al., 1991) that are also direct bandgap, lattice-matched heterostructures grown on optically absorbing GaAs substrates (AS-AlInGaP). However, in the case of AlInGaP the direct-indirect transition occurs at a higher energy than for GaAsP and AlGaAs so high-performance devices can be grown in the red, orange, and amber (- 590 nm) spectral regions. A further improvement in AlInGaP devices was achieved in 1994 ( Kish et al., 1994) with the introduction of transparent substrate devices (TS- AlInGaP), which are fabricated by removing the GaAs substrate and replacing it with an optically transparent GaP substrate that is wafer bonded to the AlInGaP epitaxial structure. This gave an additional 2 x performance improvement. Amber AlInGaP devices have been used for applications such as highway signs, and red AlInGaP devices have been used in traffic signals and automotive applications where they have largely replaced AlGaAs due to their higher performance and better reliability. The most recent breakthrough in LED technology is the emergence of the InGaN material system. The nitride system has been investigated since the early 1970s, but the lack of an appropriate substrate and difficulties in achieving p-type bonding delayed the demonstration of high-efficiency devices. A determined effort by Akasaki and his team resulted in a lowtemperature buffer layer growth process that enabled the growth of improved quality epitaxial films even on badly lattice mismatched sapphire substrates (Amano et al., 1986), and in a process for achieving p-type

4

M. G. CRAFORD et al.

FIG. 2. The performance evolution of different colors of LEDs compared with the performance of incandescent lamps both with and without filters. We can see that in recent years LED performance exceeds even unfiltered (white) incandescent lamps.

material by annealing utilizing LEEBI (Amano et ul., 1989). Following this work, high-performance blue and green LEDs were commercialized by Nichia in late 1993, based on the work of Nakamura and coworkers (Nakamura et al., 1994a, 1994b). As a result, it is now possible to fabricate high-performance LEDs throughout the visible spectrum. The performance exceeds that of filtered (colored) incandescent lamps and in the best devices even exceeds even the performance of unfiltered white incandescent lamps. This comparison is illustrated in Fig. 2, which shows data from Fig. 1 in a linear plot along with the performance of filtered and unfiltered incandescent lamps. This figure clearly shows the revolution in LED technology that has occurred within the last decade and suggests that LEDs can be expected to have an impact on many new applications in the future, depending largely on further improvements in the performance and reductions in cost, as discussed later in this chapter.

11. LED Basics

In this section, we review some of the basic materials issues relevant to the design, manufacture, and performance of visible LEDs. This includes a brief overview of the physics and typical operating characteristics of 111-V LEDs, which have been successfully commercialized. A more detailed discussion of high-brightness AlGaInP and InGaN LED technology can be

1 VISIBLE LIGHT-EMITTING DIODES

5

found in other chapters of this book. There are also a number of books and review articles which document the evolution of other mature visible LED technologies such as GaAsPN and AlGaAs (Craford and Steranka, 1994; Bergh and Dean, 1976; Craford, 1992; Stringfellow and Craford, 1997).

1. MATERIAL PROPERTIES To manufacture a practical visible LED a number of basic materials properties must be well understood, and technical challenges must be overcome. Critical material properties include electron and hole transport, recombination mechanisms (radiative vs nonradiative), and the role of defects in these processes. All IIILV visible LEDs are p-n junction devices fabricated from epitaxial material. Practical requirements for each LED technology include availability of a suitable substrate and an epitaxial growth technique capable of reproducible growth of high-quality layers. Design issues include optimization of the internal efficiency and light extraction efficiency, given the restrictions imposed by the epitaxial growth technology and the properties of the semiconductor materials. Finally, the LED structure must be low-cost and reliable.

a. Bandgap and Lattice Constant The bandgap ( E G ) and corresponding wavelength (A) of the As-, P-, and N-based 111-V binaries and ternaries is plotted versus lattice constant (a,) in Fig. 3. The bandgap is the single most important material parameter in the design of 111-V heterostructure devices, and the lattice constant is typically the parameter that determines compatibility of different 111-V alloys in an epitaxial structure. As a result, this figure is often referred to as the “111-V roadmap.” As discussed later, the most efficient light emitters are the direct bandgap alloys (represented by solid lines in Fig. 3). The Sb-, As-, and P-based compounds are composed of the zinc blende crystalline structure. In general, the lattice constant decreases and the energy gap increases as we move “up” column V of the periodic table from the Sb-based materials (typically used in IR-emitting devices) through the Asand P-based materials (common in visible light emitters). The most common substrates for visible LEDs are GaAs and Gap, which are both manufactured in high volume predominantly by the liquid-encapsulated Czochralski (LEC) technique. “Lattice match” refers to a condition where the difference in lattice constant relative to the substrate, Aala,, is maintained low enough (typically d 0.1%)

3

I AIP

I

.

6.5

I

40.2

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AIN

AWnPtGaAs

0.6 E

a

::f

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0.9 3

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m

1.5

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53

5.5

5.7

htticeconstant,A

Q

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5

0.3

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as

I 26

32

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3.6

4.0

+

+I 4.4

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FIG. 3. Bandgap energy (EG)and corresponding wavelength (i versus ) lattice constant for 111-V materials commonly used to produce visible LEDs: (a) As- and P-based materials (zinc blende) and (b) N-based materials (wurtzite and zinc blende).

1 VISIBLELIGHT-EMITTING DIODES

7

so that the epitaxial layers remain coherently strained. When the lattice mismatch Au/a, becomes too great, lattice relaxation occurs via generation of dislocations in the epitaxial layers, leading to a degradation in electrical and optical properties (Lester et al., 1995). Direct-bandgap Al,Ga, -,As (red and near IR) and (Al,Ga, -,)o,51n,,5P (yellow-green, amber, orange, and red) layers are grown lattice matched to GaAs substrates. The covalent bonding radius is nearly identical for Ga and A1 in 111-V alloys. As a result, the bandgap in these systems can be tuned by changing the A1 content (x) in the alloy, without a significant change in the lattice constant. This enables growth of multilayer heterostructures that are lattice matched to the GaAs substrate. Indirect bandgap G a P layers can be made to emit green or red light via doping with N isoelectronic traps or Zn:O donor-acceptor pairs, and are of course lattice matched to a GaP substrate. Yellow, orange, and red Ga,As,-,P LEDs (direct gap and N-doped indirect gap) are also common, and are grown lattice mismatched to either GaAs or G a P substrates. The N-based semiconductors continue the trend of smaller lattice constant and higher bandgap. However, the most common crystal structure for N-based 111-Vs is the wurtzite structure, due in part to the stronger ionic character of the bonding between N and the column I11 elements. Growth of bulk GaN substrates is in practice not feasible due to the extremely high vapor pressure of N at the growth temperatures required for GaN. In fact, lack of a suitable substrate for growth of high-quality epitaxial layers was one of the main factors that delayed the development of useful N-based devices until the mid-1990s (Amano et al., 1986; Nakamura et al., 1994a). Today, efficient blue and green InGaN LEDs grown on insulating A1,0, and on conducting SIC are commercially available.

b. Doping and Defects

The ability to control conductivity and conduction type is critical in LED devices. Doping of 111-V materials is achieved by intentional introduction of shallow level impurities during growth, or via diffusion or ion implantation. Common n-type dopants (donors) include S, Se, Te, Ge, and Si. Column VI impurities (such as S, Se, and Te) typically reside substitutionally on the column V sublattice. The extra valence electron creates an electron state just below the conduction band edge (a “shallow” state, with Ec - ED 5 meV), and this electron is thermally excited into the conduction band, creating a free electron (concentration n), which may contribute to n-type conductivity of the material and participate in the process of e- -h+ recombination. The ionized donor atom (concentration N , ) then

-

8

M. G . CRAFORD et al.

becomes a fixed positive charge in the crystal. Column IV impurities (such as Si and Ge) may incorporate into the column I11 sublattice, where they donate an electron to the conduction band in a similar manner. Common p-type dopants (acceptors) include Cd, Zn, Mg, and Be (on the column I11 sublattice) and C and Si (on the column V sublattice). These impurities create an electron state just above (typically E , - E , 30 meV) the valence band edge, and this state will accept an electron thermally excited from the valence band, creating a free hole (concentration p ) , which may contribute to p-type conductivity of the material and participate in the process of electron-hole recombination. In some high-bandgap p-type materials such as GaN, the acceptor states are deep ( E A - E , 2 160meV >> kT) and are only partially ionized so that p << N , , making p-type doping very difficult. Note that it is possible for column IV impurities such as C, Si, and Ge to act as donors or acceptors, depending on how they are incorporated into the crystal. These elements are referred to as amphoteric impurities. The carrier mobility p (pn for electrons, p p for holes) in 111-V materials is determined mainly by the band structure, phonon scattering, and ionized impurity scattering. As a result, p depends on alloy composition, temperature (p T 3 I 2 ) , and the doping concentration (p decreases with increasing doping). Hole mobilities are lower than electron mobilities due to the large effective mass of holes relative to electrons in most semiconductors. This also results in lower resistivity ( p = l/qnp,) in n-type material than in p-type layers ( p = l/qppp). Typical room-temperature electron and hole mobilities in GaAs are pn * 5000 cm2/V s in n-type GaAs doped at 1x 1017~ m - and ~ , pp 300 cm2/V s in p-type GaAs doped at 1 x 10’’ cm-3. The carrier mobilities in inorganic semiconductor materials are approximately six to eight orders of magnitude higher than in molecular and polymer organic LED materials, where carrier hopping is the dominant conduction mechanism. In practice, all 111-V crystals contain residual impurities that act as donors and acceptors. For example, a p-type layer intentionally doped with a concentration N , of shallow acceptors may also contain a smaller residual donor concentration N,. In this case, the layer is compensated and has a reduced hole concentration of p N , - N,. The hole mobility p p will also be reduced due to the presence of additional ionized impurities. Majority carrier (electrons in n-type material and holes in p-type material) transport determines the resistivity of the layers in a 111-V LED structure. High doping concentrations (and carrier mobility) are desirable for keeping the LED series resistance low. Minority carrier transport properties (holes in n-type material and electrons in p-type material), though more difficult to measure, are also very important in p-n junction devices such as LEDs. The N

-

-

-

1 VISIBLELIGHT-EMITTING DIODES

9

minority carrier properties of interest are the minority carrier lifetime z and diffusion length L. These two parameters are related by L = (Dz)’”, where D = (kT/q)p. In this case, p is the minority carrier mobility ( p p in n-type material, p, in p-type material), which is typically lower than the majority carrier mobility. For example, the 300 K minority electron mobility in p-GaAs with p = 1 x l O I 7 cm-3 is p,, 3000 cm2/V s, while the majority electron mobility in n-GaAs with n = l x 1017cm-3 is pn = 5000cm2/Vs (Lancefield, 1996). The total carrier lifetime (typically of the order of 10 ns) places a limitation on the ultimate switching time or modulation speed of an LED, while L describes the typical distance a minority carrier will diffuse in the absence of an applied field before recombination occurs. Typical electron diffusion lengths (L,) in high quality As- and P-based LEDs are of the order of 1pm. Deep levels are electron states within the energy gap, which are separated from the conduction and valence band edges by more than -5kT These states may be associated with impurities, native point defects (such as interstitials, vacancies, and antisite defects), extended structural defects (such as threading or misfit dislocations), or surfaces. Deep level defects do not act as efficient dopants at room temperature, but may act as efficient compensating centers, thereby increasing the resistivity in doped layers. Oxygen is a particularly troublesome deep-level impurity, which is difficult to eliminate from III-V epitaxial processes. More importantly, deep-level defects adversely affect the radiative efficiency, as discussed later.

-

c. Direct vs Indirect Recombination

The most efficient materials for LEDs are direct bandgap semiconductors. A schematic representation of the band structure of a direct bandgap material is shown in Fig. 4a. In a direct gap semiconductor, the conduction band minima is in the r valley, where the electron momentum is nearly zero. Electron-hole recombination may occur radiatively or nonradiatively. Of course, both energy and momentum must be conserved in each recombination event. The probability of radiative recombination is highest in direct gap semiconductors. In a radiative recombination event, an electron from the valley recombines with a hole in the valence band. Energy is conserved in this process by creation of a photon. The photon energy E = hc/A is approximately equal to the bandgap E,. The probability of such an event is proportional to the product of the electron and hole densities. For example, in p-type material with an excess minority electron density of An << p ,

R, = An/z,

= An.B.p,

z,= (B.p)-’

Electron energy

F

0

I

\ I

EIectro~momenhm

(a) Direct bandgap semiconductor

(b) Indirect bandgap semiconductor

(c) Near the direct-indirect transition

-

FIG. 4. Simplified band structure and most likely recombination paths for (a) a direct bandgap semiconductor ( E x - E , >>kT),(b) an indirect bandgap semiconductor ( E , - Ex >>kT),and (c) a semiconductor near the direct-indirect transition (& Ex).

1 VISIBLE LIGHT-EMITTING DIODES

11

where Rr is the radiative recombination rate (cm-3 s - I ) , zr is the radiative recombination lifetime (s), p is the majority hole concentration ( ~ m - ~and ), An is the minority electron concentration (cm-3). Here B is a recombination rate constant that is typically near lo-'' cm3/s in direct gap 111-V semiconductors. This process is referred to as spontaneous emission, and results in photon emission in a random direction in most LED structures. As a result, LEDs produce incoherent radiation. This process is distinct from stimulated emission, which requires the interaction of a photon, electron, and a hole to create a replica photon with energy, direction, and phase that are identical to the original photon. Stimulated emission is the dominant radiative recombination process in semiconductor laser diodes. Electrons in the r valley may also recombine nonradiatively with holes in the valence band. This process occurs via defect-induced states (deep levels) within the bandgap. This process is known as Hall-Schockley-Read (HSR) recombination, and the nonradiative recombination rate R,, is dependent on the location of the defect state in the bandgap, as well as the kinetics of the carrier capture process (Hall, 1952; Shockley and Read, 1952). Nonradiative recombination is of course highly undesirable in LEDs, so the minimization of impurity and defect concentrations is extremely critical at all stages of the LED fabrication process. When radiative and nonradiative processes compete, the total minority carrier recombination lifetime T is expressed as 1/T = l/zr l / T n r

+

where z,, is the nonradiative recombination lifetime. Another nonradiative recombination mechanism, Auger recombination, is typically not important in visible wavelength LEDs due to the high bandgap E,. A band diagram for an indirect bandgap semiconductor is illustrated in Fig. 4a. In this case, the X valley is significantly lower in energy than the r valley. As a result, conduction band electrons will populate the X valley. Radiative recombination is extremely unlikely in this case, since conservation of momentum cannot be satisfied without annihilation of a phonon. Nonradiative HSR recombination will dominate in a typical indirect gap semiconductor (z z znr). Figure 4c shows a semiconductor near the direct-indirect transition, where both the r and X valleys have a significant electron population. In this case, only the electrons in the valley have an opportunity for radiative recombination. However, all electrons may recombine nonradiatively via trap states. As a result, the overall radiative efficiency r r a d is given by Vrad

+ Rnr) = (nl-/zr>/[(nl?/zr) = 1/C1 + (1 + n x / n r ) ( t r l z n r ) l = Rr/(Rr

+

+ nX)/znrl

12

M. G. CRAFORD et al.

+

where n = n, n,, n, is the concentration of electrons in the r valley, and nx is the concentration of electrons in the X valley. The distribution of electrons in the two conduction band valleys is given by

where N is the relative density of states in the X and r minima (typically N > lo), and Er - Ex is the energy difference between the two minima. Due to the large density of electron states in the indirect minima ( X ) relative to the direct minima (r),the population of electrons in the X minima may be very high even when the semiconductor is still technically “direct bandgap.” Thus, the efficiency of LEDs will decrease dramatically as the direct-indirect transition is approached. In the case of GaP:Zn:O and GaP:N LEDs, impurities are used to enhance the probability of radiative recombination, but these “tricks” still result in LEDs that are much less efficient than direct bandgap devices such as AIGaAs, AlInGaP, and InGaN.

2. LED STRUCTURES The first practical visible-wavelength LED was fabricated in 1962 from “bulk” GaAsP (Holonyak and Bevacqua, 1942). Since that time, advances in the physical understanding of performance limitations, together with tremendous advances in alloy crystal growth and device fabrication technology, have led to fantastic improvements in LED efficiency (see Fig. 2). The basic structure (a p-n junction diode) has remained constant. a.

Homojunction LEDs

Figure 5a shows an energy band diagram for a forward-biased p-n junction LED. This device is termed a homojunction LED since the bandgap (E,) is the same in both the p - and n-type layers. In a forwardbiased homojunction, electrons will be injected from the n-type layer into the p-type layer where they are minority carriers. This minority carrier injection process will be followed by recombination of the electron with a (majority) hole. Similarly, holes may be injected from the p-type layer into the n-type layer, where they recombine with majority electrons. A third mechanism for current flow is injection of electrons and holes into the space-charge region (the region of material very near the p - n junction which is depleted of free electrons and holes) where they may recombine. In each case, the electron-hole recombination may occur either radiatively or

13

1 VISIBLE LIGHT-EMITTING DIODES

n-type

P-tYPe

Ec EF Ev

(a) Forward-biased homojunction

8 E

c

~e

Q

e

Q

8

e

I

EF

(b) Forward-biased single heterostructure (SH)

Ec EF

(c) Forward-biased double heterostructure (DH) FIG.5. Band diagram for (a) a forward-biased homojunction LED, (b) a forward-biased single-heterostructure (SH) LED, and (c) a forward-biased double-heterostructure (DH) LED.

14

M. G. CRAFORD et al.

nonradiatively, depending on the band structure, doping, and defect concentrations. Many important elements of a practical LED are not shown in the simplified one-dimensional band diagram of Fig. 5a. Low-resistance ohmic contacts must be made to both the n- and p-type layers to create a device with a low operating voltage of V - E , = hv = Iic/A. The n-type and p-type doping levels throughout the LED must also be optimized. For example, high doping is desirable for low series resistance and a short radiative lifetime (q).However, high doping may also lead to increased defect concentrations (and a correspondingly low rnr) and excessive free-carrier absorption of the photons generated by the LED. The primary advantage of the homojunction LED structure is its ease of fabrication. It consists of only a single crystal composition. In the case of a binary such as Gap, there is no need to worry about the difficulties associated with lattice matching. This greatly simplifies the epitaxial growth process. However, the structure also has several drawbacks. First, carrier injection is controlled by the doping near the p-n junction, and occurs for both electrons and holes. As a result, it becomes critical to maintain careful control over doping concentration and defect levels on both sides of the junction. Second, there is no mechanism for confining the injected electrons or holes to the vicinity of the p-n junction, so they may diffuse a distance of up to several diffusion lengths (L,or L p ) before recombination occurs. This makes it difficult to achieve a high injected minority carrier density, and also means that defects located far from the junction may contribute to nonradiative recombination and reduced LED efficiency. Thus the “active region” (the region where radiative recombination may occur) is very thick. Third, in a direct gap semiconductor the emitted photons may be reabsorbed via band-to-band absorption before they have an opportunity to be extracted from the LED.

b. Single-Heterostructure LEDs An energy band diagram for a single-heterostructure (SH) LED is shown in Fig. 5b. In this example, a wide-bandgap n-type layer forms a p-n heterojunction with a narrow-bandgap p-type layer. The bandgap difference between the two layers is AE, = E , (n-type layer) - E , (p-type layer) = AE, AE,,. The conduction band offset A E , provides an increase in the electron injection current under forward bias, while the valence band offset AEv helps to suppress the hole injection current. As a result, the ratio of electron current to hole current is enhanced by a factor of exp(AE,/kT). Electron-hole recombination is essentially limited to only one side of

+

1 VISIBLELIGHT-EMITTING DIODES

15

the junction. In this example, the active region consists only of the p-type layer. Use of an SH enables doping levels to be optimized with very little effect on the minority carrier injection process (assuming AE, >> kT). In Fig. 5b, the n-type layer may be optimized for low resistance, and high defect concentrations will have little impact on LED efficiency since there are very few minority holes. The p-type layer doping level may be optimized solely for high radiative efficiency. Self-absorption is also reduced in this SH LED, since E , in the n-type layer is greater than the photon energy. The fabrication of an SH LED requires a heteroepitaxial growth process capable of growth of low-defect-density alloys. In addition, the p-n junction must be placed precisely at the location of the heterojunction, or the device becomes a homojunction and the advantages of the SH will be lost.

c. Double-Heterostructure LEDs An energy band diagram for a forward-biased double-heterostructure (DH) LED is shown in Fig. 5c. In this example, a wide-bandgap n-type layer forms a p-n heterojunction with a narrow-bandgap p-type layer, as in the case of the SH LED. The narrow bandgap layer again becomes the active region. In this DH, the p-type active region is also bounded by a widebandgap p-type layer. Majority holes are easily transported across this p-p heterojunction. However, minority electrons may be effectively confined to the narrow-gap active region by the large conduction band barrier formed at this heterojunction. The wide gap layers are often referred to as confinement layers. Thus, the DH affords improved control over minority carrier injection and confinement. The active region may also be made thin to reduce internal band-to-band absorption and to increase the minority carrier density. A high minority carrier concentration is advantageous in that the radiative recombination rate can be enhanced, while at the same time the nonradiative rate may be suppressed due to saturation of deep level traps. The basic DH structure is employed in virtually all high-performance visible LEDs based on AlGaAs, AIGaInP, and AlGaInN. The design of the active region varies greatly, from a thick (1- to 2-pm) homogeneous layer in red AlGaAs LEDs, to a series of thin ( < 100A) quantum wells in blue and green AlGaInN LEDs. In some cases (such as red AlGaAs and yellow-green AlGaInP), electron confinement in the active region becomes difficult due to the small electron effective mass and restrictions on the availability of high-bandgap, lattice-matched confinement layers. The epitaxial growth sequence for an optimized DH LED structure may become quite complex. The DH structure is also utilized in 111-V infrared

16

M.G. CRAFORD et al.

LEDs and laser diodes. The LED industry has contributed to, and benefited greatly from, the development of techniques for epitaxial growth of other types of 111-V heterostructure devices such as laser diodes, photodetectors, solar cells, heterojunction bipolar transistors, heterojunction FETs, etc. These techniques are described later in this chapter and in other chapters of this volume. d. Other Design Issues A number of basic LED design issues can be illustrated and understood using simple energy band diagrams. However, optimization of an LED structure requires consideration of a number of other practical issues and limitations imposed by crystal growth and device fabrication techniques. These issues include the ability to control doping, composition, thickness, and defect concentrations in the epitaxial growth process and dopant diffusion, contact formation, and etching-sawing-cleaving for device processing. Light extraction from an LED is also a challenging problem. Photons generated via spontaneous emission are emitted isotropically from the active region into the surrounding semiconductor material. They are subject to reabsorption (band-to-band or free-carrier), reflection or scattering (at interfaces and defects), and in some cases waveguiding (which is desirable in laser diodes but undesirable in most LEDs). When a photon reaches a surface of the LED, it encounters a large change in index of refraction (from n1 2 3.5 in the semiconductor to n2 z 1.5 in the clear epoxy typically used in LED packaging). If the angle of incidence is greater than the critical angle oc = sin-’(n,/n,) 25”, determined by Snell’s law, the photon will be reflected back into the LED chip, where it is again subject to loss via absorption. Thus, an LED structure is unfortunately an efficient “photon trap.” The main macroscopic features of most LEDs, including contact geometry, surface texture, physical location of the p - n junction, substrate choice, and the use of reflectors are engineered to increase the light extraction efficiency. In spite of these efforts, the light extraction efficiency (defined as Vextraction, where Vtotai = rinternal x Vextraction) is only 3-30% in most commercial LEDs. Two other powerful factors determine the commercial success of a particular LED design: cost and reliability. Cost is primarily a function of LED process complexity, process yield, and whether the process can be successfully scaled to high manufacturing volumes. The major advantage of LEDs over other electroluminescent device technologies is reliability. Although LED lifetime varies depending on material type and operating

-

1 VISIBLELIGHT-EMITTING DIODES

17

conditions, most LEDs retain a brightness of at least 50% of the initial level after 105-106 h. Even among semiconductor LED technologies, the “winner” is often determined by which technology provides the best reliability. For example, 11-VI structures based on ZnSe have been used to produce high-performance laser diodes and LEDs in the blue-green part of the visible spectrum. However, they have not provided a commercial threat to GaNbased devices due to the inherent stability advantage of the 111-V nitrides over 11-VI materials. Another example is red emitters for traffic signals, automotive brake lights, and outdoor signs. AlGaAs-based LEDs were the first to enter these markets in the late 1980s, but have since been replaced by AlGaInP LEDs due to their superior performance in high-humidity environments.

e.

Common LED Designs

Key elements of a number of common commercial LED structures are listed in Table I and shown schematically in Fig. 6. A typical homojunction GaP LED structure is illustrated in Fig. 6a. In this structure, below bandgap light is emitted in the green via isoelectronic N traps, or in the red via Z n - 0 donor-acceptor pairs. In both cases, the GaP is a transparent substrate. A GaAsP-based homojunction LED is shown in Fig. 6b. In this structure, yellow, orange, and red emission may be achieved via N traps in indirect gap alloys, or red emission may be obtained in direct gap alloys. Growth on GaAs results in an absorbing substrate LED, while growth on GaP results in a TS LED. A single-heterostructure AlGaAs LED is shown in Fig. 6c, along with a DH device in Fig. 6d. The DH LED comes in two types: AS and TS. In AS AlGaAs, the DH is grown directly on a GaAs substrate. TS AlGaAs is fabricated by first growing a thick ( > 100pm) nonabsorbing AlGaAs layer, followed by growth of the DH. The GaAs is then removed by chemical etching, leaving the thick AlGaAs layer to serve as the substrate during subsequent device processing. An AS AlGaInP LED structure is shown in Fig. 6e. In this structure, the AlGaInP DH is grown lattice-matched to a GaAs substrate. A distributed Bragg reflector (DBR) is sometimes employed to aid light extraction by reflecting a portion of the downward-emitted photons away from the absorbing GaAs and toward the top surface of the LED. Lateral current spreading and light extraction are also aided by the use of a thick p-type GaP or AlGaAs window layer. The light extraction efficiency may be roughly doubled by removing the GaAs substrate and bonding the epitaxial layers to nonabsorbing GaP to create a TS AlGaInP device, shown in Fig. 6f.

TABLE I

PRopwtr~sOF COMMERCIALLY AVAILABLEVISIBLELEDs

Peak wavelength Type

(nm)

Type/ structure

GaP GaP:N G a p :Zn :0 GaAsP GaAsP:N

555 570 700 655 585-630

Indirect/H Indirect/H Indirect/H Direct/H IndirectIH

GaP GaP Gap GaAs GaP

Yes Yes Yes No No

Low Low Low High Low

LPE LPE LPE VPE VPE

AS AIGaAs AS AlGaAs TS AlGaAs

650 650 650

Direct/SH DirectpH DirectDH

Yes Yes Yes

High High Low

LPE LPE LPE

AS AlGaInP

560-650

DirectpH

GaAs GaAs AlGaAs (epitaxial) GaAs

Yes

High

MOVPE

TS AlGaInP

560-650

Direct /DH

GaP (waferbonded)

Yes

Low

MOVPE

InGaN

470-535

DirectiDH

" U A

No

Low

MOVPE

Substrate

Lattice matched?

Internal absorption

c

cx,

The efficiency values are an estimate of typical commercial performance in 1998.

Growth technique

Typical quantum efficiency (%) at 20 mA 0.1 0.4

2 0.2

0.4 at 585 nm 0.3 at 605 nm 0.7 at 630nm

Typical luminous efficiency (lm/W) at 20mA 0.8 2.5 0.4 0.15 2 at 585nm 1 at 605nm 1 at 630nm

4

L

8 16

4 8

2.5 at 588 nm 5 at 6 0 5 m 7 at 626nm 5 at 590nm 10 at 616nm 14 at 628 nm 11 at 470 nm 8 at 505 nm 6 at 525 nm

12 at 588 nm 18 at 605 nm 14 at 626 nm 22 at 590nm 30 at 616nm 27 at 628 nm 6 at 4 7 0 m 15 at 5051x11 22 at 525nm

F-7 p-typo G8A.P or GaA3P:N

Layer (x > 0.6)

n-fypa GaP:N or GaP

n-type QaP m~b6lfBlI) (wensparent)

FIG. 6. Device structures for common LEDs: (a) GaP:N or GaPZn:O homojunction, (h) direct GaAsP or indirect GaAsPN homojunction, (c) SH absorbingsubstrate (AS) AICaAs, (d) DH AS or transparent-substrate (TS) AIGaAs, (e) DH AS AlGalnP, (f) DH TS AIGalnP, and (9) DH InGaN.

20

M. G. CRAFORD et al.

AlInGaP LEDs are described in greater detail in Chapter 2 of this volume. A simplified schematic of an InGaN DH LED is illustrated in Fig. 6g. In this structure, both the anode and cathode are on the top surface, since the sapphire substrate is insulating. Growth of highly conductive p-type GaN layers is very difficult, so current spreading is often achieved using a thin semitransparent electrode. Light extraction from this structure is very efficient due to the low absorption and low index of refraction in the nitride layers and A1,0, substrate. The emission wavelength is controlled by changing the In content in the InGaN quantum well active region. GaNbased LEDs are also described in greater detail in Chapter 3 of this volume.

111. LED Technology

This section is devoted to an overview of some basic technologies required for the manufacture of LEDs. Excellent reviews exist for each of the topic areas and the references cited will provide a more in-depth presentation of the material. The basic technology for LEDs remains the crystal growth process, hence the areas of bulk crystal growth and epitaxy are covered with the most emphasis. Following the discussion on crystal growth is a short discussion of the basic wafer processing steps that are used in LED fabrication. Finally, product packaging issues are briefly discussed.

1. CRYSTAL GROWTH Although a key technology for manufacturing LEDs is the deposition of epitaxial films, the process of LED production begins with the bulk crystal operation, where the substrate wafers are produced. Substrate technology for visible LEDs generally means the bulk crystal growth of GaAs and G a P materials. The 111-V substrate technologies are mature at this point with commodity trading of a variety of types of GaAs (2- to 6-in.-diameter) and GaP (2- to 3-in.) wafers. More recently, high-brightness LEDs based on InGaN utilize either sapphire or Sic substrates. The technology for the manufacture of sapphire and Sic is not discussed. For GaAs, the three most successful methods for producing bulk crystal are liquid-encapsulation Czochralski growth (LEC), horizontal Bridgman (HB) growth, and gradient freeze (CF). Other possible methods for producing bulk GaAs include liquid-encapsulated Kyropoulus (LEK), float zone and horizontal and vertical zone melting. Generally speaking, the methods

1 VISIBLELIGHT-EMITTING DIODES

21

are divided into nonpulling and pulling techniques. LEC and LEK are pulling techniques, while the rest are nonpulling. Zone melting methods begin with polycrystalline GaAs created by combining gallium metal and the vapor from solid arsenic at high temperatures. The polycrystalline synthesis process is performed in a sealed quartz ampoule that maintains an arsenic overpressure at about 1 atm. The arsenic vapor is generated from heating solid arsenic. The arsenic must be present in a stoichiometric excess of the gallium in the ampoule. Next, single crystal GaAs is generated from a melt of polycrystalline GaAs at 1238C. In horizontal Bridgman growth, the sealed ampoule is either moved through a three-zone furnace or the furnace is moved the length of the ampoule (Weisberg et al., 1960). In the first zone, which is maintained over the solid arsenic source at about 620"C, the arsenic vapor sublimes and diffuses to the gallium melt. The melt stoichiometry (Ga/As ratio) can be controlled by the rate of sublimation. The third zone located over the gallium is maintained at 1240°C enables the gallium melt to become saturated with arsenic. The middle zone is maintained at the temperature for crystallization of the solid GaAs, 1238°C. Crystallization proceeds from the solid seed crystal, which has been preplaced on the low-temperature side of the melt crucible, as the melt isotherm moves along the length of the saturated melt. Ingots of HB GaAs crystals conform to the shape of the crucible with the crystal orientation determined by the seed crystal. In the case of gradient freeze methods, the process proceeds with programmed cooling without movement of ampoule or furnace. The vertical gradient freeze (VGF) technique is used widely for semi-insulating and conducting substrate applications. The strength of the technique is its simplicity. The process begins with polycrystalline GaAs. The crystallization front, produced by a temperature gradient, moves across the ampoule at rates about 0.5 mm/h. VGF is generally considered to have better scale-up characteristics than HB (Szweda, 1996). The liquid-encapsulated Czochralski method involves the pulling process starting with a seed crystal that is lowered into the GaAs melt and is then raised slowly. The LEC technique is named for the liquid boron oxide (B,O,) that floats (encapsulates) on the surface of the molten GaAs and prevents the volatile arsenic from vaporizing (Mullin et al., 1968; Weiner et al., 1971). Generally, the crystal is rotated slowly to help average out temperature fluctuations in the melt. The in situ synthesis of GaAs can be achieved by vaporizing arsenic contained in a quartz cell into a crucible containing gallium encapsulated by the B,O, (Swiggard et al., 1977; AuCoin et al., 1979). The strengths of LEC is the relatively fast pull rate and and ease of wafer scale up. The pull rates are generally in the 5- to 15-mm/h range. Diameter control is achieved by careful control over temperature and

22

M. G. CRAFORD et al.

the pull rate. Continuous feedback of the ingot mass as measured by a load cell helps to control pull rates for better diameter control. The GaAs crystallization is only part of the process that dictates the properties of the resulting finished wafers. The ingot cooling rate and degree of the impurity hardening by silicon doping determine the dislocation density and the electrical properties of the final GaAs wafers (Suzuki et al., 1978). Crystals may also be doped with a variety of other n-type and p-type dopants. Typical n-type impurities are tellurium, sulfur, and siiicon; whereas, for p-type conductivity zinc is generally used. The production of GaP wafers for LEDs still represents a very large fraction of compound semiconductor substrate use. Today, the most successful method for manufacturing GaP wafers is by high-pressure LEC (Nygren, 1973). The apparatus for GaP crystal growth must accommodate the high phosphorus vapor pressure (35 atm) at the GaP melt temperature of 1470°C. Hence, the crystal pulling is accomplished inside a high-pressure vessel. GaP crystals are general doped with sulfur (intentionally) and silicon (unintentionally). GaP substrates are used for a variety of visible LEDs, including GaP:ZnO, GaP:N, and GaAsP:N (Bergh and Dean, 1976). Once single crystal ingots are produced, the next steps in the process are grinding and cropping the ingot to a uniform diameter, slicing wafers, edge profiling each wafer, and finally, lapping and polishing the wafers to a uniform thickness, and surface flatness. The process flow is illustrated in Fig. 7. Advances in slicing include the use of wire saws to reduce sawing

FIG. 7. The basic steps of the producing substrates includes the bulk crystal growth process, grinding and cropping into a flat oriented cylinder, slicing and edge profiling each wafer, and lapping and polishing.

1 VISIBLELIGHT-EMITTING DIODES

23

waste. The polishing process may include chemical-mechanical polishing. It is this polishing step that determines the final wafer quality. Once produced largely for LED applications, GaAs demand is now split fairly evenly between semi-insulating and conducting substrates (Szweda, 1996). The use of larger diameter wafer sizes (6-in. GaAs) is predominately driven by the GaAs IC manufacturers. With the relatively small LED chip size ranging from 0.010 to 0.015in. square, and with the epitaxial films exhibiting a high degree of stress and wafer bow due to lattice mismatch and thermal expansion mismatch, LED manufacturers generally have less motivation for using large wafer sizes. In addition, nonuniformity in emission wavelength and efficiency becomes more pronounced for epitaxial growth on large wafers. As a result, most LEDs are produces on 2- and 3-in. wafers.

2. EPITAXIAL FILMGROWTH Epitaxial LED structures may be produced by liquid phase epitaxy (LPE), vapor phase epitaxy (VPE), metal-organic vapor phase epitaxy (MOVPE), and molecular beam epitaxy (MBE). For commercially available visible LEDs, the first three methods are the most important. In Table I, the combination of substrate and epitaxy (epi) technology is presented to compare the performance of LEDs. Among other data, the table illustrates the wide variety of technologies that have been employed for commercially successful LEDs. a.

Liquid Phase Epitaxy

The first and most widely used production epitaxial growth technique, liquid phase epitaxy, was developed for the manufacture of GaP and AlGaAs LEDs. The reasons for its popularity are its simplicity and cost. Although not the most flexible technique for producing epi films, it has found a niche in binary systems such as GaAs, Gap, and InP and in the inherent lattice matched ternary AlGaAs material system. Basically, LPE growth involves the precipitation of material from a solution onto a substrate. For GaAs or Gap, the temperature of the saturated solution before being brought in contact with a substrate is generally 700-900°C. There are two methods that are currently being used in high-volume manufacturing. Each method relies on a temperature gradient as the driving force for epitaxial growth. In the “slow-cooling technique” first proposed by Nelson (1963), the solution or “melt” is gradually cooled and the epitaxial material precipitates out onto the substrate. In the “temperature difference method” (see Nishizawa et al., 1973, 1977), a

24

M. G. CRAFORD et al.

thermal gradient is established between the melt and the substrate and again, the material to be deposited precipitates out onto the substrate. For reviews on the growth of single-composition LPE techniques consult Nelson (1963) or Woodall et al. (1969). Device structures that require multiple layers (e.g., AlGaAs heterostructures for LEDs and lasers) can be produced using a horizontal slider configuration. A substrate is placed in a recess that can be made to slide over the required number of melts. Each melt is saturated with a different composition material and contains a different dopant concentration. The time that the growth surface is exposed during the transfer from one melt to the next is kept as small as possible to avoid the formation of oxides, which can form defects at the interfaces between the layers. LPE has several advantages over other epitaxial techniques including high deposition rates (up to 2pm/min), the use of solid or liquid source materials (no compressed, toxic gases), and no need for ultra-high-vacuum equipment. However, because it is an equilibrium growth technique, it is difficult to control the composition of some alloys where the distribution coefficients of the constituent elements vary significantly (e.g., AlInGaP). It is also difficult to reproducibly control layer thickness below 1.0pm with LPE. LPE continues to be a “workhorse technology” primarily because it is the process of choice for the production of several high-volume visible and infrared LED device types. These include green GaP:N, red GaP:Zn:O, and red AlGaAs visible LEDs as well as GaAs and AlGaAs infrared LEDs. Other reviews of the LPE growth process can be found in books by Casey and Panish (1978), Kressel and Butler (1977), and Stringfellow (1979).

b. Hydride Vapor Phase Epitaxy Vapor phase epitaxy is one of the most widely used and successful techniques to produce compound semiconductor materials. The concept involves gas-phase chemistry in an open tube configuration enabling semiconductor crystal growth several hundred degrees below its melting point. The technique has been developed in the 1960s and 1970s for the high-volume production of the first GaAsP red, orange, and yellow LEDs. One inherent advantage of VPE over LPE is the degree of flexibility for producing mixed-alloy crystals where VPE is not limited by solid-liquid equilibrium phase diagrams and offers better control over composition and doping. There are two variations of process chemistry known as hydride VPE and chloride VPE. For chloride VPE AsCI, is passed over the hot gallium melt,

1 VISIBLELIGHT-EMITTING DIODES

25

as well as being injected into a hot zone ahead of the wafers, to form the GaCl and arsenic species (DiLorenzo, 1972; Knight et al., 1965). For hydride VPE, anhydrous HC1 is used to generate GaCl from hot gallium ahead of the wafers. The ASH, or AsCl, is cracked in a hot zone in the reactor to form As, and As, compounds. The GaCl and the arsenic compounds react on the wafer to form GaAs. The sources for hydride VPE are hydride gases arsine, phosphine, ammonia, and HCI from high-pressure cylinders. The gallium source comes from pure gallium metal that resides in the reactor chamber. If AsC1, is used, it is delivered by vapor transport with hydrogen carrier gas. For LED production, hydride VPE is the technique of choice due to the ease of delivering arsine and HCI. In Fig. 8, a schematic of a large-scale production reactor is presented. Typical production scale reactors rely on a barrel configuration with a wafer capacity of more than 150 square inches. The carrier rotates to ensure uniform temperature on the wafers and between the wafers. The reactor chamber is usually constructed of quartz when heated by resistance furnaces or requires a graphite susceptor shell to re-radiate the heat generated from a RF induction-heating coil. In hydride VPE, it is important to keep the reactor walls hot to avoid condensation of the reactants.

FIG. 8. A schematic figure of a production hydride VPE reactor. High-volume reactors are capable of growth on greater than 100in.’ per batch. VPE deposition chemistry is in near equilibrium with vapor etching of the semiconductor.

26

M. G. CRAFORD et a/.

The chemistry of hydride VPE is also illustrated in Fig. 8. The gallium chloride transport reaction begins with high purity gallium that resides in the reactor. It reacts with the injected high purity, anhydrous HCl gas. HC1+ Ga,,, -+ GaCl

+ tH,

The efficient conversion of the HC1 to GaCl and H, has important implications for the downstream semiconductor deposition reaction. The deposition reaction is an equilibrium reaction. That is, semiconductor material may be deposited or etched depending on conditions immediately over the wafers. Hence, if the HC1 does not react efficiently with the gallium melt in the chamber, unreacted HCl may cause etching of the epitaxial layers rather than deposition. At higher wafer temperatures etching is favored over deposition. Typical growth temperatures are around 700°C for GaAs and above 800°C for Gap. The VPE technique is considered flexible owing to the ability to create semiconductor alloys. Through the use of mass flow controllers and the inherent VPE property of fast growth rates, many devices take advantage of grading of either the arsine or phosphine. This grading technique is important for the realization of GaAsP LEDs. For instance, the LED emission wavelength can be tailored by choosing the relative flow of ASH, and PH, to achieve the desired GaAs,P,-. composition (x). The VPE technique is very effectivein producing thick semiconductor films. In GaAsP LED structures, a thick (20- to 30-pm) graded layer is necessary for growth of relaxed lattice-mismatched alloys with low dislocation densities. A constant composition GaAsP region is then grown, and a homojunction LED structure is formed via postgrowth Zn diffusion into the n-type epitaxial film. More recently, VPE has found other uses. High-brightness AlGaInP LEDs can be fabricated by utilizing VPE to produce a thick p-type GaP film on top of an MOVPE-grown AlGaInP DH (Fletcher et al., 1991). The thick ( 350pm)GaP layer plays an important role in current spreading on the p-side of the DH and as an optical window for the extraction of photons. It also enables further fabrication of transparent substrate AlGaInP LEDs by providing a supportive film during substrate removal and wafer bonding (Kish et al., 1994). Another area where VPE development is continuing is in the production of GaN films. Very thick epitaxial GaN films have been produced for potential use as “substrates” for MOVPE growth of InGaN based LEDs and lasers. (Detchprom et al., 1993; Molnar et al., 1997; Lee et al., 1997). While VPE offers the advantage of high growth rates, between 5 and 120 pm/h, and good composition control for mixed group-V alloys, it suffers from an inability to produce abrupt interfaces. Several clever methods have

1 VISIBLE LIGHT-EMITTING DIODES

27

been used to produce nearly abrupt interfaces including switching the wafer between streams of two different gas concentrations (Beuchet, 1985). However, the technique has not been adapted for commercial applications. Another limitation of VPE is difficulty in growing aluminum-bearing materials due to the reactivity of the AlCl, source with the quartz walls and with residual oxygen in the growth chamber. Finally, VPE uses some of the most toxic industrial gases: arsine and phosphine. The use of these gases ultimately leads to the need to handle toxic materials in the downstream gas lines. By-products of the process are toxic compounds of arsenic and phosphorus.

c. Metal-Organic Vapor Phase Epitaxy Metal-organic vapor phase epitaxy has rapidly become the epi growth technique of choice among many device engineers owing to its improved flexibility over LPE and hydride VPE. MOVPE processes offer good control of the growth rate over a wide range (0.1 to 10 pm/hr) and excellent control over composition, doping, and interface abruptness. The epitaxial growth of aluminum-containing alloys is made possible by MOVPE. In fact, MOVPE has shown to be capable of producing high-quality films over a wide range of III-V and II-VI compound semiconductors (Stringfellow, 1989). The cation (group 111) elements Al, Ga, and In are supplied by precursors derived from a broad family of chemical compounds called metal alkyls. Hence, the name metal-organic vapor phase epitaxy, is used. Other names for this technique include organometallic vapor phase epitaxy as well as the acronyms based on CVD (chemical vapor deposition) such as in MOCVD. Common precursors for III-V film growth are trimethylindium (TMIn), triethylgallium (TEGa), trimethylgallium (TMGa), and trimethylaluminum (TMA1). Generally, the anion species used are the hydrides ASH,, PH,, or NH,. The earliest reported use of the technique is credited to Manasevit (1968). The MOVPE growth process is composed of a combination of chemical, thermal, hydrodynamic, and mass-transport processes. The deposition zone in a reactor is designed to have a well-controlled gas flow pattern free of eddies and recirculation cells. The mass-transport processes occur through a concentration-gradient boundary layer where the driving force is the difference between the mean precursor concentration in the flow stream and a reduced level at the semiconductor surface. This condition leads to an effect of depleting the gas stream as it flows across a wafer. Reasonable uniformity on a static wafer can be achieved by using high gas flows and

28

M. G. CRAFORD et al.

low reactor pressure (commonly 0.1 atm). For a reactor geometry that relies on a horizontal flow, the horizontal gas-source depletion effect can be countered by substrate rotation to produce highly uniform films. Typical growth temperatures in MOVPE growth of As- and P-based alloys are between 600 and 800°C. In this temperature regime, the growth rate is controlled by the delivery rate of the group I11 species to the growth front. Hence, it is called the mass-transport-limited regime and is characterized by the growth rate being fairly temperature-independent. At temperatures much lower than 600°C, the pyrolysis of the metal alkyl species may be limited and the growth regime becomes kinetically limited. The basic MOVPE system is composed of three parts: the gas manifold, the reactor chamber, and the exhaust handling system. The gas manifold delivers the hydride gas to the reactor from high-pressure cylinders and delivers the alkyl sources using a carrier gas (usually H,). The gas manifold is also designed for efficient delivery of the organometallic sources through a zero dead space, run-vent configuration that enables rapid and reproducible switching of gases injected into the reactor chamber. The decomposition of the organometallic source occurs near the heated substrate to produce indium, gallium, or aluminum species. For the group I11 precursors, the molecule thermally decomposes as it approaches the hot wafer surface. The deposition reaction goes to completion under nonequilibrium conditions (Kuech and Veuhoff, 1984).

+ CH, . In(CH,), + In(CH,) + CH, . In(CH,) + f H 2 + In + CH, In(CH,),

-+

In(CH,),

While the first two equations can be considered as gas-phase reactions, the third equation is a heterogeneous reaction, involving the presence of a surface (Stringfellow, 1989). Likewise, through the pyrolysis of the hydride gas (e.g., ASH,), the As, and As, compounds are formed. The group V source is supplied in excess of the desorption vapor pressure of the group V element at the wafer surface. It is not unusual for the ratio of the group V to group 111 species to be greater than 100 to achieve this desired condition. The ratio of the group V to I11 also affects the incorporation of intentional and unintentional impurities in the semiconductor film. The organometallic source chemicals are chosen by their vapor pressure so as to be delivered conveniently from bubblers using hydrogen and in some cases, nitrogen carrier gas (Hollfelder et al., 1997; Hardtdegen and Giannoules, 1999). In the source bubbler, the carrier gas (H, or N,) bubbles through the liquid (or flows past the solid) metal alkyl source where it forms

1 VISIBLELIGHT-EMITTING DIODES

29

a near equilibrium vapor concentration (Mason, 1987). Good control of the temperature and pressure inside the source bubblers are necessary to control the concentration of the metal alkyl in the gas stream. It is not uncommon for MOVPE systems to be built with up to 10 metal alkyl sources to enable a maximum degree of flexibility of setting up flows for delivery into the reactor. One of the most recent developments in the field of MOVPE is the introduction of production-worthy, high-volume reactors. MOVPE is no longer simply a tool for research or small-scale production. Today, highcapacity MOVPE reactors provide enormous production volume in solar cell applications, which are mostly GaAs-based structures on germanium substrates. High-brightness LEDs based on AlInGaP and InGaN are another high-volume MOVPE application. The advent of high-volume MOVPE reactors has helped to accelerate the adoption of compound semiconductor devices by meeting the tight uniformity requirements for epi growth (McKee et al., 1992; Schmitz et al., 1992). The most popular MOVPE reactor design among early researchers was the horizontal tube design since the chamber was easy to build and maintain, as shown in Fig. 9. This type of reactor is usually limited to one wafer. The tube, usually rectangular, is configured with a susceptor that is heated by rf induction, IR lamps, or a resistance heater. A key feature of the design is that the reactor walls are cooler than the cracking temperatures for the organometallic source to minimize parasitic reactions (Li and Giling, 1995). Design and modeling of commercial reactors must take into account the flow mechanics, thermal fields, mass transport, and chemical kinetics. It is important to control each of these transport and chemical processes (Jensen, 1994). Other designs that have been explored include inverted wafer tube reactors (Arai et al., 1997); up-flow (chimney) reactors that flow gas along the direction of upward convective forces; traditional barrel configurations; large-area, low-volume, vertical downward stagnant flow design; planetary design with radial gas injection and horizontal flow (Bergunde et al., 1997); and vertical downward flow reactors with high-speed rotating disk susceptor used for pumping the gas across the wafers (Thompson et al., 1997). The latter two concepts have proven to be the most scalable and commercially successful reactor designs. Both planetary and rotating disk reactors are capable of film growth over 100 in.2 of substrates per run. MOVPE has demonstrated the technical ability to tackle the most demanding optoelectronic and electronic device requirements. The scalability of the technique has been demonstrated over the 1990s with the commercialization of several reactor design concepts. MOVPE offers the ability to produce a wide array of 111-V compound semiconductor alloys with excellent control of growth rates and composition. I n situ process

30

M. G . CRAFCIRD et al.

FIG.9. A schematic of a horizontal tube, MOVPE reactor. Flexibility in design, ease of construction, and capability to be used at low pressures are some of the advantages of this design. The conditions at the liner walls (temperature) and difficulty in controlling flow patterns and depletion across the wafers are some of its drawbacks.

monitoring techniques are being developed for MOVPE to help provide feedback to the epi grower about changes in the epi growth rate, composition, surface roughness, etc. These tools will improve MOVPE even further. Application-specific features of MOVPE growth for both AlInGaP and InGaN LEDs are presented in Chapters 2 and 3, respectively.

3. WAFERFABRICATION PROCESSES Traditional LED wafer processing involves the application of metal contacts for the anode and the cathode of the device. In some homojunction LEDs, such as GaAsP by VPE, a diffusion process is required before the metallization steps to form the p-n junction. There are many considerations

1 VISIBLELIGHT-EMITTING DIODES

31

that must be satisfied in wafer fabrication to successfully produce optimum device performance from the semiconductor films. The LED chip design must allow for the maximum extraction of photons. For early GaAsP LEDs, this means proper optimization of the diffused p-n junction depth. For direct bandgap GaAsP grown on GaAs, the diffusion depth is kept shallow to minimize band-to-band absorption in the player, while for transparent substrate nitrogen doped GaAsP grown on Gap, the diffused junction may be much deeper to enable current spreading and improved extraction of photons from the side walls of the chip. This challenge of optimizing absorbing substrate designs and transparent substrate designs has been replayed in AlGaAs and AlInGaP LED developments. In Fig. 6, a variety of AS and TS LED designs are illustrated. The coupling of an efficient chip design into an efficient package is also important. The LED wafer fabrication process is relatively simple compared to other types of compound semiconductor devices. It often consists of applying a top metal contact and alloying followed by wafer thinning and the application of a back metal contact and alloying. The metal contact geometry can be formed by either shadow masking or photolithography. The wafer can then be diced into chips by a sawing or cleaving process. Each type of LED has its own wafer fabrication process flow for various reasons relating to light output performance and cost. The metal contact area on the top surface of an LED chip must be minimized to reduce shadowing effects, yet it must enable sufficient current spreading to take advantage of the entire chip area. For transparent chip designs, the fabrication of metal contacts on the substrate is equally important. Traditional designs of backside, or substrate, contacts on Gap, TS AlGaAs, and TS AlGaInP LEDs usually employ a small dot pattern to reduce absorption of photons emitted in the direction of bottom contact. For n-type contacts, gold-germanium alloys (0.5 to > 10% germanium) are common, while for p-type contacts gold-zinc (1 to 6% zinc) can be used. Alternatives to the traditional metal contacts have been developed to enable efficient current spreading, while still enabling photon extraction. These include the use of a transparent contact, such as, indium tin oxide (ITO) (Lin et al., 1994) or the use of very thin, semitransparent metal films. Although it can transmit greater than 90% visible light, the ohmic I T 0 contacts work best on GaAs. Hence, I T 0 application for visible LEDs must involve some absorptive losses under the contact if a thin layer of GaAs is employed. Thin, semitransparent metal films have been used on some InGaN LED chip designs to aid in current spreading. Another important'area of wafer processing is the wide use of etching, both wet and dry. Wet-etching processes are used for removal of native

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M. G. CRAFORD et al.

oxide layers, thinning wafers to the final die thickness, selectively removing certain semiconductor layers, and the removal of saw damage. For InGaNbased LEDs, reactive-ion etching is commonly used to expose the lower n-type film for the dual top contacts of the device when it is grown on insulating sapphire. Etching may also be used to form mesa structures that confine current and aid in extraction of photons.

4. PACKAGING ISSUESFOR LEDs Traditional LED chip packaging involves placement in a metal reflector cup, die attach, wire bonding, and encapsulation in a transparent epoxy. With the rapid improvements in LED light output performance in all colors based on AlInGaP and InGaN materials, new applications emerge that require new packaging technologies. A review of LED packaging and applications has been written by Hodapp (1997). Packaging considerations for even the simplest LEDs must include effects of the epoxy. Besides protecting the semiconductor chip and wire bond, the purpose of the epoxy is to improve the extraction of photons from the high-index semiconductor material to air. By coating the semiconductor with the epoxy, the escape cone for photons is increased. However, the epoxy can have a dramatic effect on the reliability of the diode due to a number of factors including mechanical stress that the epoxy imparts on the semiconductor chip, resistance to moisture and environmental forces, and stability to UV or visible radiation from the sun or the device. As LEDs are being driven at higher currents for high-power applications, the thermal resistance of the package can limit device performance (Hofler et al., 1998). Large chip designs are already being commercialized that dramatically reduce the number of emitters needed to compete with the performance of filtered incandescent lighting. Without proper heat sinking, the increased power dissipation may lead to excessive junction heating and decreased light output performance for most LEDs.

IV. LED Performance and Related Technologies

A suggestion has been made that the direct bandgap 111-V alloy semiconductor has the potential to be the ultimate form of lamp (Holonyak, 1995). This premise is based on its potential ability to convert electron-hole pairs to photons with 100% internal quantum efficiency and then extracting the emitted photons from the high-index semiconductor with very high effi-

1 VISIBLELIGHT-EMITTING DIODES

33

ciency. To date, infrared-emitting GaAs-based devices have best approached this potential. However, AlGaInP LEDs have made significant progress toward this goal in the visible regime with the best red-emitting devices exhibiting an external efficiency of 30% (photons/electron), corresponding to an estimated (90% internal efficiency and a -30% extraction efficiency (Gardner et al., 1999). Yellow and orange spectrum AlGaInP devices possess about two to three times lower external quantum efficiency, but exhibit some of the highest luminous efficiencies (20-70 lm/W) as a result of the increased eye sensitivity in this spectral regime. Furthermore, high luminous efficiencies have also recently been realized in the InGaN materials system, resulting in green- and blue-spectrum LED devices with efficiencies in the ranges of -20-35 and -3-71m/W, respectively (Nakamura et al., 1994a, 1995). As a result, the luminous efficiencies of both AlGaInP and InGaN LEDs are competitive with many conventional commercial lighting technologies as shown in Table 11. This comparison must be qualified in that LEDs are colored (nearly monochromatic) lighting sources, whereas the comparison technologies are generally white. However, LEDs possess a distinct efficiency advantage over filtered halogen and incandescent sources. A current drawback of LEDs is that they are low-power devices. The highest power commercial LEDs are rated at a maximum input power of -0.67 W (Hofler et al., 1998) compared to 10-10,000 W for the conventional lighting technologies of Table 11. The gap between conventional lighting sources and LEDs should further decrease as higher power LEDs are developed. Furthermore, the high efficiency available across the entire visible spectrum should enable high-efficiency, high-power “white” solid-state light emitters that are truly competitive in performance with conventional white light sources. However, this achievement will require not only the development of highly efficient, high-power colored devices, but also new techniques for integrating them into a “white” source. Such advances will require many more revolutionary and fundamental advances in the field of visible optoelectronics. Despite such challenges, solid-state emitters remain one of the most promising lighting sources for future improvement. This potential is evidenced by the fact that of all lighting technologies shown in Table 11, LEDs have by far shown the largest improvement in luminous efficiency since 1970. Undoubtedly, much more work is to be done in this field, and if successful, may indeed yield the ultimate lamp. Several different forms of solid-state emitters are currently produced. Figure 10 (adapted and updated from Moon, 1997) shows the power conversion efficiency versus peak wavelength for the most common solidstate emitter architectures: LEDs, vertical-cavity surface-emitting laser diodes (VCSELs), and edge-emitting laser diodes. Each data point represents the best-reported efficiencyunder DC room-temperature operation. As

-

34

M. G. CRAFORD et al. TABLE I1 MODERNEVOLUTION OF LIGHTING TECHNOLOGY PERFORMANCE

Lighting Technology Incandescent Halogen

Light-Emitting Diode’ Blue emitting Red, yellow, green emitting White emitting (phosphor-converted blue)

Average Commercial Efficiency (Im/W)’

Improvement since 19701

8-22 13-39

1.4 x

3-1

1.5~

30 x (all colors)

20-30 5-10 (> 70, best device)

High-pressure mercury Metal halide Fluorescent High-pressure sodium Low-pressure sodium

36-60 68-95 41-104 46-140 100-200

1.2x 1.3x 1.3 x

1.4~ 1.2x

1. Data for conventional lighting sources are based on 1993 data from Lighting Manual, 5th ed. (Phillips Lighting B. V., Eindhoven, NL, 1993), pp. 3-71. These data include the full efficiency range of di5erent lamp types and powers for each technology. Generally, the highest power (wattage) lamps exhibit the highest efficiencies. 2. Blue and green LED data are based on InGaN/Al,O, emitters and yellow and red LED data are based on transparent-substrate AlGaInP/GaP emitters. “Best device” represents the highest efficiency measured reported in the literature (Gardner et al., 1999).

shown in Fig. 10, light-emitting diodes are the only devices have been produced across the entire visible spectral regime (blue, green, yellow, orange, red emission). The availability of the full color range combined with the low cost of LED light (typically measured in lumens per dollar) makes LEDs currently the technology of choice for most solid-state lighting applications. In the short- and long-wavelength realms shown in Fig. 10, all technologies (LEDs, VCSELs, edge-emitting laser diodes) have relatively similar performance. Currently, both edge-emitting laser diodes and VCSELs are employed primarily in communication applications, Infrared devices are typically employed in free-space and fiber-optic communications applications as well as optical memory systems. Short-wavelength laser diodes are targeted primarily for use in next-generation optical memory systems. Since efficiency is not a primary distinction for these long- and short-wavelength devices, other factors (e.g., power output, beam characteristics, modulation speed, cost) are the primary determinants for their selection in a given application.

35

1 VISIBLELIGHT-EMITTING DIODES

FIG. 10. Power conversion efficiency (or “wallplug” efficiency) versus peak emission wavelength for the best reported results for various solid-state emitter technologies (LEDs, VCSELS, and edge-emitting laser diodes). All results are for DC room-temperature operation. The various emitter technologies compare favorably in terms of performance; however, only LEDs span the full spectrum of visible wavelengths (adapted from Moon, 1997).

V. Applications of Visible LEDs The earliest applications of LEDs were in uses for indicators and relatively simple low-density numeric and alphanumeric displays for both consumer and industrial applications. More recently, improvements in LED efficiency have resulted in the ability to employ these devices in an increasing number of applications. These include automotive interior and telecommunications applications (e.g, cellular phone displays and backlighting) as well as backlighting liquid crystal displays (LCDs). Furthermore, high-brightness LEDs are making significant inroads in power signaling applications, including commercial and traffic moving message panels (MMPs), traffic signals, and automotive exterior lighting. The markets for display and signaling applications can be segmented based on information density and viewing distance as shown in Fig. 11. Note that high information density and short viewing distance segment is dominated by LCDs and CRTs. Inorganic LEDs displays must be fabricated with one chip per pixel, making the realization of cost-effective, high-density displays prohibitive. Over the next 510 years, organic LED displays (wherein multiple colors can potentially be more readily integrated on a single substrate) may begin to penetrate this market. Despite the myriad of new applications that are emerging for LEDs, traditional indicator and display applications utilizing the

-

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M. G. CRAFORD et al.

Display and Signaling Market Segmentation

FIG.11. Segmentation of display and signaling applications by information density and viewing distance. Long viewing distance applications are best suited to high-brightness LEDs. Low-information density and short viewing distance applications typically employ conventional lower performance LEDs (GaAsP and Gap), although high-brightness devices are emerging in several applications (e.g., automotive interior lighting, display backlighting).

lower performance older technologies (Gap and GaAsP) dominate both in terms of shipment volumes and total revenue.

1. POWER SIGNALING u.

Electronics Signs and Signals

One of the fastest growing applications for LEDs is their application in traffic signals. The major value offered by LEDs in this market is both energy and maintenance savings resulting from the high efficiency and long lifetimes of LEDs. In the United States, a typical red LED traffic signal consumes 15 W compared to a 150 W incandescent solution (LED turn arrows consume only -5 W). Despite their higher initial cost, the energy savings offered by these devices results in a lower total cost of ownership. The development of high-efficiency red-spectrum AlGaInP emitters has reduced the time required to recoup the increased initial cost in energy savings to 1-2 yr (compared to -3-5 yr when red traffic signals based on AlGaAs emitters were first introduced). Currently, red LED traffic signals have had the highest adoption since they are operated -55-60% of the time, and hence, consume the most power. Green signals are operated -35-40% of the time, and are beginning to be converted to LEDs as the

-

1 VISIBLELIGHT-EMITTING DIODES

37

price and performance of the cyan ( N 505-nm) InGaN emitters continues to improve. The power savings for these devices (as well as other electronic highway signs) is so significant that many utility companies are offering municipalities incentives to purchase LED signals. Such initiatives reduce the total energy consumption in the region and avert the building of additional infrastructure by the utility companies. In Europe, the main value of LED traffic signals is reduced maintenance. Most existing installations utilize higher efficiency 50-W halogen sources and the signaling infrastructure requires LED traffic signal heads to dissipate 25 W (lower power dissipation is designed to indicate a failed bulb). As a result, reduced maintenance by the long-lived LEDs is the primary value proposition for solid-state emitters. Accordingly, both red and green LED signals are required before for this value can be realized. Since cost-effective green LED traffic signals are just beginning to emerge, the market adoption in Europe has been slower than that in the United States. Electronic signs (often referred to as moving message panels) are another high-growth market for high-brightness LEDs. Typically, the signs are either monochromatic (traffic or commercial signs) or full color (video). Full-color video signs generally contain millions of LEDs and are typically composed of multilamp pixels (in ratio 1 b1ue:l green:2 red lamps per pixel currently). LEDs are valued in these applications for their low energy consumption (- 10% of a bulb solution), long lifetime (> 30,000 hr compared to 1,500-8,000 hr for minibulbs), ruggedness (lack of glass envelopes and filaments), and lower maintenance costs. Furthermore, the high brightness offered by LEDs compared to bulb solutions enables better viewability in direct sunlight. LED signs also possess the additional flexibility of being capable to be operated on solar power during the day and by batteries (charged via solar power) at night.

b. Automotive Exterior Signaling High-brightness yellow- and red-emitting AlGaInP LEDs are being increasingly employed in automotive exterior lighting applications. AlGaAs emitters were first utilized on the outside of a vehicle in a center high-mount stop light (CHMSL) in 1986. Since this time, an increasing number of automotive exterior colored lighting functions have employed LEDs. Figure 12 shows the first vehicle model year that a given lighting function utilized LEDs. By the early 2000s, at least one example of LEDs are expected to be have been used for each colored lighting application on the outside of a vehicle. Market adoption has increased for these applications as a result of the introduction of high-brightness AlGaInP LEDs in 1991. These devices

38

M.G . CRAFORD et al.

FIG. 12. Time line showing the first model year that LEDs were employed in various automotive exterior lighting applications. By the early 2000s, at least one example of each colored light function on the exterior of an automobile is projected to have employed LEDs.

offer improved reliability (relative to AlGaAs emitters, especially in highhumidity environments) as well as superior cost per lumen. The automotive industry is highly sensitive to any increase in cost on a vehicle. Consequently, TS AlGaInP red- and yellow-emitting LEDs (which offer the lowest cost per lumen at the system level) are currently the technology of choice for most automotive exterior LED lighting applications. To date, the relatively high cost (dollars per lumen) of LED solutions has limited their application to high-end and luxury vehicles. As the system level pricelperformance ratio continues to be reduced over the next 5-10 yr (via performance improvements and cost reduction of the LEDs), significant penetration of middle and lower cost vehicles is anticipated. Unexpectedly and by far the primary value of LEDs on the outside of a vehicle is their styling benefits. Accordingly, car manufacturers are able to differentiate their vehicles with unique appearances, including thin strips of light, extreme three-dimensional geometries (without sacrificing the highly valued depth in the chassis), clear lenses, and a variety other effects. Other benefits of LED lighting are improved safety, durability, low (thin) profiles, and cost. Safety is improved by the faster turn-on times relative to incandescent bulbs, resulting in an increased stopping distance of 5 m at highway speeds. The high reliability of LEDs result in lifetime lighting

-

1 VISIBLELIGHT-EMITTING DIODES

39

solutions for a vehicle, thus reducing warranty costs. Furthermore, manufacturing costs are improved since the need to provide service access to the lighting function is eliminated. The low profile offered by LEDs results in increased trunk space as well as reduced costs in metal forming. An additional cost savings occurs since smaller components and wires can be utilized as a result of the reduced power consumption (improved efficiency) of the LEDs relative to filtered incandescents.

2. LIGHTING AND ILLUMINATION The revolution in luminous performance (in lumens per watt) of LEDs in the 1990s has resulted in a distinct threat to many conventional lighting sources. An evidence of the viability of this threat is the establishment of joint ventures teaming various lighting and LED companies. In the period from 1997 to 1999, all three of the world’s largest lighting companies (Philips, Osram, General Electric) formed joint ventures with different LED manufacturers. The first inroads are being made in colored lighting applications where LEDs have the strongest advantage relative to conventional sources. In addition, white light LED solutions are beginning to compete with low-power incandescent and halogen solutions. Currently, the majority of these solid-state lighting applications are utilizing white light generated by conversion of the blue emission of InGaN LEDs to white light via a phosphor. In the future, the mixing of individual LED colors (e.g., red, green, blue) to obtain white light may offer significant advantages in terms of performance and cost. The value of LEDs in the lighting and illumination industry is similar to many of the applications previously discussed, although the merits are surely to be weighted differently. Lower cost of ownership will occur through energy savings and reduced maintenance. Furthermore, LEDs offer improved reliability and ruggedness relative to most conventional light sources. Safety is improved for LEDs since they typically operate at much lower temperatures than other sources (especially halogens). Another advantage is the small size and low profile of LED sources. These characteristics improve styling flexibility. Light-emitting diodes also can be operated over a wide dynamic range (for dimming) and are capable of essentially instant turn-on times. White LED sources derived from a mixed color approach offer the potential ability to tune the color temperature and color rendering of the light as well as the ability to “tune” across the full spectrum of colors. Surely, many more benefits will be discovered as the use of LEDs in lighting and illumination applications becomes more prolific.

40

M. G . CRAFCIRD et d.

Future Trends

VI.

1. FUTURE LED PERFORMANCEAND COST

The theoretical limit for LED performance is the CIE curve shown in Fig. 13. If every electron injected into an LED is converted to a photon in the active region (100% internal quantum efficiency), and if every photon is successfully extracted from the semiconductor chip (100% extraction efficiency), then one would have the 100% external quantum efficiency. If one has low-resistance ohmic contacts and low series resistance in the device structure such that the applied voltage is no higher than the quantum (energy gap) voltage, then one would have the perfect emitter with 100% wallplug efficiency (watts output per watts input), and the lumen per watt performance shown by the CIE curve. What can we realistically expect to happen as the technology continues to evolve? The internal quantum efficiency is impossible to measure directly and can only be inferred based on external quantum efficiency measurements combined with modeling calculations about extraction efficiency. Based on measurements of this type the internal quantum efficiency of the best red AlInGaP devices is estimated to be in excess of 90% and the extraction efficiency is about 30% resulting in an external quantum efficiency of around 30%. While some incremental improvement in internal quantum efficiency is probable, any big gain in performance will have to come from the extraction efficiency. Extraction efficiency improvement can

-1 -t ;".\ (CIE)

High Pressure Sodium (1kW1 Fluorescent 140W) MercuryVapor(1kW)

Alln

~

I/

AlGaAs

Red-Filtered Tungsten (60W)

inGaN

1

I'

400

3

Haloaen 130W1 Tungsten (6OW)

\ 500

I

800 Peak Wavelength (nm) 600

700

FIG. 13. Luminous performance of InCaN and AlInGaP and AlGaAs LEDs compares to the theoretical limit for performance (the CIE curve) and to a variety of conventional light sources.

1 VISIBLE LIGHT-EMITTING DIODES

41

be obtained by reducing absorption losses in the chip. Optimization of the epitaxial layer structure, ohmic contacts, surface texture, chip shape, and encapsulating medium all can contribute to improved extraction efficiency. Amber AlInGaP LEDs may have a larger opportunity for efficiency improvement than red AlInGaP since the internal quantum efficiency is presently much lower (- 30 vs > 90%), yielding external quantum efficiencies of around 10% for the best devices. The direct-indirect transition clearly is the dominant factor in limiting the performance of amber AlInGaP, but with improved carrier injection and confinement some improvement is likely in addition to the extraction effects discussed earlier. Overall external quantum efficiencies of around 50% in the red and 25% in the amber seem plausible in the future. Since in these devices the applied voltage is generally not substantially higher than the quantum voltage, the wallplug efficiency could be expected to be comparable to the external quantum efficiency. The improvement in the nitride emitters is harder to predict, although there appears to be more room for improvement than in the AlInGaP system. One of the difficulties in estimating improvement is that the recombination mechanism in the GaInN emitters apparently occurs at small In rich regions in the active region. This results in emission that is substantially lower than the bandgap in energy and which is most efficient at very low current densities. Since this is not the more familiar near-bandedge radiation, it is hard to predict whether it can be improved to close to 100% and whether high efficiency can be maintained at high current densities. The external quantum efficiency of the best InGaN devices is the range of 15% for blue, 10% for green, and 3% for amber. The efficiency decreases with increasing wavelengths, apparently due to the difficulty of adding increasing quantities of indium, but this may be an engineering issue and not a fundamental problem. The internal quantum efficiency of these devices is hard to estimate since the extraction efficiency is not easily modeled. The epitaxial structure itself has a very low absorption, but a semitransparent ohmic contact is generally used that absorbs light, but is required to improve current spreading. If the extraction efficiency is about 50%, which seems plausible, this would imply that the internal quantum efficiencies range from 10 to 30%. Thus it quite plausible that 3-5 x further improvement in external efficiency is possible, particularly in the green and amber spectral regions where the present performance is lower. The forward voltage in the nitride devices is typically significantly higher than the quantum voltage due to resistance in the contacts and epitaxial layers. Consequently the wallplug efficiency is generally 20 - 40% lower than the external quantum efficiency. Clearly if the improvements discussed above can be realized in AlInGaP and nitride devices, LEDs could become competitive with such high-

-

42

M. G. CRAFORD et al.

performance conventional light sources as fluorescents on the basis of luminous performance. Cost may be an even more challenging issue than performance when comparing LEDs to conventional lighting sources. LEDs have historically been low-power devices with power dissipation of about 0.1 W and costs of about $0.5 for high-performance devices. Assuming that the LED has a performance of 10 lm/W this results in a cost of 0.5 $/lm. This can be compared to conventional 100-W light bulbs which cost about $0.5 for 1500 lm, more than three orders of magnitude less than LEDs on a dollars per lumen basis. However, LEDs are evolving into higher power devices with power dissipation in the range of several watts. As this happens and the materials technology improves the cost can be expected to drop to below 0.1 $/lm. There may still be a premium for LEDs compared to highpower conventional sources, but in lower power applications the cost differential will be smaller or nonexistent and the advantages of LEDs, such as long lifetime, more flexible design flexibility, and energy savings, should enable LEDs to penetrate in an ever-increasing part of the lighting market. In higher power applications LEDs may continue to have a difficult time gaining market penetration since the LED cost increases nearly linearly with power dissipation because the device area of the semiconductor chip largely determines the cost. For conventional sources this is not true. It does not cost l o x more to make a 200-W light bulb than a 20-W bulb. In fact, it probably costs very little more. This is a big and quite fundamental advantage for conventional lighting technology in high-power applications. The opportunity for LEDs to compete in this market will likely be determined by the ability of LEDs to provide energy savings to offset higher installation cost.

-

2. WHITELIGHT

There are two general approaches to using LEDs to generate white light. One approach is to mix red, green, and blue LED sources to create white. The other approach is to use an LED source to excite a phosphor to create white. The color mixing approach has the advantage that no energy is wasted in exciting the phosphor. It has the disadvantage that it may be difficult and expensive to optically mix the three colors to get white, and it may be hard to maintain the stability of the color over time or if the power level is changed. In some applications, such as mood lighting, the ability to tune the color by changing the relative intensities of red, green, and blue may be an advantage. The alternative approach of creating white light by using an ultraviolet or blue LED to excite a phosphor has the advantage of simplicity since there

1 VISIBLELIGHT-EMITTING DIODES

43

is a single LED source. It also has the advantage of color control and stability if a ultraviolet source is used because the color is completely controlled by the phosphor. The UV-phosphor approach has the disadvantage that a substantial fraction of the light, probably half or even more, is lost. Some light is lost in the conversion process although the phosphors are quite good, and efficiencies in the 90% range are not uncommon. A larger loss is the quantum deficit loss in down-converting from UV to blue, green, and especially red. This loss can be several tens of a percent, depending on the pump energy. Finally, a substantial fraction of light is lost due to backscattering from the phosphor grains and reabsorption in the LED chip or package. The combination of these effects results in a loss of half of the light or more. If a UV pump source can be made that is 2 x or more higher in efficiency than visible sources, then the net white luminous efficiency could be comparable to the red, green, blue approach and LED UVphosphor lighting could have a broad impact on the lighting market. If the UV source efficiency is comparable to or lower than the visible source efficiency, which is the situation at this time, then the UV-phosphor approach may be limited to a smaller market where simplicity and color control are the key customer requirements and efficiency is less important. An intermediate approach is to utilize a blue LED to pump a broadband yellow phosphor for indicator applications requiring white light. This approach is more efficient since part of the blue light is allowed to escape directly and not lost in the conversion processes. The remainder of the blue light excites the yellow phosphor and the yellow and blue are mixed to create a simple and relatively inexpensive white lamp. One problem with this source is that color control is difficult since the fraction of blue light that escapes must be controlled by precisely controlling the thickness of the phosphor layer. It is reasonable to expect that all of the approaches for white light discussed here will coexist in the market place. For applications requiring a simple, low-power source, the phosphor approach may dominate. For higher power applications where energy efficiency is more of an issue and color mixing complexity can be afforded, the mixing approach may dominate. A key variable which is yet to be determined is the relative efficiency of LED sources of different wavelengths, particularly UV sources, which have not been studied as extensively as visible LEDs. 3. ALTERNATIVE TECHNOLOGIES

The AlInGaP and AlInGaN technologies that presently dominate highperformance LEDs seem likely to remain as the leading inorganic LED

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M. G. CRAFORD et al.

technologies for the foreseeable future. As discussed earlier, it is likely that further improvement can be expected out of both technologies. There does not appear to be a viable alternative inorganic semiconductor material system on the horizon. Group II-VI materials have been used to make high-performance LEDs and lasers in the short-wavelength region (Eason et al., 1995), but reliability issues have prevented their introduction into the commercial market. Organic LEDs have made remarkable progress in recent years and devices with luminous performance comparable to inorganic LEDs in many spectral regions have been achieved. Operating lifetime is a concern, but has improved enough such that it appears to be satisfactory for some displays applications. However, the series resistance that exists in organic LEDs appears to limit their use to relatively low current applications such as displays rather than high-power signaling, exterior automotive, and illumination markets. Some small-area displays have high pixel densities that are not economical for semiconductor LEDs, which must be fabricated using one chip per pixel, but are well suited to organic LED fabrication, which is a batch process involving film depositions on large sheets of glass. The primary competition for organic LEDs in this market is likely to be LCD displays backlit with LEDs. As LEDs become brighter, backlighting becomes increasingly energy efficient and by alternatively turning on red, green, and blue LEDs, relatively inexpensive full-color LCDs with LED backlighting can be fabricated. The major area of competition in the high-performance LED technology area may be between AlInGaP and InGaN technologies in the amber and possibly the red spectral regions. Amber InGaN LEDs have been reported (Mukai et al., 1998) and although they are lower in performance than amber AlInGaP LEDs, the To is higher which makes them suitable for hightemperature operation. Also, since the nitride technology is somewhat less mature than the AlInGaP technology, more improvement may occur in the nitride technology, particularly in the amber area. Of course, there is also a cost issue, and at the present time AlInGaP is substantially less expensive than InGaN.

4. MARKETGROWTH The visible LED market has been growing steadily for three decades and presently is in the range of $2 billion per year. This corresponds to over 30 billion LED chips. The market growth for the mature GaAsP and GaP technologies has been slowing down in recent years with growth rates of less

1 VISIBLE LIGHT-EMITTING DIODES

45

than 10% per year. In 1997 and 1998 the market may actually have decreased for the mature technologies, although they continue to dominate the market in terms of chip volume. The market for the new high-brightness technologies is growing rapidly, with growth rates of greater than 20% for markets such as exterior automotive lighting and large area signs. This growth can be expected to continue into the near future, since the LED market share in those areas is still relatively small and the advantages of LEDs are only recently becoming appreciated. In addition, new markets are emerging that are enabled by LEDs such as markers in the center of roadways. A major issue for LEDs is whether LED technology will be able to substantially penetrate the general illumination market. This will depend on the evolution of performance and cost reductions, as discussed earlier.

ACKNOWLEDGMENTS The authors are grateful for the full support of the Hewlett-Packard Company, Optoelectronics Division. Assistance from D. Hamby, E. Anaya, and E. Wang in the preparation of this manuscript is also greatly appreciated. REFERENCES Alferov, Zh. I., Andreev, V. M., Garbuzov, D. Z., and Rumyantsev, V. D. (1975). Sou. Phys. Sernicond. (Engl. Transl.) 9, 305 Amano, H., Sawaki, N., Akasaki, I., and Toyoda, Y. (1986). Appl. Phys. Lett. 48, 353. Amano, H., Kito, M., Hiramatsu, K., Sawaki, N., and Akasaki, I. (1989). Jpn. J . Appl. Phys. 28, L2112. Arai, T., Hidaka, J., Tokunaga, H., and Matsumoto, K. (1997). J. Cryst. Growth 170, 88. AuCoin, T. R., Ross, R. L., Wade, M. J., and Savage, R. 0. (1979). Solid State Technol. 22, 59. Bergh, A. A,, and Dean, P. J. (1976). Light-emitting Diodes. Clarendon Press, Oxford. Bergunde, T., Dauelsberg, M., Kadinski, L., Makarov, Yu. N., Weyers, M., Schmitz, D., Strauch, G., and Juergensen, H. (1997). J. Cryst. Growth 170, 66. Beuchet, G. (1985). Semiconductors and Semimetals, vol. 22A, p. 261. Academic Press, Orlando. Casey, H. C., Jr., and Panish, M. B. (1978). Heterostructure Lasers. (Academic Press, New York). Craford, M. G. (Sept. 1992). IEEE Circ. Devices 24. Craford, M. G., and Groves, W. 0. (1973). Proc. IEEE, 61, 862. Craford, M. G., and Steranka, F. M. (1994). Encycl. Appl. Phys. 8, 485. Detchprohm, T., Amano, H., Hiramatsu, K., and Akasaki, I. (1993). J. Cryst. Growth 128, 384. DiLorenzo, J.V. (1972). J . Cryst. Growth 17, 189. Eason, D. B., Yu., Z., Huges, W. C., Roland, W. H., Boney, C., Cook, J. W., Jr., Schetzina, J. F., Cantwell, G., and Harsch, W. C., (1995). Appl. Phys. Lett. 66, 115. Fletcher, R. M., Kuo, C. P., Osentowski, T. D., Huang, K. H., Craford, M. G., and Robbins, V. M. (1991). J . Electron. Mater. 20, 1125.

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Gardner, N. F., Chui, H. C., Chen, E. I., Krames, M. R., Huang, J.-W., Kish, F. A,, Stockman, S. A., Kocot, C. P., Tan, T. S., and Moll, N. (1999). Appl. Phys. Lett. (to he published). Groves, W. O., Herzog, A. H., and Craford, M. G. (1971). Appl. Phys. Lett. 19, 184. Haitz, R. H., Craford, M. G., and Weissman, R. H. (1995). In Handbook of Optics, ed. M. Bass, Vol. I, p. 121. (McGraw Hill, Inc.). Hall, R. N. (1952). Phys. Rev. 87, 387. Hardtdegen, H., and Giannoules, P. (1999). 121-Vh Rev. 11, 34. Hodapp, M. (1997). In Semiconductors and Sernirnetals, vol. 48, p. 227, ed. G. B. Stringfellow and M. G. Craford. (Academic Press, San Diego). Hofler, G. E., Carter-Coman, C., Krames, M. R., Gardner, N. F., Kish, F. A,, Tan, T. S., Loh, B., Posselt, J., Collins, D., and Sasser, G. (1998h). Electron. Lett. 34, 1781. Hollfelder, M., Hon, S., Setzer, B., Schimpf, K., Horstmann, M., Schapers, T., Schmitz, D., Hardtdegen, H., Lnth, H. (1997). J Cryst. Growth, 170, 103. Holonyak, Jr., N., (1995). Era Kappa N u Bridge 91, 4-10. Holonyak, N., Jr., and Bevacqua, S. F. (1962). Appl. Phys. Lett. 1, 82. Jensen, K. F. (1994). In Handbook of Crystal Growth: Growth Mechanisms and Dynamics, ed. Hurle, D. T. J., vol. 3B, p. 541. (North Holland Elsevier Science). Kish, F. (1995). Encycl. Chem. Technol. 15, 217. Kish, F. A., Steranka, F. M., DeFevere, D. C., Vanderwater, D. A., Park, K. G., Kuo, C. P., Osentowski, T. D., Peanasky, M. J., Yu, J. G., Fletcher, R. M., Steigerwald, D. A,, Craford, M. G., and Robhins, V. M. (1994). Appl. Phys. Lett. 64, 2839. Knight, J. R., Effer, D. and Evans, P. R. (1965). Solid-State Electron. 8, 178. Kressel, H., and Butler, J. K. (1977). Semiconductor Lasers and Heterojunction LEDs. (Academic Press, New York). Kuech, T. F. and Veuhoff, E. (1984). J . Cryst. Growth 68, 148. Kuo, C. P., Fletcher, R. M., Osentowski, T. D., Lardizahal. M. C., and Craford, M. G. (1990). Appl. Phys. Lett. 57, 2937. Lancefield, D. (1996). Properties of Gallium Arsenide, 3rd ed., eds. Brozel, M. R., and Stillman, G. E., p. 41 (INSPEC, IEE, London). Lee, H., Yuri, M., Ueda, T., Harris, J. S., Jr., Sin, K. (1997). J . Elec. Marl. 26, 898. Lester, S. D., Ponce, F. A,, Craford, M. G., and Steigerwald, D. A. (1995). Appl. Phys. Lett. 66, 1249 Li, Y., and Giling, L. J. (1995). J . Cryst. Growth 156, 177. Lin, J.-F., Wu, M.-C., Jou, M.-J., Chang, C.-M., and Lee, B.-J. (1994). Electron. Lett. 30, 1973. Logan, R. A., White, H. G., and Wiegmann, W. (1968). Appl. Phys. Lett. 13, 139. Manasevit, H. M. (1968). Appl. Phys. Lett. 12, 156. Mason, N. (1987). The Chemistry of the Semiconductor Industry, Chapter 7 (Chapman and Hall, New York). McKee, M. A., McGivney, T., Walker, D., Capuder, K., Norris, P. E., Stall, R. A,, and Rose, B. C. (1992). J . Electron. Mat. 21, 289. Molnar, R. J., Gotz, W., Romano, L. T., Johnson, N. M. (1997). J . Cryst. Growth 178, 147. Moon, R. L. (1997). J . Cryst. Growth 170, 1-10, Mukai, T., Marimatsu, H., and Nakamura, S . (1998). Jpn. J . Appl. Phys. 37, L481. Mullin, J. B., Heritage, R. J., Holliday, C. H., and Straughan, B. W. (1968). J . Cryst. Growth, 34, 281. Nakamura, S., Mukai, T., and Senoh, M. (1994a). Appl. Phys. Lett. 64, 1687. Nakamura, S., Mukai, T., and Senoh, M. (1994h). J . Appl. Phys. 76, 8189. Nakamura, S., Senoh, M., Iwasa, N., Nagahama, S., Yamada, T., and Mukai, T. (1995). J p n . J. Appl. Phys. 34, L1332. Nelson, H. (1963). RCA Rev. 24, 603.

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Nishizawa, J., Shinozaki, S., and Ishida, K. (1973). J . Appl. Phys. 44, 1638. Nishizawa, J., and Suto, K. (1977). J . Appl. Phys. 48, 3484. Nygren, S. F. (1973). J . Cryst. Growth, 19, 21. Schmitz, D., Lengeling, G., Strauch, G., Hergeth, J., and Jurgensen, H. (1992). J . Cryst. Growth 124, 278. Shocklet, W., and Read Jr., W. T. (1952). Phys. Rev. 87, 835. Stringfellow, G. B. (1979) In Crystal Growth: A Tutorial Approach, ed. Brandsley, W., Hurle, D. T. J., and Mullin, I. B., p. 217 (North-Holland Pub., Amsterdam). Stringfellow, G. B. (1989). Organometallic Vapor Phase Epitaxy: Theory and Practice. Academic Press, San Diego. Stringfellow, G. B., and Craford, M. G. (1997). High Brightness Light Emitting Diodes. Semiconductors and Semimetals, vol. 48 (Academic Press, San Diego). Sugawara, H., Ishikawa, M., and Hatakoshi, G. (1991). Appl. Phys. Lett. 58, 1010. Suzuki, T., Akai, S., Kohe, K., Nishida, Y., Fujita, K., and Kito, N. (1978). Sumitomo Elect. Tech. Rev. 18, 105. Swiggard, E. M., Lee, S. H., and Von Batchelder, F. W. (1977). Inst. Phys. Conf: Ser. 33b, 23. Szweda, R. (1996). III-R Rev. 9, 28. Thompson, A. G., Stall, R. A., Kroll, W., Armour, E., Beckman, C., Zawadski, P., Aina, L., and Siepel, K. (1997). J . Cryst. Growth 170, 92. Weiner, M. E., Lassota, D. T., and Schwartz, B. (1971). J . Electrochem. SOC.118, 301. Weissberg, L. R., Rosi, F. D., and Herkart, P. G. (1960). In Properties of Elemetal and Compound Semiconductors, vol. 5, p. 25, ed. Gatos, H. C. (Interscience, New York). Woodall, J. M., Rupprecht, H., and Reuter, W. (1969). J . Electrochem. SOC. 116, 899.

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SEMICONDUCTORS AND SEMIMETALS. VOL. 64

CHAPTER 2

High-Efficiency AlGaInP Light-Emitting Diodes H . C h i . N . E Gardner. P . N . Grillot. J . W. Huang. M . R. Krames. and S . A . Maranowski OPTOELECTRONICSDIVISION HEWLETT-PACKARD COMPANY SANJOSE. CALIFORNIA

I . INTRODUCTION . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11. ALGAINPMATERIAL PROPERTIES . . . . . . . . . . . . . . . . . . . . 1. Band Structure and Lattice Constant . . . . . . . . . . . . . . . . . 2. Optical Properties . . . . . . . . . . . . . . . . . . . . . . . . . 3. Mobility and Carrier Trunsporr . . . . . . . . . . . . . . . . . . . 111. ALGAINPLED DEVICE DESIGN . . . . . . . . . . . . . . . . . . . . 1. Introduction and Overview of Device Design . . . . . . . . . . . . . . 2 . Double Heterostructure Devices . . . . . . . . . . . . . . . . . . . 3. Active Layer Design: Multiquantum Wells . . . . . . . . . . . . . . . 4. Current Spreading . . . . . . . . . . . . . . . . . . . . . . . . . 5. Light Extraction . . . . . . . . . . . . . . . . . . . . . . . . . . IV . EPITAXIAL GROWTH . . . . . . . . . . . . . . . . . . . . . . . . . 1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2. Growth of AlGaInP Materials . . . . . . . . . . . . . . . . . . . . 3. Growth of Device Structures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4 . Manufacturing Issues . . . . . V. DEVICE FABRICATION AND PACKAGING . . . . . . . . . . . . . . . . . 1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2. Die Definition and Contacts . . . . . . . . . . . . . . . . . . . . . 3. Wafer Bonding . . . . . . . . . . . . . . . . . . . . . . . . . . 4.Dicing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5. Packaging Issues . . . . . . . . . . . . . . . . . . . . . . . . . . VI . ALGAINPDEVICE PERFORMANCE . . . . . . . . . . . . . . . . . . . . 1. LED Eficiency and Current- Voltage Characteristics . . . . . . . . . . 2 . Eficiency vs Wavelength . . . . . . . . . . . . . . . . . . . . . . 3. Luminous Eficiency . . . . . . . . . . . . . . . . . . . . . . . . 4. Thermal Management . . . . . . . . . . . . . . . . . . . . . . . 5. High-Power AlGaInP LEDs . . . . . . . . . . . . . . . . . . . . . 6. AIGaInP Degradation . . . . . . . . . . . . . . . . . . . . . . . VII . CONCLUSIONS . . . . . . . . . . . . . . . . . . . . . . . . . . . . REFERENCES. . . . . . . . . . . . . . . . . . . . . . . . . . . . .

so 51 51

54 51

51 57 58 63 64

66 12 12

13 88 92 99 99 99

100 100 101 102 102 108 109 112 116 118 123 123

49 Copyright 02000 by Academic Press All nghts of reproduction in any form reserved. ISBN 0-12-752173-9 ISSN 0080-8784/00$30.00

50

H. CHUIet al.

I. Introduction

The AlGaInP quaternary alloy system has been widely used for visible wavelength optical devices such as lasers and light-emitting diodes (LEDs). Commercial AlGaInP LEDs were independently developed and introduced to the marketplace by the Hewlett-Packard Company (Kuo et aE., 1990) and Toshiba Corporation (Sugawara et al., 1991) in the early 1990s, nearly 30 years after the first red GaAsP devices were introduced in small quantities in the early 1960s by General Electric (Holonyak and Bevaqua, 1962).The original vapor phase epitaxially (VPE) grown GaAsP LEDs were only -0.3 lm/W (0.4% radiant efficiency) in brightness, but their wavelength range and efficiency increased in the early 1970s with the addition of nitrogen as an isoelectronic impurity (Logan et al., 1968; Groves et al., 1971). GaAsP and GaP doped wlth nitrogen made devices in the red, orange, and yellow-green a possibility. However, the low efficiency of the devices (-1 lm/W, 1.4% radiant efficiency at 650nm, 0.2% at 590 nm) restricted their uses to applications such as indicators and indoor displays where high brightness is not required. In the mid-l980s, the introduction of AlGaAs direct bandgap red LEDs, first as absorbing substrate (AS) devices and later as transparent substrate (TS) devices, grown by liquid phase epitaxy (LPE) resulted in improved efficiencies of 10 lm/W (- 14% radiant efficiency) (Alferov et al., 1975; Nishizawa and Suto, 1977). These were the first LED devices to take advantage of the efficient doubleheterostructure (DH) design, and made outdoor applications such as red traffic signal lights and automobile brake lighting a possibility. Finally in the 1990s, researchers combined the DH design with kinetically controlled crystal growth processes such as metalorganic vapor phase epitaxy (MOVPE), molecular beam epitaxy, and gas source molecular beam epitaxy (Chen, C. H. et al., 1997) to produce high-efficiency (Al,Ga, -x)o.sIno.5P LED devices. When lattice matched to GaAs, these LEDs cover a wide range of the visible spectrum from -1.9 eV (653 nm, red) for x = 0 to 2.2 eV (555 nm, green) for x -0.53, although device efficiency decreases considerably in the green. The AS AlGaInP technology enables light output in the range of 10 lm/W (- 14% radiant efficiency at 650 nm, 2% at 590 nm) for the red through the yellow wavelengths. For the red-emitter application, the AlGaInP material is a very attractive alternative to the AlGaAs alloy system, which has a higher surface recombination velocity and is more prone to hydrolyzation (Hobson, 1995). The performance of AlGaInP LEDs increased by a factor of 2 to a typical value of >20 lm/W (-28% radiant efficiency at 650 nm, -4% at 590 nm) with the introduction of TS AlGaInP in 1994 (Kish et al., 1994). This technology has also enabled

-

-

N

2 HIGH-EFFICIENCY AlGaInP LIGHT-EMITTING DIODES

51

the highest LED efficiencies (50 lm/W, 14% radiant efficiency) yet observed in the orange (1-605 nm) wavelength range (Kish et al., 1996) and the luminous performance of AlGaInP LEDs is now approaching or exceeds that of traditional lighting sources such as fluorescent and incandescent light bulbs (Craford, 1997). This chapter focuses on the key issues regarding the development and performance of AlGaInP LEDs. We begin with a discussion of the basic AlGaInP material properties and issues related to the design of an AlGaInP DH LED. A survey of the methods used to achieve good current spreading and light extraction in AlGaInP LEDs follows. Since the realization of high-brightness AlGaInP LEDs depends heavily on a good epitaxial growth process, a section describing the important features of the MOVPE growth for AlGaInP LED devices is included. Device fabrication and packaging are also described, followed by a description of the state-of-the-art device performance including color, brightness vs wavelength, reliability, and highpower devices. We conclude with a discussion of the potential and future direction of AlCaInP LEDs.

11. AlGaInP Material Properties AlGaInP is an important optoelectronic material for light-emitting diode and visible laser diode applications. In this section, we consider the material properties of AlGaInP alloys, beginning with a discussion of the band structure and lattice constant of AlGaInP. We next discuss the photoluminescence of AlGaInP, and we conclude this section with a discussion of mobility and carrier transport properties.

1. BANDSTRUCTURE AND LATTICE CONSTANT

Figure 1 illustrates the bandgap and lattice constant values attainable with arbitrary composition AlGaInP alloys. Note from Fig. 1 and Table I that a range of bandgap values from 2.45 eV (Alp) to 1.35 eV (InP) is achievable, and that the lattice constant may be varied from 5.4511 to 5.8686 A at 300 K. More realistically, production AlGaInP devices are grown on GaAs substrates, and the growth of high-quality AlGaInP alloys thus requires lattice matching to GaAs at the epitaxial growth temperature (Bour and Shealy, 1988). Since AlP and GaP have virtually identical lattice constants, A1 and Ga can be traded off for one another without affecting the alloy lattice constant. This lattice-matching requirement hence constrains

52

H. CHUIet al. 2.50

5- 2.25 *

0)

2.00

e

a,

2 1.75 m

U

m

1.50

1.25 5.30 5.40 5.50 5.60 5.70 5.80 5.90 6.00 Lattice constant (A) FIG. 1. Bandgap vs lattice constant of the AlGaInP material system with binary endpoints Alp, Gap, and InP. AlAs and GaAs are also indicated. Solid lines represent direct bandgap alloys, while dotted lines represent indirect bandgap alloys. The dashed vertical line represents the lattice constant of GaAs.

only the In composition of ( A l x G a l ~ x ) y I n l ~toy Pbe fixed at y-0.50, with the exact lattice-matching composition depending on the epitaxial growth temperature. Unless stated otherwise, an alloy composition of (Al,Ga, -,)o,sIno,sP is assumed throughout the remainder of this chapter. The lattice-matching requirement just discussed constrains the commercially viable AlGaInP alloy system to the vertical line in Fig. 1, hence reducing the available bandgap values to the range from 1.91 eV (Ga,,,In,,,P) to 2.28 eV (Alo.51no,5P).A more complete understanding of TABLE I LATTICE CONSTANT, THERMAL EXPANSION COEFFICIENT, AND BANDCAP OF SEVERAL RELEVANT SEMICONDUCTOR MATERIALS AT 300 K Equilibrium lattice constant, Compound

a,&

GaAs AlAs GaP AlP InP

5.6533 5.661 1 5.4512 5.45 11 5.8686

Lattice mismatch to GaAs at 300 K (YO)

Thermal expansion coefficient, ~ c ( l O -K-'Y ~

Bandgap, E,(eV)b

0 0.14 - 3.1 - 3.7 3.1

6.86 5.2 5.91 4.50 4.15

1.42 2.16 2.26 2.45 1.35

"Adachi (1985) and Chen, C. H. et al. (1997). bStreetman (1990) and Adachi (1982).

2 HIGH-EFFICIENCY AlGaInP LIGHT-EMITTING DIODES

53

the band structure of (AI,Ga, -,)o~51no,5P can be acquired by considering the compositional dependence of the r- and X-conduction band minima in Eqs. (1) and (2). Alloy ordering, as discussed in Section IV, can significantly decrease the bandgap of AlGaInP alloys, and this has led to some confusion regarding the band structure of AIGaInP. Equations (1) and (2) are widely accepted as correctly describing the r- and X-band minima of lattice matched (Al,Ga, -,)o,sIno,5P, however, assuming a completely random, disordered alloy. 1.91 + 0 . 6 1 ~

E,(x)

=

E,(x)

= 2.19

+ 0.085~

(eV)

(1)

(eV)

(2)

Note that Eqs. (1) and (2) predict a crossover from the direct bandgap r minimum to the indirect bandgap X-band minima at a composition of x-0.53. (Cao et at., 1990; Bour and Shealy 1988; Asahi et at., 1982; Kish and Fletcher, 1997). The use of MOVPE and the ability to trade off A1 for Ga in the AlGaInP LED structure without affecting the lattice constant provide a major advantage over more conventional optoelectronic devices in that the composition, and hence the bandgap of the alloy layers can be tuned, either slowly or abruptly, over the range from 1.91 (x = 0) to 2.28 eV (x = 1). By using wide-bandgap confining layers above and below the active layer, one can confine both electrons and holes to a well-defined and localized volume, and enhance radiative recombination relative to bulk semiconductor devices, as described in Section 111. Careful consideration of the confining layer composition and resulting bandgap offsets is thus an important parameter in optimal device design and Fig. 2, along with Eqs. (3) through (5) show the approximate conduction band and valence band energies as a function of aluminum composition. (Kish and Fletcher, 1997). AE,(x) = 0.369x(eV)

for x Q 0.53

(3)

AEc(x) = 0.285 - 0.157x(eV)

for x > 0.53

(4)

AE,(x) = 0.241x(eV)

for 0 >, x >, 1

(5)

Note that these energies are given relative to the valence band energy of Ga,,,In,,,P while Eqs. (1) and (2) are given relative to the valence band energy of (Al,Ga, -,)o,51no,5P, which varies with A1 composition, x. For small aluminum composition, x-0 to 0.3, the r- and X-band minima are well separated in energy, and the 300-K photoluminescence intensity from such alloys shows little compositional dependence as discussed in more detail later (Cao et al., 1990).

54

H. CHUIet al. 2.2

,

I

E,(x=1)=2.28 eV

-0.4 0

0.2

0.4

0.6

0.8

1

Al mole fraction of (AI,Ga,,)o,51no~5P, x

FIG.2. Energy of conduction and valence band extrema vs A1 mole fraction of AlGaInP lattice matched to GaAs. Note that the maximum conduction band offset occurs for an A1 mole fraction -0.53, where the conduction band minima change from direct bandgap to indirect bandgap.

2. OPTICAL PROPERTIES The light-emitting properties of AlGaInP layers have been characterized by several groups using temperature-dependent photoluminescence (PL) studies (Cao et ul., 1990; Naritsuka et al., 1991). In these experiments AlGaInP films are photoexcited using a 488-nm Ar' ion laser beam of photon energy higher than the bandgap of the AlGaInP layer of interest, typically referred to as the active layer. Electron-hole pairs are generated in the material and recombine either radiatively or nonradiatively according to the internal quantum efficiency of the active layer. Figure 3a shows the PL intensity of (Al,Ga, -,)o,51n,~,P single layer films on GaAs as a function of x, the A1 composition of the alloy (Cao et al., 1990). Data is shown for both low-temperature ( T = 10 K) and roomtemperature (300 K) photoluminescence. At 10 K, the PL intensity is observed to be bright and relatively constant with increasing A1 composition of the AlGaInP films below the direct-indirect crossover at x -0.53. This suggests that the AlGaInP material system is capable of high internal quantum efficiency throughout the direct bandgap regime when nonradiative recombination associated with the indirect minima and other recombination centers is avoided. The constant intensity from x 0 to 0.5 is expected at a low temperature since carrier population of the indirect

-

2 HIGH-EFFICIENCY AlGaInP LIGHT-EMITTING DIODES

55

Aluminum Mole Fraction, x FIG.3. (a) Photoluminescence intensity as a function of A1 composition for single-layer films of AlGaInP on GaAs at T = 10 and 300 K. (b) Photoluminescence intensity as a function of A1 composition of the active layer for AlGaInP double heterostructures on GaAs at T = 12, 50, 100,200, and 250 K.

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conduction band minima is proportional to exp( - AE/kT), where AE is the separation between the direct and indirect transition energies. Also, the PL intensity is high for all compositions in Fig. 3a because the activation of nonradiative defect centers is suppressed at this low temperature. At room temperature, carrier population of the indirect minima becomes observable and, for x =. -0.3, the PL intensity decreases rapidly as the direct bandgap carrier population is depleted relative to that of the indirect minima. Moreover at room temperature, PL intensity is seen to decrease by almost two orders of magnitude, even at low A1 compositions, due to the activation of nonradiative defect centers within the material. Because of the low level of PL excitation (50 to 200 W/cm2, corresponding to a forward injection current density of roughly 1 Ajcm’) there is no saturation of trap centers. Thus, nonradiative lifetimes are very short and the internal quantum efficiency is very sensitive to the presence of these traps. However, the thermalized population of the indirect minima with increasing A1 composition is a fundamental mechanism that limits the light output efficiency of AlGaInP-based LEDs. Note that near the directto-indirect crossover (x -0.53) the PL intensity approaches zero. Figure 3b shows data from another PL study on AlGaInP films (Naritsuka et al., 1991). In this case, the AlGaInP active layer is sandwiched between two higher bandgap layers, which serve to confine carriers to the active layer and enhance internal quantum efficiency. Also, the upper confining layer serves to protect carriers from nonradiative recombination at the exposed top surfxe of the sample. As in Fig. 3a, PL intensity is found to be constant with increasing A1 composition at T = 12K. At higher temperatures, PL intensity begins to drop significantly with a dramatic dependence on A1 composition. This drop with increasing A1 composition is observed immediately with x > 0, unlike the case for Fig. 3a. The difference is attributed to an increased sensitivity to nonradiative traps with increasing temperature and is consistent with the fact that impurity incorporation increases with increasing Al content in these films. in fact, oxygen and silicon are common impurities that incorporate more readily with increasing A1 content of the film during growth. Another possibility for the increased dependence of PL intensity with A1 composition x in Fig. 3b is the lower excitation level used for this study (2 to 70 W/cm’, which corresponds to a forward injection current density of > 1 A/cmZ), for which carriers are even more sensitive to the presence of traps. Figures 3a and 3b serve to indicate the fundamental loss mechanisms present in the AlGaInP material system that limit the performance of devices at shorter wavelengths. Another fundamental mechanism that limits performance in practical

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devices is carrier confinement, which is discussed in Subsection 2 in Section 111. 3. MOBILITY AND CARRIER TRANSPORT

Carrier transport properties have a significant effect on device design and optimization, as will be discussed in Section I11 of this chapter. To achieve optimum LED performance, it is important to consider the differences between electron and hole mobilities in AlGaInP alloys. Hall effect measurements on AlGaInP alloys typically indicate a hole mobility 10 cm2/V s, while electron mobilities tend to be 50 to 500 cm2/V s (Ohba et al., 1986; Suzuki et al., 1994; Maranowski et al., 1997). This higher electron mobility, relative to holes, means that it is easier to inject electrons into p-type AlGaInP than it is to inject holes into n-type AlGaInP (Kish and Fletcher, 1997), and the efficiency of AlGaInP LEDs with a p-type active layer may be slightly higher than AlGaInP LEDs with an n-type active layer. Finally, note that the carrier effective masses of AlGaInP alloys are significantly higher than those of AlGaAs, and this can have important implications on current spreading in AlGaInP LEDs as discussed in Section I11 (Kish and Fletcher, 1997).

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111. AlGaInP LED Device Design 1. INTRODUCTION AND OVERVIEW OF DEVICE DESIGN This section describes the essential device design features of high-performance AlGaInP LEDs. The design and construction of AlGaInP LEDs has evolved to address fundamental issues related to achieving true highbrightness devices, namely current spreading and light extraction. A depiction of the two most common high-performance AlGaInP LEDs is shown in Fig. 4. In both of these designs, current spreading and light extraction are enhanced relative to the basic DH device. In this section, we begin by discussing the fundamentals of DH devices, including carrier injection, carrier confinement and active layer design. We next review device designs used to enhance current spreading, including the use of both a thick p-type window layer and a current-blocking structure, as shown in Fig. 4. Finally, the section concludes with a description of techniques used to improve light extraction from AlGaInP LEDs such as transparent substrate technology, or the use of distributed Bragg reflectors (DBRs), also shown in Fig. 4.

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(a)

Direction of light emission

wafer-bonded interface

/

n-GaP substrate -2oOpm

FIG. 4. (a) Wafer-bonded transparent-substrate AlGaInPiGaP LED structure employing a GaP window layer for improved current spreading and light extraction. (b) Absorbing substrate AlGaInP/GaAs LED structure employing an AlGaAs window layer and a bottom DBR mirror.

2. DOUBLE HETEROSTRUCTURE DEVICES u.

Background

High-quality AlGaInP material is produced by growing an alloy of (Al,Ga, -,)o,51no,5Plattice matched to a GaAs substrate. The dominant epitaxial growth technique for AlGaInP LEDs is MOVPE, as described in

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detail in Section IV. Like most other 111-V material systems (with the notable exception of InGaN materials), the radiative efficiency of AlGaInP is strongly dependent on the density of dislocations and other crystalline defects (Lester et al., 1995). While a number of attempts have been made to grow high-quality GaInP and AlGaInP on mismatched substrates such as GaP (Nuese et al., 1972) or thick compositionally graded GaAsP on GaAs (Macksey et al., 1973; Lin et al., 1993), the efficiency of these devices is roughly an order of magnitude lower than the best commercially available lattice-matched AlGaInP devices. The available colors of (Al,Ga, -,)o,51no,5PLED devices range from red to yellow-green and are limited by the direct-indirect transition at x = 0.53, as described in Section 11. Devices with the highest commercial potential span the red (-640 nm) to the yellow (-588 nm) wavelengths. The yellow-green devices (- 560-570 nm) are considerably less efficient due to the nonradiative recombination of carriers in the indirect X-band minimum. LED efficiency versus wavelength will be described in detail in Section VI. The bandgap of AlGaInP material is reduced (by up to 190 meV) by an effect known as atomic ordering, or the formation of a monolayer scale Gap-InP (or Alp-InP) superlattice on (111) crystal planes (Gomyo et al., 1986). Ordering is more likely to occur under certain MOVPE growth conditions and will be described in detail in Section IV. Because AlGaInP LEDs are particularly useful for short-wavelength applications, disordered material is generally preferred and is assumed for the following discussion. The AlGaInP system is suitable for growth of DH LEDs with the benefits of increased injection efficiency and carrier confinement (Kroemer, 1963). Since the advantages of the DH were well known during the development of the AlGaInP materials, virtually all LEDs in this system employ a DH. The benefits of the heterostructures were first applied to LEDs in the AlGaAs system (Nishizawa et al., 1977, 1983), first using single heterostructures (SH) and later using a DH. High-efficiency AlGaAs LEDs with external quantum efficiencies of 18% at 300 K were achieved by using a DH (Cook et al., 1988), while the earlier SH device efficiencies were more commonly in the range of 4-7%. To realize these high efficiencies, highAl-composition AlGaAs confining layers must efficiently inject carriers into a lower A1 composition active layer and effectively confine the recombination of the injected carriers to the active layer. The first reported AlGaInP LEDs also utilized a DH active region (Kuo et al., 1990; Sugawara et al., 1991). Since lattice matching to GaAs can be maintained for all ratios of A1 to Ga, confining layers of both (Alo,7Gao,3)o,51no,5P and Alo,51no,5Pare commonly used in DH devices. While Alo,51n,,5P confining layers can be used to maximize the bandgap difference, AEg, to the AlGaInP active layer, it is more difficult to achieve a

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high hole concentration in Alo,sIno,5Pcompared to (Alo,7Gao,3)o,sIno,5P. High conductivity in the DH confining layers is important to achieve good current spreading and low electrical resistance. This is especially important for the case of the low-mobility holes. A band diagram of a typical (Al,Ga, -,)o,sIno~,PDH structure is shown in Fig. 5. The structure is under forward bias since the electron and hole quasi-Fermi levels in the confining layers are separated. For AI,,51no,5P confining layers, we expect approximately two-thirds of the bandgap discontinuity between the confining and active layers to occur in the valence band (Kish and Fletcher, 1997). The active layer can be tuned for emission from 650 (x = 0) to 555 nm (x-0.53) and is typically -1 pm thick. The electron confining potential Ae is also shown in Fig. 5. This parameter is quite important in describing the hightemperature operation of AlGaInP LEDs and is discussed in more detail later in this section. Doping densities of the confining layers are commonly in the range of -2 x l o r 7 to - 5 x lO1'cmP3 for the n-type confining layer and -4 x lo1' to - 2 x 10" ~ r n for - ~the p-type confining layer. However, the acceptor concentration of Alo.51no,5Pis usually limited to t 2 x 1OI8 ~ r n - ~ under typical MOVPE growth conditions. Most reports describing AlGaInP LEDs refer to the active layer as unintentionally doped. In practice, the active layer of the device is commonly lightly doped n- or p-type at ,< 10'' cm-3 with the p-n junction located inside the active layer. Thus, the doping scheme may consist of a lightly doped p-type active layer (n, p - , p ) , a lightly doped n-type active layer (n, n-, p), or a junction in the middle of the active layer (n, n-, p-, p ) . b.

Carrier Injection

To illustrate the increase in carrier injection efficiency obtained using a DH structure, consider the injection of holes into a lightly doped n-type

FIG. 5 . Band structure of an AlGaInP LED.

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AlGaInP active layer (n, n - , p). LEDs are usually operated at low current densities (< 110 A/cm2) and low injected carrier densities ( < 7 x 1 O I 6 cmP3). Injection of holes will occur across the interface between the p-type upper confining layer and the n-type active layer. The ratio of injected hole current to injected electron current is given by (Kressel and Butler, 1977):

where J , and J , are the injected electron and hole current densities; De and D, and Le and L, are the minority electron and hole diffusion coefficients and diffusion lengths, respectively; N , and N , are the electron and hole doping densities; and m, and m, are the electron and hole effective masses (on the n-side in the numerator and the p-side in the denominator). This expression shows that the injection of holes is favored from the widebandgap p-type into the n-type active layer by an exponential factor containing AEg, the energy gap difference between the upper confining layer and the active layer. For most AlGaInP LEDs this factor is quite large resulting in essentially single-sided injection. For a green AlGaInP LED operating at 560 nm with AlInP confining layers, the exponential factor is 150, so most of the current will consist of holes injected into the active layer, and not electrons injected from the active layer into the p-type confining layer. The analysis using Eq. (6) assumes an ideal heterojunction without interfacial recombination, and neglects the effects of conduction and valence-band offset spikes.

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c.

Carrier Confinement

In addition to providing efficient injection of minority carriers, the DH structure also serves to confine the injected carriers to the active region. Injected electrons and holes in the active layer may recombine either radiatively or nonradiatively, or may escape into the confining layers where they will either recombine nonradiatively (in the indirect material) or return to the active layer. The DH structure enhances the device’s efficiency by reducing the carrier leakage from the active layer. Reduced leakage decreases the possibility of nonradiative recombination in the confining layers and also increases the density of injected carriers in the active layer. A high injected carrier density is preferred to help saturate nonradiative recombination centers present in the active layer. To realize the benefits of carrier confinement, the active layer thickness should be less than the injected minority carrier diffusion length. To achieve high-brightness AlGaInP

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LEDs, the increased injected carrier densities in the D H must be combined with low densities of nonradiative recombination centers in the active layer and confining layers, as described in Section IV. Although the DH reduces carrier leakage into the confining layers, it does not eliminate carrier leakage. In general, minority carriers in the active layer may escape to the confining layers by both diffusion and drift processes. As the hole mobility in the confining layers tends to be much smaller than the electron mobility, the dominant leakage mechanism is that of electrons into the p-type confining layers. In AlGaInP laser diodes, this lack of electron confinement results in higher threshold currents and a reduced maximum operating temperature (Bour, 1993). The laser diodes typically have higher injected carrier densities (from both the high current densities and thinner active layers) and higher electric fields than LEDs, and are therefore subject to both drift and diffusive leakage. For the lower injected carrier densities in LEDs, the drift leakage component may be neglected as an approximation. The size of the electron diffusive leakage from a lightly doped p-type active layer into a p-type upper confining layer (UCL) is proportional to the minority carrier concentration np at the edge of the depletion region in the UCL (Hagen et al., 1990). This concentration is given by

where N c ( u T 3 / 2 )is the conduction band density of states and A& is the potential barrier to electron flow, as depicted in Fig. 5. The quantity A& can be expressed in more fundamental terms as

where E,, and E,, are the quasi-Fermi levels for electrons and holes in the layers of interest (“al” is the active layer and “cl” is the p-type confining layer). To reduce the electron leakage, the quantity np should be minimized, which occurs when the barrier A&is maximized. This is accomplished by a large AEg, by a large term { ( E f P T ’- (E,,)”} that results when the p-type confining layer is heavily doped and the active layer is lightly doped, and by a small term (E,,)”’ that occurs as the injected electron density in the p-type active layer decreases. To translate these trends into device design choices for a low-leakage AlGaInP LED, a very heavily doped p-type AlInP confining layer, a lightly doped p-type active layer, and a low injected carrier density (a function of both current density and active layer thickness) are preferable. However, as discussed in Section IV, it is difficult to achieve

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high p-type doping in AlInP, which often leads to the choice of (Alo,7Gao,3)o~51no,5P for the p-type confining layer. For most AlGaInP LEDs operated at room temperature with 2 > 590 nm, with typical doping levels and injected carrier densities, and with 1 pm active layer thickness, computer simulations indicate that the electron leakage current is less than 20% of the total injected current density (Kocot et al., 1998). However, for shorter wavelength devices, this leakage does begin to limit device efficiency. Suzuki et al., (1994) reported a 50% relative increase in external quantum efficiency when increasing the confining layer composition of a 566-nm AlGaInP LED from (Alo,7Gao~3)o,51no,5P to AlInP. This is consistent with a reduction in electron leakage current due to the larger AEg obtained for AlInP confining layers. Amber (2-588 nm) AlGaInP LEDs may also begin to suffer in efficiency due to the loss of electron confinement in the limit of higher injected electron densities, low p-type confining layer doping and lowAl-composition confining layers. Several methods have been employed to reduce carrier leakage in AlGaInP heterostructures, one of which is the use of strained-layer confining layers (Chang et al., 1997a). In these structures, the In mole fraction of a 200-A-thick region in the p-type confining layer is chosen to be Alo.651no,35Pto provide tensile strain to shift the band structure and increase the confining layer barrier height for electrons and holes. Another method is to utilize multiquantum well barrier (MQB) structures, which act as “reflectors” for electrons of certain energies much like a DBR acts as a reflector for photons of certain wavelengths (Iga et al., 1986; Morrison et al., 1996). These structures are placed in the p-type confining layer to provide an effectively increased potential barrier to minority carrier electrons, forcing them to remain in the active region, even under higher drive or higher temperature situations, so that they may have a more favorable chance of recombining radiatively. An interesting feature of the MQB structure is that improvements in threshold current and characteristic temperature occur even when the MQB region is “leaky,” or not optimized for electron reflection. Although these techniques have shown improvements relative to “control” LED or laser devices in each case, an increase in the record performance of LED devices has not been reported by using these structures, so the ultimate potential of the improved-confinement techniques for AlGaInP LEDs is difficult to measure.

-

3. ACTIVELAYERDESIGN: MULTIQUANTUM WELLS Although the DH structure has been the standard for commercial AlGaInP LEDs, improvements to D H performance have been attempted by

employing a multiple-quantum-well (MQW) active layer. Sugawara et al. (1994a) fabricated MQW AlGaInP LEDs with 3 to 40 ( A l o . ~ G a o , * ) ~ , ~ I n ~ , ~ P wells separated by ( A l ~ , 5 G a ~ , 5 ) o , ~ Ibarriers. n o , ~ P Because quantum efficiency decreases with increasing A1 composition in an AlGaInP active layer, the quantum size effect (that shortens the wavelength for a given A1 composition) realized in these MQW devices results in a benefit in achieving high-efficiency short-wavelength LEDs. Performance of these LEDs was improved relative to DH control samples at the same wavelength, but was limited to external quantum efficiency qelt of -1.6% at 600 nm. An interesting feature of these results is the tendency toward increasing brightness as the number of wells is increased, which suggests that interfacial recombination is not a significant loss mechanism for these LEDs. 4. CURRENT SPREADING A useful LED requires a practical means of making electrical contact to the p-type and n-type regions. The choices of contact geometry and contact-layer design have significant impact on the performance of LEDs due to their effects on current spreading and light extraction. Unlike the case for edge-emitting devices such as laser diodes, LEDs emit potentially useful light from the active layer in all directions, not just the desired vertical direction. A problem here involves light extraction in that a significant portion of the light generated by carriers injected directly below the top contact is occluded by the metallization at the surface. The high resistivity of AlGaInP layers exacerbates this problem because current cannot spread away from the top contact within thin layers so a majority of the carriers recombine directly below the electrical contact. In addition to reducing light extraction, the resulting nonuniform current density in the active region hinders LED reliability and limits the maximum possible drive current. As discussed in Subsection 6 of Section VI, current density plays a major role in AlGaInP device degradation, and a nonuniform current density reduces the upper limit on the allowable total drive current compared to the case of uniform injection across the active layer. This is a compound problem in that the region of highest current density (thus, greatest photon density) is that directly beneath the top contact, where light extraction through the top surface is poorest. A number of methods have been proposed and have proved to be manufacturable solutions to the problem of current spreading. One method is to provide a thick, conductive, transparent superstrate for the device. This approach is powerful in that it attacks both the current-density uniformity problem and also provides for increased light extraction through the sides

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65

of the superstrate, also referred to as a “window” layer. The degree of improvement in light extraction depends on the geometric aspect ratio of the device, that is, the height of the transparent window relative to the size of the chip. The degree of uniform current spreading depends on the thickness and sheet conductivity of the window as well as the top contact configuration. Kuo et al. (1990) describe a method wherein a latticemismatched (Au/u-3.7%) GaP layer is successfully grown on top of an AlGaInP heterostructure to serve this purpose. The relatively high hole mobility in p-type GaP ( p p 62 cm2/V s) makes this material far superior to p-type AlInP (pP 10 cm2/V s) as a current-spreading layer. Part of this technique’s success lies in the fact that the misfit dislocations caused from the lattice-mismatched growth remain within a few thousand angstroms of the mismatched interface and do not show evidence of propagating downward into the GaAs-lattice-matched layers (Fletcher et al., 1991). Further development of this technique, including hydride vapor-phase epitaxy (HVPE) for faster growth rates, has resulted in GaP superstrates up to -45 pm or more in thickness and enabled record-level achievements in external quantum efficiency (Fletcher et al., 1991). A similar approach to this thick GaP window has been demonstrated using an MOVPE-grown latticematched transparent Al,Ga, _,As (x > 0.7) superstrate (Sugawara et d., 1991, 1992a). The slower growth rate of MOVPE relative to HVPE, places a practical upper limit on the thickness of the AlGaAs window layers (- 7 pm) that are significantly thinner than that available in the HVPE G a P window approach. Hence, the current spreading capability is not as effective as in the VPE GaP window approach. Figure 6 compares theoretical calculations of the current spreading provided by various window technologies (Kish and Fletcher, 1997).

-

3

-

-I

fa) -p (50 pnp (b) f i O . @ O . d ~ (7 Prro Alo,,h# (5Cm l (d) f i ~. , dn~.(1~P prrp

7

0

0

50 75 1m Distance from Chip Center, x (pm)

25

I 125

FIG.6. Current spreading with various p-type window layer technologies.

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Other approaches to the current spreading problem have been implemented using transparent conducting films deposited after epitaxial growth. Films such as indium tin oxide (ITO) with low resistivity and reasonable transparency may be deposited by either evaporation or sputtering to form a combination contacting and current-spreading layer (Lin et al., 1994a). While this technique provides a significant improvement over LEDs without any other means of spreading current, this method is inferior to the transparent window approach discussed earlier in the chapter. Finite transparency within the I T 0 (a-2-5 x lo3 cm-') and the requirement of a highly doped low-bandgap material (e.g., GaAs) for low-IT0 contact resistance results in significant optical absorption at the top surface of devices employing this technique (Chua et al., 1997). Another method demonstrated to increase current spreading employs current-blocking structures, which have been used successfully in AlGaInP laser diode structures to provide lateral current confinement. For LEDs, such structures are implemented directly beneath the top contact to force current around and outside them so that carrier recombination occurs in regions not occluded by the top contact. This method is particularly effective in devices with thin superstrates ( e g , 7-pm AlGaAs) but becomes less important in structures with thick superstrates because of reduced solid angle occlusion by the top contact in these cases. Current-blocking structures may be provided by several means. One method comprises a two-step growth process wherein a counterdoped capping layer is patterned so that a second growth step, which completes the device structure, provides for selected counterdoped regions. These regions are registered to the top contact metallization. During operation the reverse-biased junction provided by the counterdoped region blocks current flow under the top contact and forces recombination in the desired areas, away from the top contact (Sugawara et al., 1992a). A similar method to this one includes providing an AlInP-GaAs heterobarrier instead of a counterdoped region to achieve the same effect. A slightly different approach utilizes dopant diffusion after a single growth step (Sugawara et al., 1992b).In this latter method, the current blocking is achieved by the insertion of a thin n-doped layer in the p-type upper layers which is type-converted during a subsequent anneal via area-selective Zn diffusion. In the areas where the Zn diffusion is inhibited, a reverse-bias junction blocks current flow directly underneath the top contact.

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5. LIGHTEXTRACTION

Light extraction, as quantified through the extraction coefficient, C,, is a fundamental problem that limits the external quantum efficiencies qext

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67

attainable from LEDs. The root of the problem lies in the large difference between the refractive indices of the light-generating semiconductor material and the ambient (typically air) where useful output light is required. Considering a simple planar interface through which light must be transmitted, the extraction of light from a semiconductor to ambient is governed by Snell’s law, which specifies a critical angle beyond which light is totally internally reflected by the high-to-low refractive index transition at the semiconductor-ambient interface. The critical angle thus defines an “escape cone” of light, which may be transmitted through the interface. Assuming isotropic emission from within the semiconductor, the portion of emitted light that is within the solid angle of the escape cone is -n,2/4n;, where n, is the refractive index of the ambient and n, is that of the semiconductor. For extraction from an AlGaInP active region to air, this corresponds to 1/(4 x 3.52) = 2% of the total light emitted from the active region. The remaining 98% of the light is either going in the other direction (toward the substrate) or is totally internally reflected (TIR) at the semiconductorair interface back toward the substrate. In cases where the substrate is absorbing (e.g., GaAs), this 98% of the generated light is irretrievably lost. For thin-window, absorbing substrate devices, the preceding planar model for extraction is a reasonable approximation since emission out the side surfaces of the LED is negligible. For such devices, then, light extraction (and thus external quantum efficiency) is limited to a maximum value of -2%. One thing that makes the situation considerably less bleak, however, is the availability of a range of encapsulation materials that can be cast or molded to form curved lenses around the LED chip to form LED lamps. These encapsulation materials have refractive indices in the range of -1.45-1.55 and, when properly shaped to form a dome lens around the LED chip, can effectively increase light extraction efficiency from the chip on the order of 1S2> 2x by coupling light efficiently from the epoxy to air. Thus, for an AS LED lamp, light extraction through the top surface (neglecting top contact occlusion) increases to 4-5%. For applications such as indicator lamps and seven-segment displays, such low efficiencies are tolerable and, indeed, the vast majority of today’s LED market is satisfied by devices with efficiencies considerably lower than 4-5% (mostly not AlGaInP). There are other applications, however, most of them emerging, wherein efficiency is a premium. These applications, such as traffic signaling, require power LED lamps to provide a large amount of useful flux at efficiencies comparable to or higher than what can be provided by incandescent sources. For such applications, which require many LEDs to produce the required flux, increased efficiency provides a double benefit. On the one hand, the increased efficiency reduces the required input electrical

-

-

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H. CHUIt?t al.

power for the application, which can result in significant energy savings. Second, the same increase in efficiency can reduce the number of LEDs needed to provide the required flux for the application, hence reducing piece-part materials and assembly costs associated with the system. For such applications, the headroom for potential increase in AlGaInP LED extraction efficiency (from 4-5% to 50% or more?) is very attractive, and considerable recent effort has been put forward by AlGaInP LED manufacturers to push the technology and maximize performance. As mentioned in Subsection 4 of Section 111, a considerable amount of side light can be extracted from an LED if the top window layer of the device is sufficiently thick. In addition to providing improved current spreading, the window thickness enables light emitted from the active layer at large angles (relative to vertical) to impinge in an escape cone at the four sides of the window layer. A simple argument -assuming isotropic emission from the active layer, a thick window layer, and considering only first-pass light -predicts about a threefold increase in extraction efficiency relative to a thin window device. Each side surface provides one-half of an escape cone and, included with the top surface emission, allows 4 x (1/2) + 1 = 3 escape cones for the photons generated within the material, compared to one cone for a thin-window AS device. The escape cone model for a thin- and thick-window AS structure is illustrated in Figs. 7a and 7b. Using this model the maximum extraction efficiency for an epoxy dome-encapsulated, thickwindow device becomes 12-15%. Indeed, high efficiencies have been observed in thick-window LEDs (Huang et al., 1992) with measured lamp external quantum efficiencies qext exceeding 12'7'0, putting this as a lower limit on the extraction efficiency C,, (Krames et al., 1998). In addition to side light extraction, considerable effort has been put into retrieving a good portion of the downward-emitted light in AS AlGaInP LEDs. This light, which would normally be absorbed by the substrate, may be partially recovered by providing a distributed Bragg reflector between the active layer and the GaAs substrate. As illustrated in Fig. 7c, the theoretical benefit is the recovery of the downward escape cone to increase the extraction efficiency by another 4-5% (a twofold improvement for thinwindow LEDs and a -33% relative improvement for thick-window devices). In practice, the benefits are smaller than theoretically predicted for several reasons. One challenge is the low refractive index contrast provided by the limited group of materials that are both lattice matched to GaAs and transparent to the AlGaInP LED emission wavelength. AlInP/AlGaInP DBRs exhibit index contrasts An less than -0.4 with decreasing contrast at shorter wavelengths due to the increased A1 composition of the low-index layer required for transparency. This low-index contrast requires a DBR to have many layer pairs to achieve high reflectivity. The large number of pairs,

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cex-12-15%

Tap Cons + BottunCbns

-

FIG. 7. Light extraction efficiency of various device structures using the “escape cone” model. The numbers assume that an epoxy ( n 1.5) domed lens is provided to increase light extraction from the chip into air.

in turn, reduces the reflectivity bandwidth (in terms of wavelength or angular extent), such that capture of the full-escape cones is not achieved (Kish and Fletcher, 1997). A higher index contrast is provided by AlInPGaInP and AlInP-GaAs DBRs, which are lossy due to absorption in the low-bandgap GaInP and GaAs layers. While peak reflectivities are relatively low due to the optical loss, the bandwidths are quite large. Attempts to combine the benefits of high-reflectivity transparent DBRs with high-bandwidth lossy DBRs has led to hybrid AlInP-AlGaInP + AlInP-GaAs mirrors. AlAs- AlGaAs mirrors provide the highest index contrast and have been shown to provide high reflectivity even at wavelengths as short as 550 nm (Young et al., 1992). Even so, the DBR bandwidth is limited and in practice reflects only a fraction of the total downward escape cone back

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a!.

toward the top surface. Another approach has been to include “chirped” DBRs, wherein different sections of layer pairs are included for high reflectivity at angles at oblique incidence in an effort to capture the full downward escape cone for an encapsulated device. The chirped structure has resulted in devices of increased light output relative to devices employing conventional DBRs (Chang et at., 1996). An important point to remember concerning devices employing DBRs is that the reflected light must pass a second time through the active layer, and a significant fraction of this light is reabsorbed: 1 - exp( - EL), where CI is the absorption coefficient at the photon energy (typically lo4 cm-’) and L is the photon path length through the active layer. The amount of reabsorbed light can be significant. For a typical active layer thickness of 0.5 pm, about 40% of the light reflected from the DBR would be reabsorbed in the active layer. Depending on the internal quantum efficiency of the active layer, this light may or may not be re-emitted. In shorter wavelength devices where internal quantum efficiency is low, most of it is not. Because of this and compromises between DBR reflectivity and bandwidth mentioned earlier, in reality the DBR approach does not provide the ful1 intended benefit of retrieving the downward escape cone. Furthermore, the DBR approaches described do not provide for recovery of the downwardemitted side cones. At the time of this writing the best reported external quantum efficiency for AlGaInP LED lamps utilizing DBRs is -7% at 610 nm (Itaya et at., 1989). A third generation of AlGaInP LEDs has been provided in TS devices. For AlGaInP LEDs, this was achieved at Hewlett-Packard Company by the development of a process in which the absorbing GaAs growth substrate is removed and replaced with a wafer-bonded transparent GaP substrate (Kish et al., 1994). The extraction efficiency of such a wafer-bonded AlGaInP device is quite high, approaching that of TS devices based on much more mature technologies such as GaAs emitters. Using the simple cone argument as illustrated in Fig. 7d, the wafer-bonded TS AlGaInP LED should provide up to -30% extraction efficiency if most of the downward escape cone can be recovered using a reflective back contact and/or DBR. External quantum efficiencies as high as 23.7% have been reported (Vanderwater et al., 1997). Indeed, we measured a 637-nm device with a peak external quantum efficiency of 29.9% at 100 mA (q = 2 5 T ) (Krames et al., 1998). Presently, transparent substrate wafer bonding provides the world’s highest efficiency commercially available LEDs in the red (650 nm) to yellow (580 nm) color regimes. In the case of a TS device structure, the simple arguments using only first-pass light to estimate extraction efficiency become inaccurate. In such devices, a significant portion of light may instead escape the chip after

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several bounces within it. This light is referred to as multiple-pass light. Because of the general roughness of the device geometry (e.g., sawn sidewalls, etc.), multiple pass light is likely rerandomized in orientation at each contact with the semiconductor-encapsulant interface. Also, for material of relatively high internal quantum efficiency, photons reabsorbed by the active layer may be reemitted. These randomizing events provide extra opportunities for photon extraction. In the case of no parasitic losses within the device, this photon randomization would allow for a 100% extraction efficiency! In practice, extraction of multiple pass light is severely dependent on the parasitic loss mechanisms present in an LED device structure. Indeed, even with extremely high quantum efficiency active layers (qin,-99.7%), TS AlGaAs devices with merely 8% estimated optical loss per pass through the structure have been measured to provide less than 25% external quantum efficiency (Schnitzer et at., 1993a). The same should be true of AlGaInP devices. The excessive reduction in external quantum efficiency is due to the relatively long path length of photons trapped within the high-refractive-index chip, coupled with the presence of many parasitic losses such as absorption at ohmic contacts, free-carrier absorption, and annihilation in the active layer. To reduce the effective path length of photons inside the chip, surface texturing has been employed to randomize the photons each time they encounter, for example, the top surface of the LED chip. Such surface texturing of TS AlGaAs devices has been shown to increase the air-gap external quantum efficiency by approximately threefold relative to devices without texturing (Schnitzer et al., 1993b). The gain would be considerably lower in end-user devices (encapsulated with epoxy domes) since the photon path lengths are considerably reduced in this case. TS AlGaInP devices employ rough (sawn) sidewalls and back surfaces that provide some randomization of the light. It is not yet clear what benefit additional surface texturing may provide for light extraction in TS AlGaInP LEDs beyond the external quantum efficiencies measured to date. Note, however, that recent thin-film LEDs employing GaAs-active regions (qint approaching 100%) have achieved external quantum efficiencies of 31% (without an epoxy dome encapsulation) in the infrared (850 nm) by the use of a textured top surface (Windisch et al., 1998). An altogether different approach to increase the light extraction in LEDs is to alter the shape of the light-emitting device so that all or most of the light generated within the device enters into an escape cone at the interface with the ambient. In a sense, the number of escape cones is increased. The ideal shape is a spherical device with a point-source light-emitting (active) region located at the center of the sphere. All light that is emitted would strike the device-ambient interface at normal incidence, thus escaping into the ambient with no total internal reflection. More practically, the device

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would be hemispherical with an outer radius sufficiently larger than the radius of the active region within, and would employ a highly reflective mirror on the flat side of the device. Such altered-geometry devices have been fabricated early on in the history of LEDs, with improvements in extraction efficiency observed (Dierschke et al., 1971; Carr, 1966). However, complex shaping techniques and inefficient device material utilization make this approach doubtful for high volume manufacturing and today no commercially available visible-wavelength LEDs are fabricated using this approach. In a section devoted to the topic of light extraction, a final device structure deserves comment and some discussion. Resonant cavity LEDs (RCLEDs), which employ a vertical cavity structure to enhance the radiative recombination rate in the vertical direction, provide an interesting attack on the light extraction problem in that side light and multiple-pass light are essentially ignored. Instead, one tries to devise a high-Q cavity wherein the high radiative recombination rate is matched to the vertically radiating (called “normal”) modes of the device, which are within the escape cone for extraction out the top surface. To the extent the resonant cavity increases the radiative rate into these normal modes (relative to recombination into lossy modes), external quantum efficiency can be increased. External quantum efficiencies as high as 20% have been measured for InGaAs infrared RCLED chips in air, with predicted extraction efficiencies approaching 50% with epoxy encapsulation (De Neve et al., 1997). On the other hand, the difficulties of providing high-reflectivity DBR mirrors with large bandwidths may limit the performances achievable by an RCLED structure in the AlGaInP material system, especially for shorter wavelength devices. Also, it is not clear whether such devices will prevail over standard TS structures for raw flux efficiency even for InGaAs infrared emitters. At the same time, RCLED structures provide desirable characteristics, such as top-surface-only emission. Also, unlike the case for conventional LEDs, which depend on side light utilization, the efficiency of RCLEDs should remain high when scaling up to the larger junction areas necessary for power devices. A detailed discussion on RCLEDs is provided in Chapter 9.

IV. Epitaxial Growth 1. INTRODUCTION Having discussed the material properties of AlGaInP and AlGaInP LED device design, we next consider the growth of AlGaInP LEDs followed by device fabrication and performance. The conventional epitaxial growth

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processes for LEDs, such as LPE and HVPE, are not suitable for the growth of Al-containing material systems such as AlGaInP. This is mainly due to the fact that AlP is much more thermodynamically stable than InP, and compositional control is therefore very difficult. This thermodynamic problem can be overcome by using a kinetically controlled process, such as metalorganic vapor phase epitaxy, molecular beam epitaxy (MBE), or gas source molecular beam epitaxy (GSMBE), in which the growth precursors or sources are present at a very high level of supersaturation in the growth ambient (Chen, C . H., et al., 1997). While the growth of phosphoruscontaining materials by MBE is difficult due to the high vapor pressure of phosphorus (Stringfellow, 1989), the MOVPE technology has advanced significantly since the 1980s and has become established as the preferred high-volume manufacturing technique for the synthesis of device quality AlGaInP alloys. This section describes the MOVPE growth of AlGaInP for LED applications. A brief overview of the MOVPE process and reactor system is presented first. The precursors for the synthesis of AlGaInP materials, as well as the n-type, p-type, and unintentional hydrogen and oxygen incorporation issues are then addressed. The important topic of AlGaInP atomic ordering is discussed, followed by the growth conditions that are required to achieve high-quality AlGaInP materials and devices. Finally, manufacturing and safety issues are covered.

2. GROWTH OF ALGAINPMATERIALS a.

MOVPE Growth Process and System

Initially demonstrated by Manasevit and Simpson (1969), the MOVPE process has evolved into a versatile technique for compound semiconductor growth that provides materials with high crystalline quality and atomically abrupt interfaces. The typical MOVPE process for the growth of AlGaInP alloys uses metalorganic (alkyl) compounds such as trimethylaluminum (TMAl), trimethylgallium (TMGa), and trimethylindium (TMIn) as the group 111 sources, and PH, as the group V source. The alkyls and PH, are diluted in hydrogen and decompose near the heated substrate to produce Al-, Ga-, In-, and P-containing species. An excess amount of PH, relative to group 111 alkyls is usually employed to stabilize the AlGaInP surface with the volatile group V element. The growth process can be viewed thermodynamically as a gas-solid equilibrium reaction, where the supersaturated ambient vapor phase provides a thermodynamic driving force for growth to occur. The maximum growth rate would be obtained if a sufficient

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amount of time were allowed for all the reactants to decompose and equilibrate with the substrate. In practice, MOVPE is not carried out under equilibrium conditions, and the realized growth rate, at a typical value of 3 ,um/h, is less than that predicted by thermodynamics (Stringfellow, 1989). The MOVPE growth process is a consequence of complex and coupled chemical, thermal, hydrodynamic, and mass-transport processes. The deposition region of the reactor is typically operated under laminar flow conditions (such as reduced pressure) to minimize local variations in the growth ambient due to turbulent flow. A concentration boundary layer is formed over which the reactant concentration changes from the mean stream value to a reduced level at the growth front. The boundary layer thickness increases along the direction of gas flow as the deposition depletes the reactive species in the gas phase. This depletion leads to a decreased mass flux to the growth front and, therefore, nonuniformity over a large substrate area. Improvements in uniformity have been made by tilting the susceptor to increase gas velocity and reduce this gas depletion effect, or by susceptor rotation during growth (Kuech, 1987). The growth temperature is usually chosen to be between 650 to 85OoC, where the growth rate is controlled by the diffusion of group II[ species to the growth front. This mass-transport-limited step leads to a regime where the growth rate is nearly temperature-independent (Stringfellow, 1989). The MOVPE system is composed of three main building blocks: the gas manifold, the reactor chamber, and the exhaust-handling system. The gas manifold is designed to controllably meter and mix the precursors in the carrier gas (H, or N,) before they enter the reactor. The group V hydrides are stored in high-pressure cylinders, and pressure regulators and electronic mass flow controllers are used to reduce the pressure and control the flow rate respectively. The group I11 metalorganic precursors are usually highvapor-pressure liquids or solids at room temperature and are stored in bubbler delivery containers. The partial pressure of a metalorganic species P,, is controlled by the temperature of a thermal bath in which the bubbler is situated. The total pressure P, over the source is maintained by a pressure controller at a value much higher than P,, to prevent boiling of the source. The carrier gas is metered and bubbled through the source at a flow rate dcg,entraining and transporting the source species into the reactor. When dcgis low enough to maintain the vapor-liquid or vapor-solid equilibrium, the volume flow rate of a precursor leaving the bubbler dm0is given by (Kuech, 1987)

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The dopants are either gases, such as Si,H,, mixed to the parts per million (ppm) level in H,, in high-pressure gas cylinders, or are liquids, such as diethyltellurium, in bubblers. All these gas constituents are further diluted and mixed within the gas manifold or reactor. The mole fraction of individual species in the growth ambient is determined by the ratio of its flow rate over the total flow rate through the reactor. The gas mixture is typically 90% H,. To grow multilayer structures with sharp interfaces, a switching manifold is usually used where the flow is selectively switched between the vent line, to bypass the reactor, and the run line, to enter the reactor. The group V hydrides are highly toxic and most of the metalorganic precursors are pyrophoric. These potential hazards necessitate the use of all-welded stainless steel lines together with bellow-sealed high-vacuum valves. These valves, as well as the mass flow and pressure controllers, are integrated into the system with fittings containing only metal-to-metal gasket seals and the entire system must be tested to be leak proof. All the sources are located inside ventilated enclosures and a toxic gas monitoring system is installed to detect any trace amount of hydride release. These features not only ensure a safe operating environment, but are also important in maintaining high material quality by preventing contaminants from entering the reactor. The MOVPE reactor can have gas flow parallel or perpendicular to the growth surface. Substrates are placed on a graphite susceptor, which is inductively heated by a radio frequency generator or resistively heated by lamps. The temperature control and monitoring are accomplished by a temperature controller integrated with a thermocouple wire or an infrared pyrometer. The typical low operating pressure of the reactor is achieved by the use of a mechanical vacuum pump and a throttle valve controller. Even when the reactor geometry allows for growth at atmospheric pressure, a vacuum system is still required to regulate and maintain the reactor pressure to give reproducible results. Particulate flakes are generated by deposition occurring on the reactor wall and graphite susceptor. These flakes are usually removed physically by spare-parts replacement or by in situ HC1 gas etching. The effluent stream of the reactor, containing toxic materials that are not consumed completely in the reactor, is typically run through a scrubber or burner before being released to the ventilation duct. The use of PH, requires special considerations since it generates elemental phosphorus, which is toxic and flammable. In this case, the exhaust system is usually heated to prevent phosphorus condensation, and to facilitate the transfer of phosphorus down the line to be physically collected or chemically converted into nonhazardous byproducts.

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Source Materials

As previously mentioned, the group I11 metalorganic precursors for AlGaInP alloys are typically trimethyl-based sources (TMAl, TMGa, and TMIn). One potential problem with these alkyl sources is the intrinsic carbon contamination from the methyl group, as in the AlGaAs system (Kuech et al., 1987). Carbon incorporation is very inefficient in In-containing alloys (Chen, C. H., et al., 1997), however, especially for the GaAs lattice-matched (Al,Ga, -,)o,51n,,,P material system, where In is a major constituent. The delivery of TMAl and TMGa using the bubbler technique is relatively reliable, since TMAl and TMGa are both liquids at room temperature (their melting points are 15.4 and - 158°C respectively). In contrast, TMIn has traditionally been used as a solid source with a melting point of 88.4"C, and TMIn source pick-up is known to decrease over the life of the source due to the decrease in internal surface area of the solid TMIn surface as the source becomes depleted (Stringfellow, 1989). TMIn has more recently become available in solution form, where TMIn has been dissolved in a solvent (such as a very low volatility aliphatic amine) to form a liquid adduct. The evaporated TMIn is thus continuously replenished by dissolution, and a consistent carryout can be maintained (Epichem, 1998). Due to their high purity and low cost, the most common group V sources for AlGaInP alloys are ASH, (for GaAs or AlGaAs growth, as discussed later) and PH,. The extreme toxicity of ASH, and PH, has, however, prompted increasing research efforts and acceptance of alternative precursors such as tertiary butyl arsine and tertiary butyl phosphine (TBA and TBP) (Chen et al., 1987, 1988). The most commonly employed n-type dopants for AlGaInP alloys are silicon (Si) and tellurium (Te). The Si source is usually delivered from a parts-per-million-level mixture of silane (SiH,) or disilane (Si,H,) in H, stored in a high-pressure cylinder. Diethyltellurium (DETe) is commonly used to deliver Te, either from a bubbler or from a parts-per-million-level mixture with H,. Zinc (Zn) and magnesium (Mg) are typically used as p-type dopants in MOVPE, and Zn and Mg delivery is usually accomplished through the use of dimethylzinc (DMZn) and bis(cyclopentadieny1) magnesium (Cp,Mg). Alkyl and hydride source purity, and the unintentional incorporation of residual oxygen into AlGaInP alloys is a very important aspect of AlGaInP growth by MOVPE. The electrical and optical properties of the A1 containing AlGaInP material system, similar to the AlGaAs or AlInAs system, are commonly found to be adversely affected by the presence of oxygen in the precursors. These observations are attributed mainly to the strong interac-

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tion between oxygen in the growth ambient and the metalorganic Al-bearing precursors. To minimize this interaction, the use of palladium-purified hydrogen as the carrier gas has become a common practice. A point-of-use ASH, purification system, based on the trimetal melt of Al-Ga-In with A1 as the oxygen getter, has been widely used with a corresponding improvement in AlGaAs material quality (Shealy et d., 1983). Effective oxygen gettering processes have also been developed and successfully applied to the purification of TMAl and TMGa for the growth of AlGaAs (Maeda et al., 1991; Roberts et al., 1994b). These efforts, as well as many others, have greatly advanced the performance of MOVPE-grown AlGaAs devices. While AlGaInP material quality may benefit from these improvements in TMAl and TMGa purity, the purity of TMIn still has room for improvement. Recent investigations of unintentional oxygen incorporation in AlInP and AlInAs have identified TMIn and PH, as the two major sources of oxygen contamination (Chen, J. C., et al., 1997; Roberts et al., 1994a). The impact of residual oxygen on AlGaInP LED performance is illustrated in Fig. 8, where a higher LED light output is observed for devices with a lower residual oxygen concentration in the AlInP confining layer (Huang et al., 1997). The variation in oxygen level is solely due to the use of different TMIn sources, while all other variables are kept constant. The variability in

=!

loo00

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rn

A

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. 1E+18

Oxygen Concentration (cm?

FIG. 8. LED light output performance as a function of oxygen concentration in the AlInP confining layer. The variation in oxygen concentration in these samples is a result of different residual oxygen contamination in the five different TMIn bubblers, labeled A through E, that were used to grow these samples. Note that the light output of the shorter wavelength LEDs (585-587 nm) is more susceptible to changes in background oxygen than that of the longer wavelength samples.

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TMIn quality is probably due to the solid nature of TMIn, which makes purification difficult. Bis(cyclopentadieny1) magnesium, which is in solid form at room temperature (melting point 176"C), can also be a source of oxygen for the same reason. It was recently reported that the oxygen content in Cp,Mg has been drastically reduced via an alternative synthesis route without the use of oxygen-containing ether-based solvents (Bartram, 1998).

c. Dopant Incorporation

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Undoped GaInP and AlInP typically have an n-type background carrier concentration in the low 10'' cm-3 and 1 x 10l6 cm-3 range, respectively, due to residual Si in the alkyl and hydride sources. LED applications require the intentional doping of both n-type and p-type (Al,Ga, - x ) o , 5 In,,,P confining layers where 0.6 6 x 61 1. Early attempts to dope high-Alcontent AlGaInP alloys were not very successful, especially for p-type alloys (Bour, 1987).These early difficulties in doping AlGalnP may have been due to the high oxygen content in the early AlGaInP materials, since oxygen is known to introduce deep level centers in AlGaInP that compensate shallow impurities (McCalmont et al., 1992). Higher bandgaps with higher A1 compositions also result in higher acceptor ionization energies, reducing free-hole concentrations. An additional reason for poor acceptor activation in AlGaInP alloys was found to be unintentional hydrogen passivation, which can be eliminated by postgrowth annealing (Minagawa et al., 1992). With the aforementioned understanding, both n-type and p-type doping of AlGaInP is now routinely achieved. The dopants with high elemental vapor pressures, such as Te and Zn, exhibit a decrease in doping level with an increase in growth temperature due to reevaporation of the atoms from the growth surface. These dopants also show an incorporation dependence on V-TI1 ratio. The group I1 element, Zn, residing on a group I11 lattice site, exhibits an increase in doping as the V-I11 ratio is raised, while the group V1 element, Te, which substitutes on a group V site, shows the opposite trend. This V-I11 dependence may result from thermodynamic processes, such as gas-solid or atom-vacancy equilibrium, or it may arise from gas phase interactions between the dopant and the group V source. Dopants with a low vapor pressure, such as Si, can exhibit various doping behaviors, depending on the exact precursor that is used. Due to the low vapor pressure of Si, it can be assumed that all the Si atoms that reach the surface remain there, and the incorporation of Si would then be limited by (1) gas phase or surface kinetics or (2) by the gas phase diffusion of Si species to the growth surface. In case 1, an Arrhenius-type temperature dependence would be noted, as seen in the use of SiH,. In case 2, a weak temperature

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dependence would be observed, as seen in the use of Si,H, (Kuech et al., 1984). n-type Doping The n-type dopants of choice for AlGaInP materials are Si or Te (Ishikawa et al., 1986; Kobayashi et al., 1985). Since Si exhibits very low diffusivity in the epilayer, and neither SiH, nor Si2H, has a memory effect, the doping profile and junction position can be easily controlled with Si doping. In contrast, Te exhibits a doping tail after the source supply is turned off. Both Si and Te lead to deep level centers (Suzuki et al., 1991; Cederberg et al., 1998) that may have adverse effects on LED performance. In contrast, deep levels are not a problem for p-type dopants, and it is therefore a common practice to have a p-type active region and inject electrons from a wider bandgap n-type confining layer (Craford et al., 1991). p-type Doping The most commonly used p-type dopants are Zn from DMZn and Mg from Cp,Mg. The effectiveness of Zn as a p-type dopant was found to decrease dramatically with increasing A1 composition in (Al,Ga, -x)o,51no,5P, and, as illustrated in Fig. 9, the highest hole concen~ . is due to the tration in Zn-doped AlInP is only -2 x 10'' ~ m - This reduced Zn electrical activity (Nishikawa et al., 1988), as well as an increasing oxygen content with higher A1 composition (Nishikawa et al., 1992a). Zn is also very volatile under the typical AlGaInP growth temperatures, and, as shown in Fig. 9, a high molar flow rate of DMZn, which is comparable to or higher than that of the group 111 precursor, is usually required to achieve a high doping level in the 10'' cm-3 range. This

0.1

1

10

[DMZn]/[TMIn+TMAI+TMGa] FIG. 9. Hole concentration of ( A l x G a l ~ x ) ~ ~ 5 1 n(xo ~=50,P 0.4, 1) as a function of [DMZn]/[III].

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behavior requires frequent DMZn bubbler changes in large-scale manufacturing reactors. An additional problem with Zn is its high diffusivity, which complicates p-type doping control. In contrast to the low hole concentrations observed in Zn doped AlInP, free-hole concentrations as high as 1 x lo’* cm-3 have been reported in Mg-doped AlInP (Ohba et al., 1988). The correlation between AlInP hole concentrations and molar ratios of Cp,Mg vs group I11 precursors is shown in Fig. 10. Note by a comparison with Fig. 9 that the required molar flow rate of Cp,Mg to achieve the 1 x 10’’ to 1 x lo’* cm-3 ranges is much less than that of DMZn. This observation is consistent with the difference in vapor pressures of elemental Mg vs Zn. While Mg also has a lower diffusion coefficient than Zn (Nishikawa et al., 1992b), p-type doping control is still problematic due to the Mg memory effect. The Mg doping profile was found to be delayed with respect to the turn-on time of the Mg source (Nishikawa, et al., 1992b), possibly due to the adhesive nature of Cp,Mg on the internal surfaces of the gas handling system (Kuech et al., 1988). This delay can be mitigated by the presence of TMAl to help occupy the surface sites in the gas manifold and reactor chamber (Nishikawa et al., 1992b; Kondo et al., 1994a). A Mg-doping tail has also been observed after the Cp,Mg source was turned off (Wu et al., 1994). Hydrogen Passivation Hydrogen, introduced into the MOVPE growth ambient via PH, or as a carrier gas, passivates shallow acceptors in AlGaInP materials during growth or postgrowth cool down. For the cases

FIG. 10. Hole concentration of AlInP as a function of [Cp,Mg]/[III].

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of Zn or Mg, the hydrogen atom is believed to be bound to a phosphorus atom and reside between the acceptor and the phosphorus neighbors. The hydrogen atom satisfies the electron pair requirement of the phosphorus neighbor, and therefore the entire hydrogen-impurity complex is neutral (Schubert, 1993). Minagawa et al. (1992) studied the effect of various cooling ambient (H,, PH,, or ASH,) and capping layers (n-GaAs cap or no cap) on Zn activation in MOVPE AlGaInP. Hydrogen passivation was found to be the least severe in H,, and the most severe in ASH,, presumably due to the better cracking efficiency of ASH, to produce atomic hydrogen. The n-GaAs cap also reduced hydrogen passivation, presumably due to a lower hydrogen solubility limit in n-type GaAs, and the tendency of the n-type GaAs to act as a hydrogen diffusion barrier during cool down. The reduced hydrogen concentration with an n-GaAs cap was verified by secondary ion mass spectroscopy (SIMS) measurements (Nishikawa et al., 1992a). Reactivation of hydrogen passivated dopant species can be achieved by post-growth annealing in H, or N, ambient at temperatures above 400500°C (Minagawa et al., 1992). During annealing, the acceptor-hydrogen complex dissociates, followed by the outdiffusion of atomic hydrogen. Hamada et al. (1992) reported that the degree of passivation is independent of the substrate orientation, but increases with A1 composition over the range of x = 0 to x = 0.65 in (Al,Ga, -,)o,51n,.5P. Since high-Al-composition-layers are typically used in LED confining layers, proper postgrowth thermal treatment is especially critical in such structures. Oxygen Incorporation The common presence and reactivity of oxygen makes it a likely contaminant in all semiconductors. This is especially true for Al-containing compound semiconductors due to the strong affinity between oxygen and metalorganic aluminum sources. Oxygen related defects are believed to be the major source of electrically active, carriertrapping centers for nonradiative recombination processes. MOVPE grown AlGaAs has long been reported to be highly compensated and have low photoluminescence efficiency. Similar effects are also observed in the (Al,Ga, -,)o,51n,,,P material system. In an LED or laser structure, oxygen incorporation is usually very efficient in the confining layers, which have high-A1 compositions (x 0.6-1). In the active layer, the oxygen level in amber and green devices with x - 0 . 3 to 0.5 is expected to be much higher than for red devices with x-0 to 0.1. This increased oxygen content together with the decreased carrier confinement, as well as the approaching directindirect bandgap transition have made the realization of high-efficiency green and amber AlGaInP devices quite difficult. This concept is illustrated, in part, in Fig. 8, where the shorter wave length LED light output is more sensitive to the background oxygen level.

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Two approaches must be simultaneously taken to reduce oxygen incorporation to an acceptable level, First, the sources of oxygen must be identified and controlled. These sources include oxygen from source materials and the carrier gas as previously discussed, leaks in the MOVPE system, and contaminated reactor components. Second, the growth conditions, including growth temperature, V-I11 ratio, and substrate orientation must be optimized to suppress oxygen incorporation. Controlled oxygen incorporation, using 0, or diethyl aluminum ethoxide (DEA10) (Huang et al., 1994), has also been used as a diagnostic tool (McCalmont et al., 1992; Kondo et al., 1994b; Stockman et al, 1999) to better understand oxygen incorporation and optimize LED growth conditions. Nishikawa et al. (1993) investigated the effects of growth parameters on residual oxygen incorporation in MOVPE (Al,Ga, -,)o~sIno,5P,and reported that the oxygen concentration increased at higher A1 compositions, and decreased significantly at higher V-I11 ratios for x = 0.7 and 1.0, while a temperature increase from 650 to 750°C induced only a slight reduction of oxygen level. Compensation of shallow Zn acceptors due to oxygenrelated deep levels was also observed. Suzuki et al. (1993) identified two oxygen-related deep levels at 0.64 and 1 eV below the conduction band edge ( E c ) . They also reported a drastic reduction in the concentration of these deep levels when a (100) substrate with 15” tilt toward (111)A was used in contrast to a nominally (100) substrate. The first report of intentional oxygen incorporation in AlGaInP was by McCalmont et al. (1992), who grew semi-insulating, Si-doped (Alo~4Gao,6)o,51no,5P with a dilute 0, source, and found a square-root dependence of oxygen incorporation on 0, flow rate. A much stronger incorporation dependence of [O] K F(02)”, where n = 4, was reported by Kondo et al. (1994b) in (Alo,,Gao,3)o,51no,5P. A more recent study of AlInP doped with either 0, or DEAlO revealed a similar superlinear dependence with n ranging from 2 to 7, depending on growth conditions (Stockman et al., 1999). Kondo et al. (1994b), using isothermal capacitance transient spectroscopy, have unambiguously identified two oxygen-related deep levels at 0.46 and 1 eV below E , for (Alo,,Gao~3)o,51no,5P. The total concentration of these two levels, however, was always less than 10% of the total oxygen concentration, suggesting that most of the oxygen is electrically inactive. The oxygen concentration, as well as the measured deep level concentrations were also verified to be significantly reduced by decreasing the A1 composition, using a higher growth temperature, a higher PH, partial pressure, and a (100) GaAs substrate misoriented toward the (1 1l)A surface. The preceding discussions clearly demonstrate the striking similarity of oxygen incorporation in AlGaInP quaternary alloys and the well-studied

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AlGaAs system (Kuech et al., 1987, 1992; Chad et al., 1991), despite the differences in gas phase precursors (e.g., PH, vs ASH,) and solid phase composition (e.g., In, Ga, and A1 vs A1 and Ga). The strong dependence on A1 composition is due to the strong A1-0 affinity. The enhanced evaporation of volatile metal suboxides and the reduced thermodynamic driving force for oxygen incorporation suggest decreasing oxygen incorporation with increasing temperature. This observation is very similar to the incorporation behavior of dopant species with a high elemental vapor pressure, as discussed earlier. The effect of V-I11 ratio can be attributed to the atomic hydrogen, from either gas-phase or surface-adsorbed PH, --x, that reacts with oxygen-bearing species and forms a volatile reaction product, preventing oxygen from incorporating into the growing crystal. An alternative explanation is to draw an analogy from the doping behavior of the group VI elements on the V-I11 ratio effect, as the competition of oxygen with phosphorus would lead to the reduced oxygen level at a higher V-I11 ratio. The reduced oxygen concentration on a (100) GaAs substrate misoriented toward the (1 1l)A surface suggests that the phosphorus sticking probability was increased on the group-III-type step edges. The subsequent decrease in the number of available phosphorus vacancies, which serve as sites for oxygen incorporation, leads to reduced residual oxygen. High concentrations of oxygen in AlGaInP can result in electrical compensation of shallow donors or acceptors. This effect is illustrated in Fig. 11 where DEAlO has been used to controllably introduce oxygen into Mgdoped p-type AlInP and Te-doped n-type AlInP (Huang et al., 1998). Two multilayer samples were grown with the Cp,Mg and DETe molar flow rates kept constant for all the layers. The hole and electron concentrations are 5.2 x loi7 and 6.1 x l O I 7 ~ m - ~ respectively, , if no DEAlO doping is provided, and the oxygen background concentration without DEAlO is below 2 x loi6 cm-, in these layers. The controlled incorporation of oxygen via DEAlO leads to compensation in both cases. In the n-type case, the ratio of compensation to the total oxygen concentration is -2O%, and an oxygen related deep level -0.9 eV below E , has been identified (Cederberg et al., 1998), consistent with the results of Kondo et al. (1994b). In contrast, the compensation is much more effective in Mg-doped AlInP. An enhancement of Mg incorporation due to oxygen doping has also been observed, as illustrated in Fig. 12 using a similar multilayer structure with a constant Cp,Mg molar flow rate for each sample. The enhancement effect is obvious only at higher Mg or 0 concentrations. One possible explanation for this enhancement is that the strong affinity of Mg to 0 results in the formation of MgO, and leads to electrically inactive complexes

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1E+17

1E+18 Oxygen Concentration (cm-3)

1E+19

FIG. 11. Room-temperature free-carrier concentrations (measured by electrochemical capacitance-voltage measurement) vs oxygen concentration for Cp,Mg and DETe-doped AlInP. Oxygen doping is provided by DEAIO. The oxygen background concentration without DEAlO is below 2 x lo’&cm-3 in these layers. The Cp,Mg and DETe molar flow rates are kept constant for all the layers. The hole and electron concentrations are 5.2 x 10’’ and 6.1 x 10’’ ~ m - respectively, ~ , if no DEAlO doping is provided.

~~

-chigher Cp,

T

E

A

flow rate

lower Cp, Iv(s flow rate

Y

----I

A

1E+17

L

1 ’ ~

1E+15

~

iE+16

I1

I

1~ ”

M ~ /

1E+17

1E+18

1E+19

o w e n Concentration (cm-7 FIG. 12. For a fixed Cp,Mg molar flow rate, the Mg incorporation is enhanced with increasing oxygen concentration in AlGalnP. Intentional DEAlO doping was used to controllably change the oxygen concentration at values above 1 x lo”.

2 HIGH-EFFICIENCY AlGaInP LIGHT-EMITTING DIODES

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(Schubert, 1993). While this Mg-0 related defect could originate from the Cp,Mg source, or from the reaction of Cp,Mg with any oxygen or water vapor in the growth system, the weaker Mg-0 bond (86.8 kcal/mole), as compared to the AI-0 bond (122.4 kcal/mole) (Lide, 1990), and the much lower concentration of Mg, compared to that of Al, in both the crystal and in the gas phase, make this explanation seem unlikely.

d.

Ordering

The MOVPE AlGaInP material system exhibits CuPt-type spontaneous atomic ordering. The related issues have been reviewed by Zunger (1997), Stringfellow (1997), and Suzuki (1997). In general, the group I11 atoms are not distributed randomly on the group I11 sites, but rather in alternating Al-Ga- and In-rich (111) planes. In this way, the total formation energy of random alloys composed of constituents that have a large lattice parameter difference can be reduced without producing a large macroscopic strain. This (111) superlattice due to the ordering planes leads to a reduction of the Brillouin zone by folding the L point into r. The most familiar consequences of the zone folding on the band structure are the bandgap reduction and the valence band splitting. The variation of AlGaInP bandgap energy due to the degree of ordering has a direct impact on the color control of LEDs and lasers. Suitable growth conditions and postgrowth processing are therefore needed to minimize the effects of ordering. Ordering is considered to be a kinetically controlled process occurring on the growth surface. The key to ordering is the formation of a (2 x 4) surface reconstruction on group V terminated (001) surfaces with [110] rows of [T10]-oriented group V dimers. Su, et al. (1995) reported that the maximum degree of ordering in MOVPE GaInP occurred at a growth temperature of 620°C. A continuous decrease in ordering with increasing growth temperature or reduced V-I11 ratio is attributed to a loss of the (2 x 4)-like surface reconstruction. Decreased growth temperature also reduces ordering, due to an excess phosphorus phase at low temperatures. Misorienting the (001) substrate by a few degrees to produce [ i l O ] surface steps also reduces the degree of ordering, while [110] surface steps enhance ordering. Other kinetic factors may also be significant in the ordering process. For example, Cao et al. (1991) studied the effect of MOVPE GaInP growth rate on the degree of ordering, and reported that ordering is reduced at rates above -4 ,um/h, possibly due to the lack of time for atoms to reach the ordered state. The CuPt ordered structure is not stable in the bulk, and can be eliminated by postgrowth anneaiing.

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Growth Conditions

The growth of high quality (Al,Ga, - x)o,51no,,Pepilayers requires proper control to ensure lattice matching to the GaAs substrate. When the lattice-matching condition is met at the growth temperature, the AlGaInP epilayer will be compressively strained at room temperature due to the larger thermal expansion coefficient of GaAs relative to those of Alp, InP, and Gap, as previously shown in Table I. The choice of lattice match at the growth temperature, rather than room temperature, is preferred for several reasons. First, dislocations are highly mobile at typical growth temperatures, while they are much less likely to nucleate or glide at room temperature. Second, if the film is lattice matched at the growth temperature, any drift in In composition, which is difficult to control, is less likely to cause sufficient lattice mismatch to induce strain relaxation. The superior material quality of AlGaInP films grown lattice matched to GaAs at the growth temperature has been experimentally verified by Bour and Shealy (1998). The resulting strain at room temperature after cool down may, however, cause strain relaxation and introduce dislocations if t he film thickness exceeds the critical thickness. Figure 13 shows a plot of critical thickness at 300 K vs alloy composition for a Ga,In,-,P film grown on GaAs. For example, a Ga,In,-,P film that is lattice matched to GaAs at 850°C has a critical thickness of roughly 1200A at 25°C. As previously mentioned, the MOVPE growth temperature is typically chosen to be between 650 and 850°C, where mass transport is the rate2500

-

"=2000 u) u)

0 E

r .-v 1500 c I-

L

E 1000

m

I

500 0

CL--

0 40%

I

+--+-+ I -I-

42% 4% 4% 48% 50%

52% 54% 56%

I

3%

Ga % in GalnP

FIG. 13. Critical layer thickness for GaInP grown on GaAs.

2 HIGH-EFFICIENCY AlGaInP LIGHT-EMITTING DIODES

87

limiting step. In this regime, the growth rate is determined by the diffusion of group-III-bearing precursors from the gas phase to the sample surface, and to first order, the film composition is therefore linearly proportional to the input gas phase composition. An In/(Al + In) and In/(Ga In) mole ratio of 0.5 is therefore a good starting point to produce GaAs-latticematched AlInP and GaInP during growth. One complicating factor in achieving lattice match is the commonly observed In reevaporation from the surface at higher growth temperatures or In-depletion from the gas phase before reaching the wafer surface. These effects can require an In molar ratio greater than 0.5, especially at higher growth temperatures. Once the correct In molar ratio is determined, the AI/(Al + Ga) ratio can then be varied with a constant A1 plus Ga molar flow rate to achieve the desired bandgap for a particular AlGaInP emission wavelength. The V-I11 ratio during the MOVPE growth of AlGaInP is generally maintained at rather high values, greater than 200, mainly due to the inefficient cracking of PH, at typical AlGaInP growth temperatures (Larsen et al., 1987), and to the beneficial effect of reduced oxygen incorporation. High PH, flow creates heavy phosphorus loading in the exhaust handling system, however, and the degree of ordering may also be enhanced at a high V-I11 ratio, as discussed earlier. Finally, PH, flow also dictates the incorporation efficiency of most of the n-type and p-type dopants. Although a higher AlGaInP growth temperature would improve the decomposition efficiency of PH,, and suppress the oxygen incorporation, the issues of In reevaporation, phosphorus deposits in the exhaust system, and shallow dopant incorporation all must be considered to determine an optimal growth temperature window. Another important topic for the MOVPE growth of AlGaInP materials is the switching between ASH, for arsenide materials and PH, for phosphide materials. This switching is necessary since ASH, is required to stabilize the GaAs substrate surface while the reactor temperature is being ramped up to initiate growth of the AlGaInP layers. This switching is also required if an As-containing layer, such as an AlGaAs window layer or a DBR, is included in the device structure. A common practice of optimization usually involves adjusting the two time periods of hydride purges (after the present layer is finished and before the new layer is started) as well as the hydride concentrations. A hydrogen-only purge step is also included to minimize the formation of an interfacial, As-P mixed layer (Hergeth et al., 1991). Severe reevaporation of the group V element can occur if the hydrogen-only purge step is too long, while a layer with large lattice mismatch and poor quality can occur if both As and P are present at the same time. Quaternary GaInAsP layers a few monolayers thick have been reported for both GaInP-on-GaAs and GaAs-on-GaInP interfaces (Bhat et al., 1992). The

+

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H. CHUIet al.

intermediate GaInAsP layer present at the GaInP-on-GaAs interface is usually less pronounced than for GaAs on GaInP, due to the lower cracking efficiency of PH, relative to ASH,. Proper gas manifold design with the appropriate sequence of gas injection valves and pressure balancing for run-vent switching, and a reactor chamber with low residence time are very critical for rapid switching. Operating the MOVPE reactor under a reduced pressure (such as 0.1 atm) also helps to enable rapid gas switching. The use of low pressure further helps to reduce gas phase alkyl depletion, hence improving thickness and compositional uniformity (Chen, C. H., et ul., 1997). 3. GROWTH OF DEVICE STRUCTURES

Epitaxial growth of AlGaInP LED device structures can be separated into three main components: the double-heterostructure light-emitting portion of the device, a window layer for current spreading and light-extraction, and a distributed Bragg reflector as a back reflector. Only the D H is necessary for a functional LED, but as we discussed in Section 111, the other two components are commonly added to improve the efficiency of these LED devices. The epitaxial growth of these components is typically done by MOVPE, although the window layer can alternatively be grown by HVPE. In this section, we discuss the materials and growth issues of these components of the AlGaInP LED. u.

Double Heterostructures

MOVPE growth of the DH proceeds by sequential deposition of (A1,Gal -x)l -,In,P layers of various compositions and doping. The DH growth starts with the lower confining layer (LCL), which is composed of AlGaInP with an aluminum composition, x, of 60 to 100% and is doped with Si or Te to a level of n - 1 x 1OI8 ~ m - As ~ .we discussed in Subsection 2 of Section IV, the choice of Te as the n-type dopant can cause problems due to Te carryover effects in the MOVPE process. In a DH structure, this can lead to a Te doping tail that trails into the light-emitting active-layer portion of the device. However, both Si and Te are commonly used n-type dopants for AlGaInP LEDs. The active layer (AL) is composed of direct bandgap AlGaInP of x = 0 to -50% for red to green wavelengths. Fine control of the A1 to Ga ratio as well as the indium composition is critical to accurate targeting of the LED wavelength. A 1% change in the A1 to Ga ratio results in an 2-nm shift in the wavelength while a 1YOchange in the In composition results in an -5-nm shift in the wavelength. Accurate calibration of the TMAl, TMGa, and TMIn flow controllers must be

-

2 HIGH-EFFICIENCY AlGaInP LIGHT-EMITTING DIODES

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maintained to retain accurate lattice matching and control of the LED wavelength over the range of compositions used. As discussed in Subsection 2 Section IV, ordering in the AL also influences the emission wavelength and can either be controlled by using the appropriate growth conditions or eliminated by postgrowth annealing. After growth of the AL, the AlGaInP upper confining layer is grown with a composition of x = 60 to 100% doped with Mg or Zn to p-2-5 x 1017 cmP3. One important consideration for the growth of the DH is the choice of growth temperature. As we discussed earlier, the growth temperature of AlGaInP tends to be in the 650 to 850°C range. Higher growth temperatures are desired to reduce oxygen incorporation, but high growth temperatures can cause problems during the gas switching from GaAs or AlGaAs to the AlGaInP epilayers due to reevaporation of As (Chen, C. H., et al., 1997; Ota, 1993). An appropriate growth temperature must therefore be chosen to accommodate both issues. An alternative approach is to use multiple growth temperatures during the DH growth, where a low growth temperature is chosen for the AlGaInP growth initiation while a high growth temperature is used to grow the AlGaInP active layer. This temperature ramping, however, can introduce problems with reproducibility since the growth temperature must be well controlled at both the AlGaInP initiation and active layer sections. This choice is thus a trade-off between manufacturability and device performance. Another important consideration is the p-n junction location in the LED structure. Since the p-type dopants diffuse much more readily during growth than the n-type dopants, the p-n junction location may differ from its intended location. Typically, the DH is grown p-up to minimize the amount of time the p-type dopants are exposed to high growth temperature, thereby reducing the magnitude of this acceptor diffusion problem. The choice of growth temperature and p-type dopant can also influence the magnitude of this diffusion problem. Zn has a higher diffusion coefficient than Mg which may make Mg preferable as the p-type dopant (Chen, C . H., et al., 1997). The turn-on delay associated with Mg can also cause problems (Nishikawa et al., 1993), however, which may require the Mg source to be switched on prior to the desired position of the p-n junction. The n-type dopant can also affect the position of the p-n junction, as for example, in the case of a Tedoping tail into the active layer of the device. b.

Window Layers

A second critical part of AlGaInP LEDs is the window layer, which can more than double the efficiency of the device through improved current

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spreading and light extraction. As was discussed in Subsection 4 of Section 111, the window layer can be composed of various transparent materials, the most common being AlGaAs and Gap. In this section, we discuss the issues associated with the growth of these window layers. The growth of transparent AlGaAs window layers is a fairly straightforward process since MOVPE grown AlGaAs is inherently p-type and lattice matched to AlGaInP on GaAs, and the window layer can be grown in the same MOVPE growth run immediately following growth of the DH. Alternatively, the AlGaAs window layer can be grown in a second MOVPE run, and a two-step MOVPE growth can be used to facilitate the fabrication of current-blocking layers that are used to further enhance current spreading. The same p-type dopant is typically used for the AlGaAs window and the DH structure, and the common choice is Zn at a concentration of p-1-5 x 10” cm-, (Sugawara et al., 1991). One concern with AlGaAs windows is that high-aluminum-compositionAlGaAs can potentially contribute to poor reliability due to oxidation on exposure to even a low level of moisture. The reliability of AlGaInP LEDs with AlGaAs window layers can be excellent, however, if the exposed area of AlGaAs is minimized by capping the AlGaAs with a thin layer of GaAs or other nonhydrolyzing material (Sugawara et al., 1992~).An additional potential issue associated with this AlGaAs window layer is group V intermixing due to PH, to ASH, switching, as discussed in Subsection 2 of Section IV. Low growth pressure and temperature, combined with the use of a hydrogen-only purge step, can be used to reduce reevaporation of phosphorous at this interface (Hergeth et al., 1991). Unlike AlGaAs window layers, the fabrication of GaP window layers involves the growth of a heavily mismatched layer on top of the AlGaInP epilayers. The lattice mismatch does introduce misfit dislocations, but these dislocations remain confined to within a few hundred angstroms of the GaP to AlGaInP interface (Fletcher et al., 1991). The growth of the GaP window layer can be done by either MOVPE in the same growth run as the DH or in a second growth step using an MOVPE or HVPE process (Huang et al., 1992). The HVPE deposition can proceed at much higher growth rates of tens of micrometers per hour so that very thick (> 50 pm) window layers can be grown at low cost. If MOVPE is used for the GaP window layer, a thinner layer is typically used to maintain a short growth time and low cost. The HVPE approach, therefore, decreases maintenance and increases throughput of the MOVPE process since only the thin DH layers need to be deposited in the MOVPE step. The HVPE process uses gallium metal and PH, for the source materials, and unlike the MOVPE process, the HVPE process uses a HC1 prereaction with the gallium metal to form GaCl, which is transported to the wafer surface to react with PH,.

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The surface morphology of the mismatched GaP window layer on AlGaInP is heavily dependent on the GaP growth conditions, and Lin et al., (1994b) have demonstrated that an increase in the growth temperature from 760 to 800°C resulted in a decrease in the hillock density by over two orders of magnitude. An increase in the substrate misorientation toward [ O l l ] from 2 to 15" off the (100) surface similarly resulted in a decrease in the hillock density by over a factor of a hundred (Lin et al., 1994b). However, the exact growth conditions chosen for the GaP window layer can be limited by their effect on the underlying D H structure. For example, a high GaP window growth temperature can improve the surface morphology but can also cause additional p-type dopant diffusion in the underlying DH layers. The choice of GaP window growth conditions must accommodate both the requirements for a smooth surface morphology and optimal device performance.

c. Distributed Bragg Rejiectors As discussed in Subsection 5 of Section 111, distributed Bragg reflectors have been used in AlGaInP LEDs to increase device efficiency (Kato et al., 1991). These LED DBR layers are grown on the GaAs substrate prior to growth of the DH and are typically grown in the same MOVPE growth run as the DH. Unlike DBRs used for vertical-cavity surface-emitting lasers (VCSELs), DBRs used for LEDs require a wide reflectance bandwidth due to the broad spectral width and wide angular distribution of the emitted light (Saka et al., 1993). AlGaInP LEDs also typically have a shorter emission wavelength than VCSELs. These differences manifest themselves in a different set of DBR materials for AlGaInP LEDs. In this section, we discuss the material choices and growth requirements for these DBRs. One combination of materials for LED DBRs is AlGaAs-A1As (Saka et al., 1993; Lee, 1995; Chang et al., 1997b). This material system is commonly used for VCSELs due to the large available index contrast and lattice matching to GaAs. In this case, Si or Se is used for the n-type dopant with SiH, or H,Se as the source. The doping and composition profile in these DBRs is commonly graded to reduce the DBR series resistance. These DBRs can be made transparent using AlGaAs or lossy using GaAs for the highindex layer (Chang et al., 1997b). However, these materials are more chemically unstable than the AlGaInP DH and can potentially cause device reliability problems (Sugawara et al., 1993). Another approach is to use transparent AlInP- AlGaInP for the DBRs (Sugawara et al., 1992a). These materials are the same as those used in the DH and are thus grown in essentially the same manner as the DH. The

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disadvantage of this approach is that the refractive index difference between AlInP and transparent AlGaInP is fairly small, and a large number of DBR pairs are needed to provide the necessary reflectivity (Sugawara et al., 1993, 1994b). This introduces a thick layer of AlGaInP material that must be accurately lattice matched to ensure that the overlying DH layers are of high quality. AlInP- AlGaInP DBRs, therefore, require tight control of the growth as well as a significantly longer MOVPE growth time. One other combination that has proven to be very effective is a hybrid AlInP-GaAs DBR, which is a lossy type DBR (Sugawara et al., 1993,1994b). In this case, the refractive index contrast is high and fewer pairs are required. However, the growth becomes more difficult due to the need to switch between arsenide and phosphide materials. Sugawara et al. (1993) used a low pressure of 25 torr combined with a moderate growth temperature of 730°C to minimize group V intermixing at the interfaces. With this approach, they have demonstrated sharp interfaces with good electrical characteristics in LED devices. Thus, several materials combinations have been used for DBRs in AlGaInP LEDs, each introducing its own set of benefits and difficulties. 4.

MANUFACTURING ISSUES

With the entry of AlGaInP LEDs into large-scale markets such as indicators, displays, power signaling, and automotive lighting, high-volume manufacturing of AlGaInP LED epitaxial wafers has become necessary. MOVPE technology, although providing numerous benefits in precise control and purity, has much reduced growth rates and throughput compared to the more traditional growth techniques of HVPE and LPE. Highvolume manufacturing with MOVPE thus requires growth on multiple wafers simultaneously in a single large chamber. For this multiwafer growth, variation of LED performance within a single growth run, or uniformity, is a major issue. In addition, variation from run to run, or reproducibility, must also be minimal for a true manufacturing process. Finally, with these large MOVPE reactors, considerable quantities of chemical inputs and waste are used and generated, and handling of these chemicals and wastes in a safe manner, while minimizing impact on the environment is a very important part of the manufacturing process. This section addresses these three primary issues of high-volume MOVPE manufacturing of AlGaInP LEDs: environmental and safety issues, uniformity, and reproducibility. a.

Environmental and Safety Issues

MOVPE manufacturing of AlGaInP LEDs raises several environmental and safety concerns. Numerous hazardous chemicals are used in the process,

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and considerable quantities of hazardous wastes are generated as a byproduct. The chemical hazards associated with AlGaInP MOVPE include toxic phosphine and arsine gases, pyrophoric and toxic metalorganics, and pyrophoric phosphorous waste. These chemical hazards and their associated treatments are discussed in this section, and a summary of the chemicals and their handling and treatment is shown in Table 11. The group V sources used for AlGaInP MOVPE are usually the highly toxic phosphine and arsine gases. These gases require both proper handling and abatement. Storage of these gases is typically in ventilated gas cabinets in a remote gas bunker, and double-contained lines are used to deliver the gases to the MOVPE reactors. Extensive toxic gas monitoring systems at TABLE I1 MAINCHEMICAL HAZARDS ASSOCIATED WITH MOVPE MANUFACTURING OF ALGAINPLEDs AND COMMON PRECAUTIONS AND TREATMENTS Prim a ry hazard

Common precautions

Phosphine/ arsine

Highly toxic

Toxic gas monitors Remote storage Double contained lines Procedural- training

Combustion and oxidation to form phosphorous and arsenic oxides

Phosphorousarsenic oxides

Toxic

Immediate abatement

Particle scrubbers

Metalorganics

Pyrophoric

Flame detectors Specialized fire extinguishing materials Standardized installation fittings Procedural- training

Consumed and decomposed

Phosphorous solid

Pyrophoric

Flame detectors Procedural- training

Condensation into cold traps Oxidation or chemical conversion to stable form prior to disposal as solid toxic waste

Hydrogen

Combustible

Hydrogen monitoring

Combustion to form water vapor

Arsenic

Toxic

Respirator protection

Solid toxic waste disposal

Chemical

Disposal

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the tool and room level are used to ensure containment of these toxic gases and to automatically shutdown the gas systems in case of leakage. In addition, only some of the phosphine and arsine gases are decomposed or consumed in the MOVPE process and much of the phosphine and arsine gas becomes waste. This waste must be treated to minimize risk to the environment, and typically a combustion chamber is used to burn the phosphine and arsine waste to form stable phosphorous and arsenic oxides. These oxides are then captured using a fine particle scrubber and disposed of as toxic waste. An alternative method is to use carbon granules or another solid material to adsorb the hydride gases. These adsorbed gases are then periodically oxidized to convert them to solid oxides. This high level of attention to safety and environmental waste abatement is currently the standard in the industry and is an absolutely necessary part of such operations (Chen, C. H., et al., 1997; Thompson, 1994). The group 111 metalorganic sources are pyrophoric and also require special safety precautions. In addition to the standard smoke and heat detectors, infrared flame detectors are commonly used to monitor areas containing pyrophorics, and fire-extinguishing materials, designed specifically for pyrophoric and metal fires are kept on hand. Well-defined procedures and training are also necessary to minimize risk of misuse or mishandling of these pyrophorics. In addition, metalorganic users and suppliers have been developing a standardized polarization of the fittings on the metalorganic bubblers to prevent improper installation (Peanasky, 1998). With continued industrywide improvements, the safety hazards associated with the handling and use of metalorganics can be minimized. Solid phosphorous is a by-product of the AlGaInP MOVPE process and accumulates downstream from the reaction chamber. Phosphorous waste in a white phosphorous form is pyrophoric and requires special safety precautions. The reactor exhaust lines can be heated to minimize condensation of phosphorous at undesirable points, and cold traps are typically used to condense out and contain the majority of the phosphorous waste. Once this phosphorous is contained, it must be chemically treated and converted to a more stable form before disposal. One method of treating the phosphorous is by burning it in a controlled manner to form stable phosphorous oxides. An alternative method is to react the phosphorous with a chemical solution to form a stable solution. Once again, a well-refined procedure and training system must be developed to minimize safety risks associated with handling and treatment of these wastes. Thus, for manufacturing AlGaInP LEDs, systems must be well established to properly address all of these safety and environmental hazards.

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Uniformity

For high-volume manufacturing of LEDs, variation of material properties and device performance within a run, both across a wafer and from wafer to wafer must be minimized to maintain high process yields. The AlGaInP material composition, doping level, and layer thickness must all be very uniform to achieve uniform device performance. Achieving uniformity in all three components simultaneously can in itself be difficult, and the problem becomes even more complex in the manufacturing environment where multiple wafers are grown in a single run and wafer uniformity must be consistent over time. To solve this complex problem, temperature uniformity across the wafers must be excellent, and source delivery and depletion issues must be addressed. Extensive work has been done over the years to solve these problems and only some highlights on AlGaInP uniformity are touched on here. Various solutions have been developed to achieve uniform temperature. Most large MOVPE chambers used for AlGaInP growth have multiple heating zones, which can be individually tuned to provide a uniform temperature profile. With resistive heaters, two or three separate heater zones are commonly used, while lamp heated chambers typically use six or more heater zones to optimize the heating profile. For chambers that are rf heated, the rf coil spacing is usually adjustable to tune the temperature profile on the graphite susceptor and hence, the wafer. The wafer temperature uniformity can also be adjusted by modifying the thermal coupling of the susceptor to the wafer. These heating profile adjustments have been mostly successful in providing a uniform temperature across multiple wafers. For example, temperature variations of less than 1°C have been demonstrated across a three inch wafer in a multiwafer reactor with a two-zone resistive heater (Gurary et al., 1995). As for source delivery and depletion, reactor design optimization, combined with process refinement, has resulted in uniform delivery across the wafers in multiwafer reactors. In particular, two basic reactor geometries have been used to achieve a high level of uniformity across multiple wafers: vertical flow rotating-disk and horizontal flow planetary geometry. A planetary reactor is basically a horizontal tube reactor extended to two dimensions rotationally so that the reactor is effectively a horizontal tube reactor without side walls. As a result, the gases enter through the center of the chamber and flow outward over the susceptor and wafers. When this reactor geometry is combined with a rotating susceptor and rotating wafers, gas depletion effects can be averaged out and high uniformity can be achieved. In a rotating disk reactor, on the other hand, the source gases are vertically incident on a high-speed rotating

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disk on which the wafers reside. This fluid flow has been extensively studied and modeled and has been shown to provide a uniform boundary layer across the rotating disk (Breiland and Evans, 1991; Tompa et al., 1994). This reactor geometry has also been used to grow very uniform films across multiple wafers. In combination with reactor design optimization, certain process parameters have also been optimized to achieve high uniformity. Low growth pressure and high rotation speed have been shown to help improve thickness and compositional uniformity. These growth conditions minimize depletion of the metalorganic species across the growth zone by increasing the gas velocity (Motoda et ul., 1994). Excellent compositional, thickness, and doping uniformity have been achieved using an optimized process in these reactor designs. As a measure of compositional uniformity, wavelength uniformity of k0.35 nm has been demonstrated for GaInP across a 2-in. wafer in a multiwafer horizontal flow reactor (Chen, et al., 1997). This corresponds to only a 0.25% variation in indium composition across the wafer. Excellent AlGaInP compositional uniformity has also been demonstrated in a vertical flow reactor with a standard deviation of only 0.6 nm in PL wavelength over 17 2-in. wafers of AlGaInP single layers (Thompson et al., 1997). As for thickness, very uniform DBRs and AlGaInP layers have been demonstrated. Using a rotating disk reactor, DBR thickness variation of less than &0.2% has been demonstrated on 3-in. wafers (Hou et al., 1996). AlGaInP layer thickness uniformity below 5% on 3-in. wafers has also been achieved (Motoda et al., 1994; McKee et al., 1992). Doping uniformity results have been excellent in these reactors with a Si- and Zn-doping uniformity as good as + l o % across a 3-in. wafer (Motoda et a]., 1994). Thus, with the extensive work to optimize process conditions along with temperature and source delivery uniformity in multiwafer reactors, excellent uniformity results have been achieved that have made manufacturing of AlGaInP LEDs practical.

c. Process Control In addition to variation within a wafer and from wafer to wafer in a single run, variation among MOVPE runs, or reproducibility is also very critical for high-process yields. In particular, run-to-run reproducibility of wavelength, doping, thickness, and source purity must be controlled. Wavelength reproducibility depends primarily on the indium composition of the active layer. Stability of the trimethylindium delivery is a key component of wavelength reproducibility, and various in situ methods have been used to ensure consistent delivery of TMIn to the wafers. One common approach to controlling the TMIn delivery is to use an in situ acoustic

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concentration monitor that provides feedback to actively compensate for drift in the TMIn concentration by controlling the TMIn flow rate (Thrush et al., 1987). Since TMIn will decompose on hot surfaces, TMIn delivery is also influenced by variations in the temperature of surfaces in the reaction chamber. Active cooling and control of the chamber wall temperatures can help reduce variation in TMIn depletion at these surfaces. Control of the quartz liner temperature can similarly be used to help reduce run-to-run variation in the TMIn depletion. This can be achieved by using a pyrometer or a thermocouple in contact with the liner to accurately measure the liner temperature. With an accurate reading of the liner temperature, the mixture of H, and N, in the space between the liner and a cooled surface can be varied to control the liner temperature (Strauch et al., 1996). The composition of the gas mixture determines the thermal conductivity of the gas and hence the temperature of the quartz liner. E x situ monitoring of the indium composition is also usually used to further help maintain control of the wavelength. Due to the ordering phenomenon in AlGaInP, photoluminescence is typically ineffective for wavelength control since the degree of ordering will change with further processing (Chen, C. H., et al., 1997). Instead, x-ray diffraction can be used to directly measure the indium composition of the layers. When this x-ray diffraction sampling is combined with statistical process control (SPC) software to control the TMIn flow setpoint, the run-to-run wavelength variation can be very minimal. A sample trend chart showing run-to-run indium composition variation and the corresponding SPC adjustments to TMIn flow in a multiwafer MOVPE reactor is shown in Fig. 14. The run-to-run standard deviation in this case is only 0.27% indium composition corresponding to a run-to-run wavelength variation of less than 1.4 nm standard deviation. A similar approach can be used for doping control. For example, capacitance-voltage measurement of the surface doping concentration can be used for ex situ control of the window layer doping. Hence, the combination of in situ and ex situ control can be used to minimize run-to-run variation in wavelength or doping in AlGaInP LEDs. Thickness control is another element of reducing run-to-run variation in LED device performance. Thickness control is especially important when DBRs are used since the center wavelength of the DBR is directly related to the thickness, and variations in the DBR center wavelength can cause variations in the LED emission wavelength and light output. A statistical process control method, as described in the preceding paragraph, can be used to control the LED thickness. In this case, the thickness can be measured using reflectance, ellipsometry, or scanning electron microscopy. Another approach is to use in situ monitoring to control thickness. I n situ reflectance has been used to measure AlGaAs and AlGaInP growth rates

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(Breiland et al., 1997). This diagnostic has been used to effectively calibrate AlGaAs growth rates for DBR and VCSEL growth (Hou et al., 1996). The in situ measurement yields an accurate determination of the growth rate. By feeding back this reflectance measurement data to the reactor control system, the thickness can be actively controlled in situ. Thus, similar to wavelength control, thickness control can be achieved using a combination of ex situ and in situ control. Even if the growth process can be perfectly controlled using the methods already described, the reproducibility of LED performance further requires close control of incoming source material quality. In particular, the oxygen impurity level in phosphine, trimethylindium, and trimethylaluminum must be maintained at a low level to ensure proper LED performance. As described earlier in Subsection 2 of Section IV, much progress has been made in the synthesis of the metalorganics that has helped reduce the problems associated with oxygen in the source materials (Maeda et al., 1991; Roberts, et al., 1994a). In addition, point-of-use purification and screening of incoming source materials can be used to further ensure source purity. Thus, a combination of efforts by the user and supplier can be used to minimize source purity variations that affect device performance.

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V. Device Fabrication and Packaging 1. INTRODUCTION LED fabrication and packaging processes are typically straightforward and optimized for high yield. AlGaInP LEDs are constructed using metal contacts (patterned and unpatterned), wafer thinning, die-singulation, and final assembly into packages designed to increase the light output and shape the light into desirable emission patterns. The most common form of AlGaInP LED that uses an absorbing substrate of GaAs follows wafer fabrication steps, as described here.

AND CONTACTS 2. DIE DEFINITION

After the completion of epitaxial growth, suitable metals (e.g., Au and Zn or Au and Be) are deposited (evaporated or sputtered) onto the top surface of the LED wafer and are patterned (using photolithography and etching) into an array of circular (or circular with finger extensions) top contacts. The spacing of this array of dot contacts determines the final die size and hence the maximum yield of die per unit area of the wafer. The top ohmic contacts are usually alloyed after the pattern definition. Once the top contacts are defined and alloyed, their ohmic nature can be tested by forcing current between two adjacent contacts and checking for linear current-voltage (I-V) characteristics. The next step is wafer thinning, which may be done mechanically by lapping, or chemically by etching, or by a combination of the two. Etching to thin the wafers tends to smooth the surface of the GaAs substrate in preparation for the n-type back contact. A typical wafer thickness at this stage of the process is 9-12 mils. The thinned wafers typically then receive a blanket deposition of n-type contact metal (e.g., Ge and Au) that is also alloyed. The wafers are now ready for singulation into die. An important concern for the formation of ohmic contacts is series resistance. This causes the slope of the I-V curve to bow over at high forward bias and leads to decreased power conversion efficiency as discussed in Section VI. This series resistance is due to bulk resistance in the semiconductor and to contact resistance at the metal-semiconductor contacts. These effects can be reduced by using heavily doped semiconductor material and by minimizing the metal-semiconductor contact resistance. Of the two metal-semiconductor contacts, the p-type contact typically dominates the contact resistance of LEDs, since the top, p-type contact generally has a smaller area than the bottom n-type contact, and the hole mobility of 111-V materials is typically smaller than the electron mobility (Kish and Fletcher, 1997). The traditional ohmic contacts to p-type 111-V semiconduc-

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tors that rely on Au-Zn or Au-Be alloys yield specific contact resistances l o w 3to IR cm’. A relatively new class of Pd-based ohmic contacts allows a reduction in specific contact resistance by a factor of 10, however, and this reduced contact resistance can be particularly important for high power device applications. (Ha0 et al,, 1996; Lei and Jeng, 1987). The fabrication process for the transparent substrate AlGaInP differs from the above description in a few respects, as described in the following.

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3. WAFERBONDING

As discussed in Subsection 5 of Section 111, considerable improvement in AlGaInP LED performance can be achieved by providing a transparent substrate to replace the absorbing GaAs growth substrate. A technique for removal of this absorbing substrate is a required step in any TS device fabrication process. The implications for practical device fabrication are considerable. The AlGaInP DH layers must be supported and provided in a useful device structure once the substrate is removed. One method proven successful has been to remove the GaAs substrate after growth of a thick (-50 pm) GaP window on top of the AlGaInP DH layers (Kish et al., 1994). The thick C a P window acts as a supporting structure so that the device can be handled in a subsequent wafer bonding process. The AlGaInP-GaP DH and superstrate is brought into contact with a GaP substrate at high temperature and pressure. An electrically conductive bonded interface is formed between the AlGaInP layers and the G a P substrate, enabling the n-type GaP substrate to be used as the n-type contacting material. Because wafer bonding is a relatively quick process in which any reasonable thickness of GaP substrate can be provided, it has considerable advantages over a technique in which such a thick GaP layer were to be grown on the AlGaInP DH. The resulting all-transparent device has the benefit of providing high extraction efficiency, as discussed in Subsection 5 of Section I11 and illustrated in particular in Fig. 7d. The wafer-bonding process described above for fabricating transparent-substrate wafer-bonded devices was developed at Hewlett-Packard Company and has provided the world’s brightest commercially available AlCaInP LEDs. For further details, see Hofler et al. (1996) and references therein. 4. DICING

The dicing process results in the separation of the patterned wafers into die that are typically from 8 to 14 mils2. This process can be accomplished

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by sawing or by scribe and break methods, but the highest yielding process is the only practical choice. Scribe and break of AlGaInP LEDs, although favorable in terms of low kerf loss, is susceptible to edge chipping and incomplete die separation. Sawing, which can be accomplished at high speeds using a diamond-embedded wheel, is also sensitive to edge chipping if not performed properly, but can be optimized to give reasonable yields and throughputs. Kerf loss for sawing is typically -1.5 mils. The sawing process is typically performed on an expandable tape that is used to separate the chips for pick and place during high-speed die-attach processes. 5. PACKAGING ISSUES The dice are next assembled into a lamp or surface mount package, as shown in Fig. 15. The chip is attached to the leadframe of the lamp or the submount of the surface mount package using a Ag-filled die-attach epoxy. The assembly is accomplished by using a die-attach machine to place a small amount of die-attach epoxy in the leadframe, pick the die from the tape, and place the die into the die-attach epoxy. The die-attach epoxy is then cured, making electrical contact between the substrate and the leadframe and holding the chip in place for subsequent thermosonic wire bonding to the circular top contact pad. The wire-bonding step is quite important for reliable LED devices, as poor bond adhesion can result in the

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FIG. 15. Schematic of lamp and surface mount packages.

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loss of electrical continuity during encapsulation or during device operation in a temperature-cycling environment. When the wire bonding is completed, the whole assembly is encapsulated in a transparent epoxy of refractive index n - 1.5. The encapsulation epoxy is also important to reliable device operation since it protects the wire bond and the chip from damage. The epoxy has a much higher coefficient of thermal expansion than the AlGaInP die. This leads to the epoxy stressing the die during severe temperature cycling or during low-temperature operation. Care must be taken to minimize the effects of these stresses to provide accepiable LED performance over a wide temperature range. Another important parameter for the encapsulation epoxy is a resistance to “yellowing” (a loss of transparency) during exposure to ultraviolet (UV) light. This is especially important for AlGaInP LEDs as they are commonly used in outdoor applications such as variable message signs. Finally, the epoxy should slow the penetration of moisture into the package, as this may have deleterious effects during high-temperature soldering processes. Fortunately, the lower A1 mole fractions used in AlGaInP LEDs makes them less susceptible to moisture-related hydrolyzation degradation than that commonly seen in AlGaAs LEDs (Dallesasse et al., 1990). Therefore, AlGaInP LEDs are the device of choice for hightemperature, high-humidity operation in the red.

VI. AlGaInP Device Performance

Having discussed the growth, design, fabrication, and optimization of AlGaInP LEDs, the final topics of LED performance and degradation are discussed here. We begin by considering LED efficiency and currentvoltage characteristics, followed by a discussion of the wavelength dependence of AlGaInP device efficiency, and the color response of the human eye. Next, we address the issue of high-power AlGaInP LEDs, and finally, we discuss AlGaInP degradation issues.

1. LED EFFICIENCY AND CURRENTVOLTAGECHARACTERISTICS LED performance is typically described in terms of device efficiency, which can be defined in a number of ways, as described shortly. Although the external quantum efficiency and light extraction have been discussed in detail in Subsection 5 of Section 111, perhaps the most fundamental measure of LED performance is’ the internal quantum efficiency qint,

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which is related to the external quantum efficiency qeXtthrough the extraction coefficient, C,,,

In a direct bandgap material such as AlGaInP, the generation of light occurs through the recombination of electrons in the r-valley of the conduction band with holes in the valence band, possibly assisted by shallow donor or acceptor states located slightly below the conduction band or slightly above the valence band, respectively (see processes la, lb, and l c in Fig. 16). This direct recombination is characterized in terms of a radiative recombination lifetime z,. Alternative means by which electrons and holes can recombine without generating light include recombination through electronic defect states located deep within the bandgap of the active layer (process 2 in Fig. 16), leakage of minority carriers out of the device active layer (processes B or D in Fig. 16) followed by recombination within some other layer of the device, and shunt leakage or recombination of carriers along the perimeter of the device or along some extended defect, which shunts through the device active layer. As in the case of radiative recombination, nonradiative recombination may be characterized by a nonradiative recombination lifetime T ~ Leakage ~ . current and shunt leakage current are somewhat more

P active layer

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FIG. 16. Schematic illustration of the various recombination pathways in an LED. Processes A and C show electron and hole injection into the active layer, while processes B and D show electron and hole leakage out of the active layer and into the confining layers, respectively. Processes la, Ib, and Ic show radiative recombination pathways within the active layer, while process 2 shows nonradiative recombination within the active layer. Those carriers that leak out of the active layer (B and D) typically recombine nonradiatively in the confining layers through process 2.

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difficult to characterize, although we describe the effect of each of these terms on qint in the following. Including the effect of each of these processes, qintcan be expressed as qin, = (radiative recombination current/total current)

where I, is the radiative recombination current, I,, is the nonradiative is the recombination current, Ishunt is the shunt leakage current, and lleakage leakage current into the confining layers. Depending on the magnitude of the total current, various terms in Eq. (11) will dominate, while other terms may be ignored. A more detailed understanding of each of these terms can be developed through the assistance of Figs. 17a and 17b, which show a circuit diagram model of an LED and the semilogarithmic I-V curve of a typical AS AlGaInP LED. Note that the total current flow in the circuit of Fig. 17a is given by

where Vf = V - ZR, is the voltage drop across the ideal diode; R, is the series resistance; R,, is the shunt resistance; R,, is the equivalent resistance, I , is the saturation current of the ideal diffusion diode, and l R - Gis the saturation current of the nonideal recombination-generation diode, respectively. Since the bandgap of AlGaInP is 2 eV, very little current flows through the device at forward bias values significantly below 2 V. At such small forward bias values (Vf< 1 V), the majority of the current flows along parallel or shunt leakage paths such as die edges, or extended defects (dislocations, stacking faults, etc.) that short through the p - n junction. This effect is modeled by the shunt resistance path Rsh, in Fig. 17a. Low shunt leakage current is, of course, a sign of good material quality, and as Fig. 17b indicates, the leakage current of AlGaInP diodes is very small. Since leakage current does not contribute to light output, LED efficiency is very poor for such low forward bias values, as shown in Fig. 18. As the forward bias is increased, the injected current density begins to increase, and nonradiative recombination begins to dominate the I-V characteristics, as shown by the n = 2 region in Fig. 17b. This portion of the I-V curve is due to electron-hole recombination within the depletion region of the p - n junction and this recombination generally takes place through nonradiative recombination-generation (R-G) centers, such as oxygen, located in or near the device active layer (Suzuki et al., 1993; Sze, 1981).This n = 2 region is modeled by the R-G diode (labeled ZR-G) in Fig. 17a.

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(b) FIG. 17. (a) Equivalent circuit model for an LED: I represents the external current source; R,,, R,, and Re, are the shunt, series, and equivalent resistances, respectively; and I, and l R - G are the ideal and nonideal diodes, respectively. (b) I-V curve of an LED diode. Note the regions labeled n = 2, where nonradiative recombination is dominant; n = 1, where radiative recombination is dominant; the rollover at high currents due to series resistance; and the shunt leakage at low currents.

As the forward current is increased beyond the saturation current of the R-G diode, IR-G, the ideal diode begins to dominate the forward I-V characteristic, and the slope of the I-V curve approaches a value -q/kT. This portion of the I-V curve is referred to as the ideal, n = 1, portion. In this region of the I-V curve, Ishunt and Ileakage are both negligible, and the expression for qint reduces to

Assuming a p-type active layer, the radiative and nonradiative current can

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be expressed in terms of the injected minority carrier electron concentration, divided by the average time required for radiative (7,) or nonradiative (znr) recombination to occur, respectively (i.e,, current = charge/time); that is,

From Eq. (16) we see that rin,is maximized by having z, << rnr,and as z,/z, approaches 0, qint approaches 1. Since the radiative and nonradiative recombination processes occur in parallel, z, and z, cannot be determined independently and one measures the recombination lifetime z,defined in Eq. (17). For direct bandgap materials, z, is l.ypically much smaller than znr, and measured values of T are hence approximately equal to 5,:

Now, the radiative recombination lifetime z, is an intrinsic property of the material in question (AlGaInP in this case) and is independent of material quality. The nonradiative recornbination rate is an extrinsic property of the material, however, and can vary widely depending on material quality (i.e., oxygen concentration, etc.). Typical recombination lifetime values for highquality AlGaInP are 5-15 ns (Maranowski et al., 1997), although z can decrease significantly below 15 ns, depending on the value of z, which is

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influenced by oxygen concentration, dislocation density, etc. in or near the active layer. This extrinsic dependence of z, on material quality is expressed through the trap concentration N,, while the radiative recombination lifetime can be expressed in terms of a fundamental (material dependent) constant B, known as the radiative recombination coefficient. That is, z, = 1/Bp for p-type material, while ,z, = l/(cruN,) where CT is the capture cross section of the trap in question, and IJ is the thermal velocity of the carriers. Substituting these terms into Eq. (16), we obtain

Note that qint increases with increasing active layer doping and decreases with increasing trap concentration. Also note from Eq. (18) that materials with a larger radiative recombination coefficient B are more immune to or more tolerant of defects than materials with a smaller radiative recombination coefficient (Grillot et a!., 1996). This may offer some insight into the higher dislocation density tolerance for phosphide-based LEDs relative to arsenide-based LEDs. Equations (16) and (18) describe qint in the regime where leakage current is small. As the forward current increases beyond the saturation current of the ideal diode I,, some carriers begin to spill over into the confining layers, and the LED efficiency decreases for a variety of reasons. First, any carriers that leak into the confining layers will recombine nonradiatively in these indirect bandgap materials and will not generate light. Second, as the forward current increases, this nonradiative recombination can lead to substantial internal heat generation, which decreases efficiency, as discussed in Subsection 4 of Section VI, with regard to To. Third, with increased forward bias, series resistance R,, begins to play a more substantial role, as indicated by the decreasing slope in Fig. 17b at large forward bias. Most of the effects described here (i.e., shunt leakage, nonradiative recombination, and forward leakage current) reduce device efficiency by reducing the number of minority carriers (electrons) that are available for radiative recombination in the device active layer. Series resistance affects LED efficiency in a slightly different manor, which can be seen through the power conversion or wallplug efficiency qWp:

where C,, is the extraction coefficient, Ephis the photon energy in electron volts, and qV’ is the forward voltage drop across the device in electron volts. Note that the series resistance will act to increase V’, and hence decrease qwp.Typical series resistance values are 1.5 to 3 mR/cm2 for AS AlGaInP

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devices and -5 to 7 mi2/cm2 for TS AlGaInP. The additional series resistance of TS AIGaInP devices is primarily due to the higher resistivity of the GaP substrate as compared to GaAs (Kish and Fletcher, 1997). 2. EFFICIENCY vs WAVELENGTH The wavelength dependence of quantum efficiency in AlGaInP LEDs is strong and is attributed mainly to fundamental physical mechanisms within the material system, as described previously. The most dominant mechanisms are nonradiative recombination associated with the indirect bandgap minima and with impurity concentration, both of which increase with increasing A1 content in the AlGaInP material system (Cao et al., 1990, 1991). At x = 0 (A-650 nm), this competitive nonradiative recombination is at a minimum, and hence the superior quantum efficiency performance of red AlGaInP LEDs as compared to amber AlGaInP LEDs, where x = 0.3 (1-590 nm). Another main contributor to reduced active layer quantum efficiency with increasing A1 content is reduced carrier confinement. Electron overflow from the active layer to the indirect bandgap confining layer effectively leads to increased total nonradiative recombination, since, as the A1 composition of the active layer is increased, the potential barrier height for retaining electrons in the active region is reduced. This leakage current becomes a significant loss mechanism that causes the internal quantum efficiency of amber LEDs to be further reduced relative to red LEDs. The reduced confinement for amber LEDs also leads to increased dependence of light output on temperature, as discussed in Subsection 4 of Section VI. Figure 19 shows the best measured external quantum efficiencies of TS AlGaInP LED lamps as a function of peak wavelength at T = 300 K and a current density of 110 A/cm2 (corresponding to 50 mA in a standard 216 x 216pm2 chip). Peak efficiencies are obtained in the longer wavelength, red color regimes where the internal quantum efficiency is highest. The highest external quantum efficiency vex, is 29.9%, suggesting that almost one-third of the electrical current that flows through this device is converted to useful light output. This efficiency approaches the limit imposed by extraction efficiency according to the simple cone model for light extraction discussed Subsection 5 of Section 111, suggesting the internal quantum efficiency in this device is quite high. The fall-off of efficiency toward shorter wavelengths is indicative of reduced internal quantum efficiency. Extraction efficiency is also penalized by decreasing internal quantum efficiency however, since any photons that are reabsorbed within the active region are less likely to be reemitted in a material with lower internal quantum efficiency. The effects combine to reduce the external

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quantum efficiencies of shorter wavelength AlGaInP devices and are the reason for the shape of the curve in Fig. 19. 3. LUMINOUS EFFICIENCY

Having established the radiometric performance available from the AlGaInP material system, we must now consider the most important factor governing LED performance in indicator, power-signaling, and illumination applications: the human eye. The luminous performance of LEDs are governed by the photosensitive response of the human eye to the wavelength of photons emitted by the LED. The normalized luminous efficacy for photopic (daylight-adapted) vision determined by CIE is denoted VA), which has a value of unity at 555 nm, corresponding to the eye’s peak sensitivity (see textbooks or in this series Hodapp, 1997). At that wavelength, 1 W of optical power is defined to be equivalent to 683 lm. Thus, the luminous flux @(A) (in lumens) of an LED can be written as

where P r a d is the radiometric power emitted from the LED in watts if the emission can be regarded as narrow in comparison with the gradient of V(d) (otherwise the product has to be replaced by a convolution integral).

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The emitted radiometric power can be written as

where qext= qin,Cexis the external quantum efficiency, qint is the internal quantum efficiency, C, is the extraction efficiency, E,, is the photon energy in electron volts, and I, is the forward current. Combining these terms we can write an expression for the luminous power efficiency (in Im/W), as

where P = I,V, is the power input to the LED, and V, is the forward voltage across the device. Note that the series resistance discussed in Subsection 1 of Section VI reduces qL through increased V, in Eq. (22). In direct contrast to the physical mechanisms that reduce internal quantum efficiency in shorter wavelength AlGaInP LEDs, the slope of the CIE photopic sensitivity curve in this wavelength region makes these shorter wavelength LEDs appear brighter to the eye, as depicted in Fig. 20. The effect of the CIE curve is to partially compensate for the wavelength dependence of radiometric efficiency for AlGaInP, giving a flatter response for luminous efficiency from 590 to 630 nm. Figure 20 shows the luminous efficiency in lumens per ampere ( q L .yf) of the devices previously represented in Fig. 19.

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550

570

590 610 Peak Wavelength (nm)

630

650

FIG. 20. Best measured luminous efficiencies in lumens per ampere of TS AlGaInP LED lamps. The dashed line is a guide for the eye.

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Electroluminescent spectra from TS AlGaInP LEDs at 110 A/cm2 are shown in Fig. 21. Typical full width at half maximum values for conventional LED structures are 12 to 20 nm, with broader spectra occurring for the longer wavelengths (Fletcher et al., 1991). The slope of the CIE curve yields a weighted spectra as seen by the human eye, which tends to appear blue-shifted relative to the true radiometric spectra as given in Fig. 21. The perceived peak wavelength is called the dominant wavelength A, and is defined by the CIE color coordinates of the LED emission (Hodapp, 1997).

554 m

500 520 540 580

580 8M 620 Wavelength (nm)

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As discussed in the next section, the quantum efficiency of AlGaInP devices is sensitive to temperature, as specified by the radiometric characteristic temperature. This is especially true for the shorter wavelength devices. In addition, bandgap shrinkage with increasing temperature results in a red shift of the emission spectra, as illustrated in Fig. 22. Because the wavelength regime of operation for AlGaInP (560 to 650 nm) is along the downward slope of the CIE curve, this red shift always serves to reduce the luminous efficiency of these devices.

4. THERMAL MANAGEMENT Light-emitting diodes are, of course, just one of many components that go into an optical system (e.g., lighting fixture, signal, display pixel). In the course of designing the optical system to meet its performance requirements (which are typically radiation pattern, color, and efficiency), thermal management is critical because temperature of the LED’S active region affects both the performance (color, power dissipation, and efficiency) and the reliability of LEDs. AIGaInP-based LEDs are similar to other material technologies (AIGaAs, GaAsP, etc.) in the way that temperature affects performance. Specifically, at higher temperatures, for a fixed drive current, the forward voltage (V’) of the LED decreases, the electroluminescence

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spectrum shifts to longer wavelengths, and the internal quantum efficiency decreases. In most applications, the quantum efficiency change is the most significant, requiring the system designer to include a sufficient number of LEDs in the system such that the minimum flux or intensity requirements are met under the worst-case conditions of p - n junction temperature. For some applications, such as turn signals for automobiles, the acceptable color range for the LEDs is small enough, and the ambient temperature range large enough, that the color shift is a limiting factor in the design. Another important device parameter that is tied to the electron confinement in the active layer is the characteristic temperature T,, which can be defined for LEDs as P ( 7 J E Pr,,(25"C)exp

[- ( r j

--025"C)

(23)

where T j and Pradrefer to the junction temperature and the radiometric light output power (in watts) of the LED, respectively. A large To value indicates a reduced temperature sensitivity of the light output (LOP), which is highly desirable in most applications. As we have seen from Eq. (7), elevated temperatures tend to increase the electron leakage current and reduce device efficiency. Thus, the To value may be a useful measure of electron confinement for AlGaInP LEDs. The measured To of AlGaInP LEDs is sensitive to the bandgap difference, AEg, between the confining layers and active layer. This is illustrated in Fig. 23, where we plot radiometric To vs active layer composition for a large sampling of AlGaInP LEDs. As expected, the red LEDs (active layer A1 composition -0.07), which have better electron confinement, exhibit the highest To values, while the amber (active layer A1 composition -0.3) LEDs have a lower To. Although the lower internal quantum efficiency of amber AlGaInP LEDs gives a lower radiometric To relative to red AlGaInP LEDs, this effect is somewhat compensated by the increased sensitivity of the human eye to amber color as compared to red. Temperature-dependent color shifts thus have a less significant effect on the perceived brightness of amber LEDs than on red LEDs, and the net effect is a decreased wavelength dependence as measured by the photometric TXh, which is defined according to the following relation ( D ( T , ) = (D(25"C)exp[-(Tj

- 25"C)/Tgh]

(24)

-

where (D is the luminous flux and 7;. is the junction temperature. For red devices (> 615 nm), TJh is typically 150"C, while Tth is typically 100°C for amber devices (- 590 nm). This temperature dependence of luminous

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.. s

. .. . .

0 .

100

-

FIG. 23. Radiometric To vs A1 composition for AlGaInP LEDs.

efficiency can be significant in many applications. For example, using these numbers for a system with a thermal impedance of 100°C/W in an ambient temperature of 50°C, an amber LED driven at 50 mA and 2.0 V suffers an 30% reduction in light output over room-temperature junction conditions. A red emitter suffers only a 20% penalty under the same conditions (Gardner et al., 1999). The junction temperature q, and the photometric quantity Tgh, described in the preceding, will determine the light output of a particular LED. The junction temperature is a function of the ambient temperature (TJ, forward voltage (V’), forward current (If), thermal resistance (Rth), and LED internal quantum efficiency (vin,), as described by

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Techniques have been developed to extrapolate this value based on Vf measurements of devices that are heated by a train of current pulses (Hewlett-Packard Co., 1996a). Figure 24 illustrates the light output of an AlGaInP LED at thermal equilibrium with its environment (‘hot’), divided by the light output at room temperature (25”C, referred to as “cold’) immediately after the LED is turned on, calculated using Eq. (25). The power dissipated in the LED chip, coupled with the thermal resistance between the chip and the ambient environment, results in self-heating of the p-n junction. As the junction temperature increases, the light output of the LED decreases. Typical TZh values of AlGaInP LEDs are used in the calculations as well as two typical drive conditions for AlGaInP LEDs. The horizontal axis of the figure is the junction-to-ambient thermal resistance, which is partly under the control of the system designer. A system designer can typically reduce thermal resistance by using metal-core printed circuit boards instead of all-epoxy boards, using wide metal traces on the circuit boards, including heat sinks or fins or ventilation 1 .o

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fans, and placing the system in a housing that reflects sunlight (this is especially true for traffic signals). The LED package is also a significant contribution to the thermal resistance. Typical T 1-3/4lamps have a thermal resistance of 250°C/W (measured between the junction and one of the pins on the package, immediately outside of the epoxy lens); this can be reduced by a factor of 2 by using heavier gauge metal for the leads (Hewlett-Packard Co., 1996b). Changing the material (copper versus steel) may also reduce thermal resistance. For automotive applications, which have high flux requirements and wide ambient temperature ranges, Hewlett-Packard Company introduced a "Super Flux" package for AlGaInP LEDs that has an especially heavy leadframe, resulting in a junction-to-pin thermal resistance of -150°C/W. For outdoor applications, a new package for large-area LEDs has been introduced that has a thermal resistance of -12"C/W (Hofler et al., 1998). Packaging advances such as these reduce the burden on the system designer and reduce the cost of producing high-flux optical systems with LEDs.

5. HIGH-POWER ALGAINPLEDs

As the efficiency of LEDs has increased, penetration into markets traditionally dominated by conventional incandescent lamps has been realized. This is made possible when the total required luminous flux in an application is achievable with a reasonable number of LEDs from a system piece-part and assembly cost point of view. Obviously, the higher the LED efficiency, the fewer the number of LEDs required and the less expensive the system becomes. Also, in applications where LEDs outperform conventional sources in terms of power efficiency, considerable savings in terms of life cost of ownership are achieved because of reduced energy consumption. This is especially true for red traffic signals, for example, where an AlGaInP red LED lamp ( 20 lm/W or 11% W/W) is more than five times as efficient as a red-filtered incandescent lamp (- 3 lm/W or 2% W/W). Since the cost of an LED lamp is composed of both the chip cost and package cost, further reduction in system cost is achievable by increasing the chip size to provide more flux per package. Hewlett Packard recently introduced a new line of power TS AlGaInP devices capable of delivering more than 10 Im per element (about five times more than conventional LED lamps). These high-power chips have a lateral junction width w j of -500pm, corresponding to a die size -5.5 x that of conventional smalljunction devices, where w j -216 p m (Hofler et a)., 1998). These lamps enable traffic signals to be made with as few as 18 elements, compared with conventional lamps in which at least 100 or more LEDs are required, N

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and present a good value proposition to the traffic signal manufacturer. Important in the development of such power LEDs is maintaining high power efficiency even at these increased light output levels. Because the input power to these lamps can be quite significant (- 3/4 W or more), the thermal resistance of the package must be made as low as possible to avoid excessive heating of the junction and subsequently decreasing device efficiency, as discussed in detail in the preceding section. Furthermore, the top contact geometry of the chip must be designed to spread current uniformly across the larger junction area. Also, an extraction efficiency as close as possible to that achieved with conventional small-junction TS AlGaInP devices must be maintained. In fact, scaling up the junction area leads to a relative decrease in extraction efficiency caused by “clipping” of the side escape cones for this lower aspect ratio chip. This side light, which cannot escape on first or second pass, is mostly absorbed by one or more of the parasitic loss mechanisms inside the chip. The efficiency derating can be observed theoretically by a number of different approaches, including the use of a photon gas model as well as Monte Carlo ray tracing (Boroditsky and Yablonovitch, 1997; Tan, 1997). Figure 25 shows the relative extraction efficiency of transparent substrate LED chips (in epoxy) as a function of junction width. For scaling from wj-216 to -508 pm, a 20-25% efficiency derating is observed. Experimental measurements of external quantum efficiency are provided and show good agreement with the theoretical prediction.

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500x500 pm2 chip FIG. 26. Schematic of high-power lamp including a large-area TS AlGaInP chip.

Figure 26 shows a schematic of the high-power lamp package introduced by Hewlett-Packard for the large-area chip. The body of the lamp is a copper slug that acts as a heat sink. Thermal contact to the LED chip is made via a low-thermal-resistance submount to which the chip is soldered. The junction-to-case thermal resistance for this device is typically 12”C/W. A deep reflector cup is provided for high optical efficiency along with a soft gel inner encapsulant to minimize stress on the die during thermal cycling. A typical light output vs current curve for such a device is provided in Fig. 27 along with that of conventional 5-mm and “Super Flux” lamps, all emitting in the 630-nm range. The operating current for the large-area device is 250 to 350 mA as compared to the conventional devices driven at 50 to 70 mA. Over 15 lm of flux are provided at 350 mA in the high-power lamp.

6. ALGAINPDEGRADATION

As an alternative lighting source, LED lifetimes far exceed those of conventional incandescent or florescent bulbs with estimated lifetimes -100,000h or more. Catastrophic failure of an LED is an extremely rare event and when catastrophic failure does occur, it is more commonly associated with packaging issues than with the LED itself. While such catastrophic failure is rare, the light output of LEDs does typically change over time, and LED lifetimes are thus typically described in terms of the time required for the light output to reach some specified fraction (i.e., 50%),

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25,

FIG. 27. Light output vs current for high-power compared to conventional TS AlGaInP LED lamps.

of its initial value. In the following section, we compare the reliability of AlGaInP LEDs to that of AlGaAs LEDs. The reliability of AlGaInP LEDs under a variety of forward drive current conditions is also considered.

a. Comparison to AlGaAs As mentioned in the introduction, the earliest DH LED devices were based on AlGaAs-emitter technology. While these AlGaAs devices increased light output by roughly an order of magnitude relative to homojunction GaP and GaAsP LEDs, these AlGaAs LEDs were highly susceptible to degradation from a variety of sources. First, the high aluminum content of the DH confining layers made these devices very susceptible to hydrolyzation, where an opaque oxide layer would form on the exposed AlGaAs layers. This oxide would absorb a significant fraction of the light generated within the LED, and on exposure to high-humidity and high-temperature conditions, the extraction efficiency of these AlGaAs LEDs would thus decrease tremendously over time. With the increased area of exposed AIGaAs, this problem was much more severe in TS AlGaAs than in AS

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Stress time (hrs) FIG. 28. Effect of 20-mA forward stress at 85°C and 8 5 % relative humidity on the light output of AS AlGaAs, TS AlGaAs, AS AIGalnP, and TS AlGaInP. Note that AlGaInP is much less sensitive to the high-humidity stress than AlGaAs.

AlGaAs. Figure 28 compares the reliability of AS and TS AlGaAs to that of a new deep red AlGaInP part under high-temperature (85"C), high humidity (85% relative humidity) operation. This new AlGaInP LED is being developed by Hewlett Packard to replace TS AIGaAs. Note that this new AlGaInP part, with an emission wavelength in the 636- to 640-nm region, has far superior reliability to AlGaAs devices due to the lower A1 mole fraction used in these AlGaInP LEDs (Grillot, et al., 1999). While this severe degradation under high-humidity conditions significantly hindered acceptance of AlGaAs LEDs for emerging markets such as traffic signals, AlGaInP LEDs do not suffer from this problem and are currently gaining wide acceptance into these new markets. An additional advantage of AlGaInP LEDs relative to AlGaAs LEDs is found with regard to dark line and dark spot defects, which are commonly observed in AlGaAs and 11-VI emitter technologies, but have not been reported in AlGaInP LEDs. In fact, phosphide-based materials are generally less susceptible, in general, to such defect formation than arsenide based LEDs (Lester et a/., 1995), possibly through the higher activation energies required for dislocation motion in these materials.

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Current Dependence

b.

As discussed in the preceding section, AlGaInP LEDs are substantially more reliable than AlGaAs LEDs under high-humidity operation, and dark spot or dark line defects are not observed in degraded AlGaInP LEDs in contrast to AlGaAs and IILVI emitter technologies. While little is known about the detailed degradation mechanisms in AlGaInP LEDs, some insight into AlGaInP degradation can be gained from Fig. 29, which shows 12,000-h reliability data of 216 x 216 pm and 173 x 173 pm AlGaInP die operated under various DC forward current densities at 55°C. Note first of all, that the majority of the light output degradation occurs in the first 500 to 1000 h of stress time, after which the light output remains almost constant. Second, more severe degradation is observed and the time required to reach this maximum degradation generally decreases as the current density is increased. That is, at 180 A/cm2, the light output of the small die is basically independent of stress time for any time beyond 500 h. In contrast, the large die stressed at 65 A/cm2, takes 1000 h to reach this quasi-time-independent regime. A third important point from Fig. 29 is that the reliability of the two die sizes is virtually identical when they are both stressed and tested at 110 A/cm2. This indicates that light output degradation of AlGaInP LEDs is a bulk effect, and is not due to surface states or

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surface leakage current. Finally we note that the current densities given in Fig. 29 (- 110 A/cm2) correspond to the maximum recommended drive current densities for AlGaInP LEDs, and as Fig. 29 illustrates, excellent reliability is observed under these drive conditions. The preceding discussion illustrates the important effect of current density on degradation of AlGaInP LEDs. An additional consideration regarding the effect of current density on LED reliability must be considered for many of the new emerging LED markets. Traditional LED applications, such as indicator signals and indoor signs, generally require only a single, constant flux while the LED is in the on state. As outdoor applications of light emitting diodes become more prolific, the demands made on LED reliability become more and more stringent. For example, automobile rear combination lamps, which employ LEDs for turn signal, stop light, and tail light applications obviously require more flux for brake lights than for tail lights. Similarly, outdoor message signs may require higher flux during the day than they require at night. Full-color displays, which may require arbitrary combinations of red, green, and blue light to achieve the full range of colors, will also require variable LED flux. The method used to achieve this time-variable flux is critical to the reliability of all LED technologies, with the possible exception of AlGaInN. Recall from Subsection 1 in Section VI and Eq. (18) that the internal quantum efficiency of an LED is limited by the concentration of traps N , in or near the device active layer, and that LEDs are particularly susceptible to defect-mediated recombination at low current densities. If any generic arsenide- or phosphide-based LED is stressed at high forward current during daytime operation, defects may be generated in or near the device active layer, which could degrade LEI) performance. While the effect of these defects may be negligible at high current densities, their effect at low current densities may be quite dramatic. For such reasons, it is essential to carefully evaluate the chosen method of intentional dimming of LEDs to adjust for changes in ambient lighting conditions. For AlGaInP LEDs, it is recommended that such “day-night’’ dimming be accomplished through pulse mode operation, where the duty factor is changed to adjust LED light output. In the event that peak current must be used to adjust the flux of AlGaInP LEDs, the maximum ratio of peak current densities for high-flux to low-flux operation should not exceed 5 to 1 (Grillot et al., 1998). With regard to further details of AlGaInP degradation, very little has been published in the literature. Deep-level transient spectroscopy studies of InGaP LEDs have revealed the formation of electron traps that increase in concentration with device operation. The authors discuss the observed defect formation in terms of recombination-enhanced defect reactions and suggest that the observed electron traps are due to the formation of point

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defect complexes (Kim et al., 1993). A study of AlGaInP laser diodes under high-temperature and high-bias operation revealed an increase in threshold current density and a decrease in emission wavelength (Endo et al., 1994). Endo and coworkers attributed this decrease in emission wavelength to disordering in the active layer with device operation, while they attributed the increase in threshold current to some type of degradation in the active layer, possibly associated with recombination-enhanced processes. The increased threshold current was found to be thermally activated with an activation energy of 0.8 & 0.05 eV. Extrapolated lifetimes for these laser diodes were estimated to be lo6 h. In a more recent study, Herrick and Petroff (1998) showed that the initial performance and reliability of AlGaInP vertical-cavity surface-emitting lasers could be improved by annealing the devices with a high-forward-bias current. The observed improvement in performance was attributed to annealing-out of intrinsic native point defects.

-

VII. Conclusions

As we have seen throughout this chapter, AlGaInP LEDs enable the highest efficiency light output of any LED technology in the red to yellow portion of the visible spectrum, and these LEDs are approaching or exceeding the luminous performance of more traditional lighting sources such as fluorescent and incandescent light bulbs. We have also seen that the reliability of AlGaInP LEDs is excellent provided that appropriate drive currents are used and thermal resistance issues are addressed. Such increased efficiency and excellent reliability compared to more mature LED technologies such as AlGaAs and GaAsP have enabled AlGaInP LEDs to penetrate lighting markets traditionally reserved for incandescent lighting, such as traffic signals, automobile lighting, and variable message displays. Further improvements in AlGaInP efficiency will continue to open up new markets such as overhead street lighting and general illumination applications. Moreover, recent advances in GaN technology have provided efficient blue and green emitters, which make full color displays possible, and the widespread availability of GaN LEDs is certain to further proliferate the use of AlGaInP LEDs.

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SEMICONDUCTORS AND SEMIMETALS. VOL. 64

CHAPTER 3

High-Brightness Nitride-Based Visible-Light-Emitting Diodes R . S. Kern. W. Gotz. C. H . Chen. H . Liu. R . A4. Fletcher. and C. P . Kuo OPTOELECTRONICS DIVISION HEWLETI-PACKARD COMPANY SANJOSE. CALIFORNIA

I . INTRODUCTION. . . . . . . . . . . . . . . . . . . . . . . . . . . . 1. Historical Background . . . . . . . . . . . . . . . . . . . . . . . 2 . Applications and Markets for Gallium Indium Nitride Light-Emitting Diodes . . . . . . . . . . . . . . . . . . . . . . . I1. 111-V NITRIDE CRYSTAL GROWTH . . . . . . . . . . . . . . . . . . 1 . Nitride OMVPE Sources . . . . . . . . . . . . . . . . . . . . . . 2. Nitride OMVPE Reactors . . . . . . . . . . . . . . . . . . . . . . 3 . Substrates for Nitrides . . . . . . . . . . . . . . . . . . . . . . . 4. Gallium Nitride Growth on Mismatched Substrates . . . . . . . . . . 5. Gallium Nitride p- and n-Type Doping . . . . . . . . . . . . . . . 6. Alloy Growth . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 . Selective Area Epitaxy . . . . . . . . . . . . . . . . . . . . . . . 111. ELECTRONIC PROPERTIES AND CHARACTERIZATION OF 111-V NITRIDE FILMS. . . . . . . . . . . . . . . . . . . . . . . . . . . . 1 . Structural Properties . . . . . . . . . . . . . . . . . . . . . . . . 2. Luminescence Properties . . . . . . . . . . . . . . . . . . . . . . 3 . Properties and Hall-Effect Characterization of Doped Ill- V Nitrides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4. Gallium Indium Nitride . . . . . . . . . . . . . . . . . . . . . . . 5 . Aluminum Gallium Nitride . . . . . . . . . . . . . . . . . . . . . IV. NITRIDE LED DEVICE STRUCTURES AND PERFORMANCE . . . . . . . . . 1. LED Structures . . . . . . . . . . . . . . . . . . . . . . . . . . 2. Substrate Selection . . . . . . . . . . . . . . . . . . . . . . . . . 3. Device Design Considerations for n- Type Layers . . . . . . . . . . . 4. Device Design Considerations for p-Type Layers . . . . . . . . . . . 5 . Design of Active Layers . . . . . . . . . . . . . . . . . . . . . . . 6. Electrostatic Discharge Sensitivity and Reliability Characteristics . . . . 7 . WhiteLEDs . . . . . . . . . . . . . . . . . . . . . . . . . . . V. SUMMARY AND CONCLUSIONS . . . . . . . . . . . . . . . . . . . . . . REFERENCES . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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I. Introduction The present development of the 111-V nitrides has allowed them to progress very rapidly from basic research and development to large-scale commercialization. Gallium nitride (GaN), aluminum nitride (AlN), and indium nitride (InN) as well as their alloys, aluminum gallium nitride (Al,Ga, -xN or AlGaN), gallium indium nitride (Ga,In, -xNor GaInN), -,N or AlGaInN), are and aluminum gallium indium nitride (Al,Ga,In, -, compound semiconductors that, when crystallized in their wurtzitic polytypes, have direct energy bandgaps of 6.2 eV for AlN, 3.4 eV for GaN, and 1.95 eV for InN at room temperature (Fig. 1). This makes the IIILV nitride materials system of great importance for high-power electronics, shortwavelength optical emitters, and detector applications. At this writing, commercial production of high-performance ultraviolet, blue, blue-green, green, and white light-emitting diodes (LEDs) and violet laser diodes (LDs) has been established, and the development of electronic devices is underway. The progress that has driven the 111-V nitrides to commercialization has primarily been the result of improvements in crystal growth, epilayer characterization, device design, and device processing. Although many fundamental questions associated with 111-V nitride materials and devices

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currently remain unanswered, this chapter focuses on the current state-ofthe-art materials and LED developments as well as proposing some of the directions and issues for future pursuit.

1. HISTORICAL BACKGROUND Single-crystal III-V nitrides are not easily grown in bulk like GaAs and Gap. Consequently, no commercial nitride substrates exist today for the epitaxial growth of AlGaInN-based devices. Most thin films and devices are therefore grown epitaxially on sapphire (A1,03) or silicon carbide (Sic) substrates. The large difference in lattice constants and thermal expansion coefficients between the sapphire or silicon carbide substrate and the AlGaInN-based semiconductor has made it difficult to grow high quality nitride films on these substrates. Furthermore, highly conductive p-type nitride alloys have historically been very difficult to produce due to the high n-type background concentration and minimal p-type dopant activation. These difficulties hindered the development of efficient heterojunction lasers and LEDs using the nitride family of semiconductors, even though the basic heterojunction device requirements have been well understood for many years. During the late 1980s, the first highly efficient nitride-based LEDs were realized when researchers discovered a method of producing high-quality nitride epilayers on sapphire substrates by inserting an initial AlN (Amano et al., 1986) or GaN (Nakamura, 1991a) buffer layer at low temperature. Growing this buffer layer on a sapphire substrate greatly improves the morphology of the subsequently grown GaN layer and reduces the n-type background concentration. This growth process, coupled with the discovery of a postgrowth low-energy electron-beam irradiation (LEEBI) (Amano et al., 1989) or thermal anneal (Nakamura et al., 1992d) to activate the p-type dopant, has made growing conductive p-type GaN possible. With the aid of these technological advances, progress in device development of AlGaInN semiconductors has greatly accelerated. Perhaps the most notable of these devices is the AlGaInN-based violet diode laser. This device has an important role to play toward greatly increasing storage density in optical disk applications. Even so, the use of AlGaInN for producing high-efficiency LEDs has a much broader role to play in numerous commercial applications. Currently, several major suppliers exist for nitride-based LEDs including Nichia Chemical Industries, Ltd., and Toyoda Gosei Company, Ltd., in Japan, Cree Research, Inc., and Hewlett Packard Company in the United States, and Siemens AG in Germany.

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2. APPLICATIONSAND MARKETS FOR GALLIUM INDIUM NITRIDE LIGHT-EMITTING DIODES There are numerous immediate applications for the long-awaited shortwavelength LEDs. These applications have been hindered by the lack of blue and green devices with light output performance comparable to the yellow to red AlGaInP and AlGaAs devices that have been available for at least 10 yr. Examples of these applications include large-screen full-color video displays, red-amber-green traffic signal lights, and interior and exterior automotive lighting. It is now possible to implement a complete lighting system for the rear of an automobile where the turn-tail-stop lights and even the white backup lights are LEDs. Highly efficient LEDs possess inany beneficial characteristics relative to incandescent filament bulbs; namely, low power consumption and long lifetime. In the area of traffic signaling, for example, replacement of only the red incandescent lamps in the signal with LEDs will reduce the power consumption by 50%. In addition, the maintenance costs associated with bulb replacement are greatly reduced, since an LED traffic signal head is expected to last 5 to 10 yr. (The incandescent bulb for a red traffic signal is typically replaced at least once per year.) Similar advantages hold for LEDs used in message signs and large video displays. One of the most exciting potential uses for LEDs, enabled by the availability of blue and green nitride devices, is in the area of white-light illumination. White light can be generated from LEDs using two methods. In the first method, the light from separate red, green, and blue LEDs is mixed to produce white (just as the emission from red, green, and blue phosphors is mixed to produce white light from a CRT display). In the second method, violet or blue light from an LED is used to pump a phosphor, which then emits a broad spectrum of light, a technique similar to the way a conventional fluorescent lamp operates. The phosphor can be selected to produce a specific type of white light, bluish or yellowish, for example. In terms of wallplug efficiency, the first method of mixing discrete red, green, and blue LEDs is superior, since there are energy conversion losses associated with the phosphor technique. Still, the phosphor method has the advantage of simplicity and lower cost, since only one LED is used, and white LEDs of this type are currently available commercially. Both methods, nevertheless, are the subject of intense investigation and development. Recent advances have resulted in the best red AlGaInP and green GaInN LEDs becoming more efficient than halogen tungsten filament bulbs (Fig. 2). Thus, matching the efficiency of typical fluorescent bulbs in the next couple of years appears feasible. In some cases where power efficiency, long

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 133

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life, and low heat generation are important, LED illumination is already of great interest as an alternative to incandescent lighting.

11. 111-V Nitride Crystal Growth

Several different epitaxial growth techniques have been used successfully to produce thin films of the 111-V nitrides. The most common among these growth techniques are hydride vapor phase epitaxy (HVPE), molecular beam epitaxy (MBE), and organometallic vapor phase epitaxy (OMVPE). High-quality epitaxial films of GaN and its ternary alloys, AlGaN and GaInN, the building blocks of all nitride-based optoelectronic devices, have been produced by each of these methods. The first GaN epitaxial layers were grown by Maruska and Tietjan (1969) via HVPE by flowing hydrogen chloride vapor over molten gallium and reacting the resultant gallium chloride vapor with ammonia downstream at the substrate. Using this technique, extremely thick films, grown at growth

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rates around 0.5pm min-’, were produced. A few years later, Manasevit and coworkers (Manasevit et al., 1971) used the emerging technique of metalorganic chemical vapor deposition (the equivalent of OMVPE) to deposit films of GaN and AlN, albeit of poor quality, on sapphire substrates. Although it took another 15 yr to adequately address the problems of nucleation and cracking during OMVPE growth of G a N on sapphire, OMVPE would eventually emerge as the dominant growth technique used to produce thin films and optoelectronic devices of the nitrides. To date, the highest quality nitride-based optoelectronic devices have been produced by OMVPE. This fact, in addition to the scalability of this technique into a process that is commercially viable, makes OMVPE the method of choice for large-scale 111-V nitride production. Currently, the OMVPE method is used exclusively in the production of commercially available high-brightness blue and green LEDs (Nakamura et al., 1994a, 1994h 1995a, 1995b, 1995c) as well as the recently commercialized violet laser diodes (Nakamura et al., 1997). As a result, in the last decade, OMVPE has emerged as the method of choice for the vast majority of 111-V nitride research, development, and manufacturing, making this technique the focus here.

1. NITRIDEOMVPE SOURCES OMVPE is a nonequilibrium growth technique that requires the vapor transport and pyrolysis of a group 111 organometallic (or metal alkyl) source and a group V hydride component via the generic reaction

In this case, R is an organic radical (typically CH, or C,H,), M is a column I11 metal atom (i.e., Ga, Al, or In) and X is the column V atom (in the case of nitride growth, N). This simplified reaction equation, while ignoring vital kinetic and thermodynamic processes involved in the overall growth sequence, provides a convenient framework for reference when discussing the OMVPE growth process. The hallmark of OMVPE growth is its use of organometallic sources, which can be either liquids or solids at standard conditions. These sources are stored in stainless steel containers referred to as “bubblers” since the method of source delivery is to flow carrier gas (typically hydrogen or nitrogen) through the container, thus bubbling the carrier gas through the source in the case of liquids. T o control the molar flow rate of the source being delivered, bubblers are maintained at a constant temperature and

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 135

pressure. The result is that the carrier gas becomes saturated with vapor from the source and is carried through stainless steel plumbing into the reactor to the substrate (which resides on a heated graphite susceptor) where it is mixed with the hydride source at temperatures suitably high for the chemical and kinetic processes necessary for deposition to occur. For nitride growth, the typical group I11 sources used are trimethylgallium (TMGa), triethylgallium (TEGa), trimethylaluminum (TMAl), triethylaluminum (TEAl), and trimethylindium (TMIn). Group V precursors are typically gaseous hydrides with ammonia (NH,) being the most common for nitride deposition. Thus, the basic OMVPE reaction for the deposition of GaN from TMGa and NH,, the most common sources used, can be written as

Dopant sources can be either organometallic or hydride sources. For nitrides, the n-type dopants, which come predominantly from group IV of the periodic table, are gaseous hydrides, such as silane (SiH,), disilane (SizH6), and germane (GeH,). The p-type (group 11) dopants are usually organometallics like diethylzinc (DEZn) and biscyclopentadienylmagnesium ((C,H,)zMg or CPzMg). 2. NITRIDEOMVPE REACTORS In addition to the dependence on reaction kinetics, thermodynamics, and growth conditions employed ( e g , temperature, pressure, ratio of group V to group I11 reactants in the gas phase, and linear gas stream velocity), hydrodynamics also play a key role in nitride OMVPE deposition. Because NH, is a very stable molecule with a low cracking efficiency, it must be delivered into the reactor in very large concentrations and the growth temperature must be kept quite high (to enhance the thermal decomposition of NH,) if solid stoichiometry is to be maintained in the growing film. As a result, NH, typically accounts for about 25-75% of the total gas flow, thus contributing significantly to the overall flow dynamics in the reactor chamber. (This is not the case for the deposition of other 111-V compounds, such as GaAs and InP, where the hydrides generally make up about 0.1- 1% of the total flow with the balance being made up by carrier gas.) Ammonia has a higher density and lower thermal conductivity relative to hydrogen (which typically makes up the balance of the reactor’s ambient), making turbulence much more likely, especially at high growth temperatures and high growth pressures, due to thermal convection.

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OMVPE systems can be categorized generically into two classes: horizontal and vertical. The difference between the two reactor classes is the manner in which the precursors are injected into the reactor. In the case of horizontal reactors, source gases are injected into the reactor parallel to the substrate and flow horizontally through the reactor vessel. Vertical reactors, on the other hand, direct the gases downward (or upward as the case may be) in a direction perpendicular to the heated substrate. Owing to the peculiarities of nitride growth (namely high growth temperature and large NH, concentration), it is often difficult to maintain laminar flow in nitride reactors. Thus, it is common to observe thermal convection (or thermal buoyancy), which causes severe ceiling deposits in horizontal reactors and recirculation cells, producing sidewall deposits, in vertical chambers. Early nitride research, pioneered by Professor Isamu Akasaki of Nagoya and Meijo Universities and Shuji Nakamura of Nichia Chemical Industries, Ltd., focused a great deal of attention on reactor design. Efforts by Professor Akasaki and his colleagues focused on atmospheric pressure growth in horizontal and vertical tube reactors with separate, thin delivery tubes for the organometallic and hydride sources, each being delivered at a very high velocity (about 5 m s-’) (Amano et al.. 1986). Although high-quality films were produced from this reactor configuration, it was very difficult to achieve a high-quality layer uniformly across the substrate. In an attempt to improve the uniformity, Nakamura proposed a novel “two-flow” atmospheric pressure reactor geometry (Nakamura, €99lb). This reactor, shown schematically in Fig. 3, introduces the source gases, or main flow, parallel to the substrate analogous to a conventional horizontal reactor. However, a second flow, or subflow, consisting of equal parts H, and N, and directed perpendicular to the substrate, is also employed to suppress thermal convection and to bring the reactants to the substrate surface. While both of these approaches involved growth at atmospheric pressure, many recent reports have also described high-quality films grown at low pressure, where nonuniformity, thermal convection, and premature reactions between the various source gases are greatly reduced, but not eliminated. Due to the success of these approaches, a tremendous increase in nitride research and development has been spawned in recent years. Many of these new researchers, when faced with the challenges of nitride OMVPE, have concentrated a considerable amount of energy and effort on reactor design. Horizontal reactors with hot walls (Chung et al., 1997), inverted susceptors (where the wafer is facing downward from the ceiling) (Niebuhr et al., 1995), as well as two-channel (Nishida et al., 1997) and three-channel (Uchida et al., 1997) systems (where two of the channels are used to deliver the alkyls and hydrides and the third injection contains carrier gas only with the intention of keeping the ceiling clean) have been constructed to combat the

N2 +

Y

4 4

CONICAL

IR RADIATION THERMOMETER

STAINLESS TUBE QUARTZ STEEL CHAMBER \

SUBSTRATE CI

W

4

N, %

-

MAIN FLOW H, + m, TWG

--

+

H2+NH, + TMG

\

ROTATING SUSCEPTOR

'

QUARTZ NOZZLE

-

HEATER

VACUUM EXHAUST

SUBFLOW

SUSCEPTOR

-

FIG. 3. (a) Schematic design of the Nichia OMVPE AlGaInN reactor chamber. (b) The action of the subflow is to force the main flow toward the substrate and suppress thermal convection above the heated substrate. (Reprinted from Nakamura, 1991b; with the permission of the American Institute of Physics.)

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problems encountered due to thermal convection. High-speed turbopumping by a rapidly spinning wafer carrier (Liu et al., 1995), closely spaced injection nozzles (Stricht et al., 1996) and inverted flow geometries (Hanser et al., 1994) (with flow from the bottom upward toward an inverted substrate) have been employed in vertical chambers to combat similar problems.

3. SUBSTRATES FOR NITRIDES One of the critical issues for 111-V nitride epitaxy is the choice of a suitable substrate. Unlike other 111-V compounds where deposition is either homoepitaxial or on lattice-matched substrates, bulk GaN (and A1N) singlecrystal substrates are not commercially available, although they have been produced in laboratories by high-pressure solution growth (Porowski, 1996), sublimation (Balkas et al., 1997; Kurai et al., 1997; Vodakov et al., 1998), and liquid phase epitaxy (Ivantsov et al., 1997). As a result, alternate substrates, none of them capable of adequately matching the lattice and thermal properties of the nitrides, have been extensively studied. Table I shows the properties of the nitride materials and the most common substrates used in nitride epitaxy (Strite and MorkoG, 1992). By far, the majority of the epitaxial growth of nitrides has been performed on sapphire and SIC. In both cases, the large differences in lattice parameter and coefficient of thermal expansion between the substrate and the epilayer create problems with respect to producing high-quality films. However, TABLE I COMPARISON OF NITRIDE MATERIALS PROPERTIES WITH THOSE OF VARIOUS SUBSTRATES AND MARKOC, 1992) (STRITE Lattice Parameter

(4

Thermal Conductivity (W cm-' K-I)

a = 3.112

2.5

Material Aln (0001)

c = 4.982

GaN (0001)

a

= 3.189

1.3

InN (0001)

c = 5.185 a = 3.548 c = 5.760

__

6H-Sic (0001) A1,0,(0001)

a = 3.080 c = 15.12 a = 4.758 c = 12.99

Coefficient of Thermal Expansion 6 - l )

4.15 x lo6 5.27 x lo6 5.59 x lo6 7.75 x 106 -

4.9 0.5

4.2 x 4.68 x 7.50 x 8.50 x

lo6 lo6 lo6 lo6

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 139

major breakthroughs in the initiation of growth and the use of intermediate nucleation layers have enabled remarkable improvements in the quality of OMVPE-grown AlGaInN layers. Growth of these nucleation layers and GaN bulk films is discussed in subsequent sections.

4. GALLIUM NITRIDEGROWTH ON MISMATCHED SUBSTRATES As discussed previously, there is no substrate (ie., GaN or AlN) commercially available that is suitably lattice matched and thermal expansion matched for AlGaInN epitaxy. Thus, 111-V nitride growth is usually performed on either sapphire or Sic. OMVPE growth of GaN directly on either substrate at typical growth temperatures (1000-1 100°C) has been shown by many researchers to result in films with very poor electrical, optical, and structural properties. However, the insertion of a thin nucleation, or buffer, layer remedies this problem and enables the crystal grower the ability to improve the quality of the layers considerably. Specifically, when using sapphire substrates, a thin ( 200- to 300-A-thick) AlN or GaN buffer layer is deposited at temperatures between 450 and 700°C prior to growing the bulk GaN film at temperatures > 1000°C. The use of these intermediate nucleation layers markedly improves the quality and properties of the subsequently grown GaN (Akasaki et al., 1989). In contrast, growth on Sic substrates is usually initiated at high temperatures (> ZOOOOC) with the deposition of a thin (< 1000-kthick) buffer layer of AlN (Weeks et al., 1995), which has a much better lattice match to Sic than GaN (Table I). Regardless of the substrate used, high-quality GaN can be grown by OMVPE with the introduction of a suitable buffer layer between the substrate and epitaxial film. Although the closer lattice match and high thermal conductivity of Sic make it a very attractive for 111-V nitride film growth and devices, the present cost (more than 20 times that of sapphire) and availability of Sic make sapphire the more common substrate. However, commercial GaN-based LEDs on both substrates are currently available. As a result, deposition on both substrates will be described here. N

a. Nucleation Layers Grown on Sapphire The 111-V nitride films generally have the hexagonal wurtzitic structure, although the cubic phase has been produced. Sapphire has a trigonal lattice structure and its (OOOl), or c-, plane, is the most commonly used orientation for crystal growth. Epilayers of the 111-V nitrides grow with their (OOOl), or basal, plane oriented approximately parallel to the substrate surface. Due to

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its high temperature stability, resistance to aggressive chemical environments, and its relative ease of production, sapphire has long been considered as a substrate material for nitride epitaxy. Despite these obvious advantages, however, the large differences in lattice parameters and thermal expansion coefficients between the sapphire and the nitride epilayers result in polycrystalline films consisting of randomly oriented nuclei that have rough surfaces and are prone to severe cracking. To grow high-quality, crack-free nitrides on sapphire, control of the interface and, consequently, film nucleation between the two materials is thought to be essential. In 1983, Yoshida et al. (1983) were the first to realize the utility of an AlN intermediate layer for the growth of GaN on sapphire by gas-source (GS) MBE. Despite the improvements in film crystallography and optical properties, they reported that their films still showed a very high background donor concentration (> loi9 ~ r n - ~ The ) . real breakthrough for OMVPE growth of GaN occurred in 1986 when Amano et al. demonstrated that by using an AlN layer (approximately 50 nm thick) deposited at low temperatures (originally about 900- lOOO"C, but later lowered to about 500-600°C; Hiramatsu et al., 1991) prior to depositing GaN at elevated temperatures ( > 1000°C), crystal quality improved, thereby enabling specular surfaces to be produced for the first time. GaN layers grown on AlN buffer layers exhibited a two order of magnitude decrease in the background carrier concentration and an order of magnitude increase in electron mobility. Bandgap photoluminescence (PL) intensity was also improved by about two orders of magnitude along with a reduced yellow (2.2eV) luminescence emission (see Section 111, Subection 2). As deposited, the AlN nucleation layer is reportedly amorphous or partially crystalline. When heated during the temperature ramp prior to epitaxy, it undergoes solid-phase crystallization creating an A1N layer that exhibits a columnar structure and provides a seed for further growth. Hiramatsu et al. (1991) and Qian et al. (1995) have shown that the resulting films are preferentially oriented along [OOOl], misoriented by <3" in the basal plane and contain low-angle grain boundaries. The initial GaN crystallites reportedly undergo geometric selection via grain orientation competition until an oriented domain structure emerges, as shown in Fig. 4. The resulting GaN undergoes a gradual improvement in quality for the first 300 nm, after which high-quality, uniform GaN is produced. The orientation relationships between the film and substrate are GaN(0001) IJAl,0,(0001) and GaN(i2iO)JJAl,O,( 10x0). Nakamura (1991a) further developed this technique by using a GaN buffer layer deposited at 450-600°C. Using GaN buffer layers, significant improvement over A1N buffer layers was demonstrated in background donor concentration and electron mobility as well as the full width half

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 141 AIN buffer layer a AI,O, ~

Nucleation of GaN

Geometric selection 4

4

4 4 C b

4

4 tGaN

Island growth

(5) Lateral growth

I

(6) Uniform growth

,Trapezoid Clystal

qplocation sound-zona

l / m

/;-/-7=% \ Semi-mnd-zonr, (150 nm)

Faultad-zom, (-50 nm) AIN (-60 nm)

FIG. 4. Schematic diagrams showing the growth process of GaN on sapphire. (Reproduced from Hiramatsu et al., 1991; with the permission of Elsevier Science.)

maximum (FWHM) of x-ray rocking curves (XRC). Correlation of optical and structural material properties as a function of buffer layer thickness determined that 20-nm-thick layers yielded the optimal results (Nakamura, 1991a). Similar results were also reported by Keller et al. (1995) who found the optimal thickness to be 19 nm (Fig. 5). Further analysis of the formation and evolution of GaN nucleation layers has been performed using analytical techniques like atomic force microscopy (AFM), reflection high-energy electron diffraction (RHEED), and transmission electron microscopy (TEM) (Hersee et al., 1995; Wu et al., 1996b, 1996~).These studies have found that the as-grown, nominally 20-nm GaN nucleation layers consist of crystalline faceted islands containing a mixture of hexagonal and cubic crystal structure, in contrast to the reported amorphous structure when AlN layers are used. Deposition of GaN on GaN nucleation layers has also been

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al.

30

10'

0 0

10 20 30 Nucleation Layer Thickness, nm

FIG. 5. The 22-K band edge to yellow (deep) luminescence peak photoluminescence intensity ratio ( I b / I d )and FWHM of the (DOX) bound exciton peak as a function of nucleation layer thickness. (Reproduced from Keller et a/., 1995; with the permission of the IEEE.)

reported to occur via geometric selection, as pictured schematically in Fig. 6 (Wu et al., 1998b). Several groups have worked on trying to create an amorphous compliant interlayer by converting the sapphire surface to AlN,O, --x by exposing the substrate to NH,, N,, or a N, plasma at elevated temperature prior to depositing the buffer layer (Fuke et af., 1998; Grandjean et al., 1996, 1997; Kawakami et al., 1988; Keller et al., 1996b; Kim et al., 1997; Kum and Byun, 1997; Masu et al., 1995; Uchida et al., 1996). Nitridation steps are credited with roughening the sapphire surface, changing the growth mode of the nucleation layer (due to changing the substrate surface energy), and increasing the crystalline quality and electrical and optical properties of the final GaN films. Keller et al. (1996b) have reported that the length of NH, exposure has a significant effect on properties of GaN, enabling films with significantly different microstructures and properties to be grown. Exposures of 3 1 min-' at 1050°C for 60 and 300s prior to the deposition of GaN buffer layers at 600°C resulted in films with dislocation densities of lo8 and 10'' ern-,, respectively. It is thought that prenitridation of the sapphire in this manner results in the partial replacement of oxygen with nitrogen at the surface. Depending on the degree of replacement, however, the resulting surface may or may not facilitate a nucleation layer that allows for the annihilation of misfit dislocations.

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES

143

OHltimlted NL After HT ExDoeure

(a) HT GaN Growth on Optimized NL (I) Initial HT Island Growth

HT GaN Island

(11) HT Island Coalescence

(b) FIG.6. Schematic of the growth process using GaN nucleation layers (NL) on sapphire showing: (a) the structure of the low-temperature (LT) GaN nucleation layer exhibiting a rough morphology and stacking disorder, (b) the initial high-temperature (HT) growth on the nucleation layer proceeding as isolated islands containing a single mixed or screw character threading dislocation, and (c) high-temperature island coalescence resulting in additional mixed and edge character threading dislocations (TD). (Reproduced from Wu et al., 1998b; with the permission of Elsevier Science.)

b.

Gallium Nitride Growth on Silicon Carbide

In addition to sapphire, S i c is also used extensively as a substrate material for growth of 111-V nitrides. Like sapphire, S i c can withstand the aggressive chemical and the extreme thermal environments used in GaN epitaxy. In addition, Sic possesses a lattice constant and a thermal expansion coefficient that are much more compatible with GaN than those of sapphire (Table I). Similar to the nitrides, Sic is tetrahedrally coordinated and has a polar surface. Sic is commercially available in either of two

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hexagonal crystallographic (polytypic) variations, 4H- and 6H-Sic. The most commonly used growth face for the nitride materials is SiC(OOO1). Nitride epitaxy is generally performed on the Si-terminated growth face rather than the C-terminated one, due to differences in surface stability. S i c substrates can also be purchased with different orientations, on-axis and vicinal, both of which are used for nitride growth. On-axis SiC(OOO1) substrates are nearly atomically smooth, while vicinal substrates are intentionally cut and polished off-axis, creating periodic “steps” on the surface. Although Sic has a much closer lattice match to GaN than sapphire, GaN cannot be grown directly grown on S i c at high temperatures due to the fact that the SiC(OOO1) has a lower surface energy than GaN, resulting in random three-dimensional nucleation and polycrystalline films. Attempts by Weeks et al. (1994) have shown that AlN and GaN buffer layers grown at low temperatures (500-950°C, similar to those used in growth on sapphire) result in polycrystalline layers that do not undergo sufficient transformation when annealed to prove useful as a nucleation layer. However, substituting a high-temperature (> llOO°C) AlN or Al,Ga, -xN nucleation layer (< 100 nm thick) prior to GaN deposition at 1050°C results in monocrystalline films with excellent electrical and optical properties (Weeks et al., 1995). Both the AIN-Sic and GaN-A1N interfaces are abrupt and coherent with the films exhibiting parallel epitaxy with the substrate and show no significant strain. Considerable efforts have been made to characterize the AlN-Sic interface and the initial nucleation and growth of the high-temperature nucleation layer on the substrate. Interfaces between AlN and Sic are generally atomically flat and planar. The most common defects are misfit dislocations, which occur about every 30 nm. The critical thickness for A1N on Sic has been determined theoretically and experimentally to be about 4.6 nm (Tanaka et al., 1996). By growing the AlN buffer layer thinner than the critical thickness (Tanaka et al., 1997b), a marked improvement in GaN epilayer properties and a reduction in dislocation density can be achieved. In addition to misfit dislocations, other defects are introduced at the interface in areas where the substrate is not atomically flat and possesses steps (i.e., vicinal substrates). Tanaka et al. (1995, 1996) have shown that stacking mismatch boundaries arise from the different nucleation processes that occur for AIN films grown on differently oriented substrates. Their study shows that growth of thin ( % 1.5 nm) AlN layers on on-axis S i c results in pseudomorphic epitaxy, nearly atomically smooth surfaces, and no misfit dislocations. However, nitride layers on vicinal Sic substrates are not smooth, but feature coalescence boundaries between islands that nucleated on adjacent terraces and misalignment of atomic planes across the step features.

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 145 c.

Growth of Gallium Nitride Heteroepitaxial Films

In addition to nucleation layer optimization, there are several additional key variables that can be employed during the growth of heteroepitaxial GaN films in order to achieve improved properties. These variables, specifically growth temperature and V: I11 ratio, are critical, regardless of the substrate choice. Ironically, the very high growth temperatures and extremely large V:III ratios employed for the nitrides are also the major OMVPE growth parameters that separate them from all other 111-V materials systems. While the crystal grower is generally responsible for determining the precise growth parameters necessary to achieve the optimal results in their reactor, a few general guidelines can be established. Typically, the growth temperature for GaN is between 1000 and 1100°C; however, many references have indicated that the optimal growth temperature is actually within the much smaller window of 1040- 1080°C. At higher temperatures, however, surface morphology and growth rate have both been demonstrated to suffer. In general, the V:III ratio exceeds 1000 and, in fact, is quite often significantly higher. Using variable temperature Hall-effect measurements, PL, and secondary ion mass spectrometry (SIMS), Keller et al. (1995) have studied the electrical and optical properties and the chemical impurities in GaN films grown at growth temperatures of 1050 and 1080°C and V:III ratios between 1287 and 5145. The best results (lowest background donor concentration and highest electron mobility) were achieved when both the temperature and V:III ratio were at their maximum values. Room-temperature PL measurements on these same samples showed that the ratio between the band edge and yellow luminescence peaks was also dependent on these factors (Keller et al., 1995).

5. GALLIUM NITRIDE P- AND N-TYPEDOPING Of paramount importance in making any semiconductor device is the precise control of dopants in the growing layer(s). Since most optoelectronic devices involve the recombination of a conduction band electron with a valence band hole to produce photons (i.e., light), the achievement of both p-type and n-type doping are vital. In general, wide bandgap materials, such as GaN, are relatively difficult to dope, due to the presence of native defects that are potentially more stable thermodynamically in wide-bandgap materials than in other semiconductor systems.

I46 a.

R. S. KERNot al. Acceptor Doping

The 111-V nitrides can be doped by the insertion of divalent group I1 elements such as Zn, Mg, Be, Ca, and Cd into the crystal lattice. These impurities, when substituted and incorporated on the cation site (i.e., the Ga site in GaN), produce acceptor levels in the bandgap. Unfortunately, all of these impurities form deep acceptors with the most commonly used dopant being Mg (typically measured between 160-170 meV by variable temperature Hall measurements (Akasaki and Amano, 1994; Gotz et al., 1995a; Nakayama et al., 1996). When GaN is Zn-doped, the films are highly resistive due to the deep acceptor levels created, rendering Zn an unsuitable dopant for conductivity control (Pankove and Berkeyheiser, 1974; Pankove et al., 1974). Although a few studies have demonstrated p-type conductivity using Be (Salvador et al., 1996) (or a Be-0 complex; Brandt et al., 1996) and Ca (Lee et al., 1996), Mg is the most commonly used acceptor impurity for the nitrides. As in GaAs, impurity elements from column IV can be amphoteric (i.e., have the possibility of being p - or n-type depending on the lattice site they occupy) in GaN. While Si and Ge are widely recognized as donors in GaN, C is expected to act as an acceptor (Boguslawski et al., 1996). Indeed, p-type behavior in GaN:C films grown by organometallic MBE (Abernathy et al., 1995) has been reported with hole concentrations as high as 3 x 10'' c m - j and carrier mobilities of 103 cm2 V-'S-' . To date, however, there are no published reports of carbon doping by OMVPE producing p-type behavior in GaN. Due to the fact that Mg is a deep acceptor dopant, heavily compensated, and only a small portion is ionized at room temperature, incorporation levels typically run about one to two orders of magnitude higher than the hole concentration achieved. Thus, films containing hole concentrations in the 1-10 x 1017cm-3 range have Mg concentration levels well above lOI9 cm-3, seriously challenging the solubility limits of the host (GaN) material (Bour et al., 1997; Yuan et al., 1995). In fact, the hole concentration has been demonstrated to increase with increasing Mg concentration to a point, before showing a negative correlation (Fig. 7), indicative of the crossing of a solubility limit.

b. Donor Doping In contrast, donor doping is relatively easy and straightforward in GaN. Despite the fact that unintentionally doped GaN is n-type as-grown, control of the doping level and conductivity is crucial to achieving proper operation

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 147

FIG. 7. Room-temperature hole concentration @ 3 0 0 ) as a function of magnesium concentration in GaNMg. (Reprinted from Bour et al., 1997; with the permission of The Materials Research Society.)

in optoelectronic and microelectronic devices. For example, the n-GaN layers grown on sapphire, which are used in commercial LEDs, are typically heavily doped (> lo1*~ r n - ~and ) constitute more than 80% of the entire device structure to achieve both a low contact resistance and a low series resistance in the structures. The most common donor dopant in GaN is Si. Several groups have shown that the electron concentration, measured by Van der Pauw Halleffect measurements, increases linearly with the Si:Ga ratio in the gas phase and is independent of the source (silane, disilane, or tetraethylsilane) (Kaneda et al., 1996; Nakamura et al., 1992c; Rowland et al., 1995; Stricht et al., 1996; Wickenden et al., 1995). By comparing the Hall-effect data and SIMS analysis, Wickenden et al. (1995) have observed that the activation efficiency for Si donors in GaN films is nearly unity, as expected for such a shallow dopant. ~ ) easily be Although very high electron concentrations (> f O I 9 ~ m - can achieved with Si, GaN films tend to show a preference toward cracking at such high doping levels, especially when deposited on sapphire. An added complication is that there is also a relationship between film thickness and cracking. In general, lightly doped films can be grown thicker before the onset of cracking. Although this relationship is assumed to be dependent on the amount of residual tensile stress locked into the GaN layer after growth, no conclusive correlation to growth parameters (e.g., growth temperature, nucleation layer thickness, or cooling rate) has been presented. Interestingly, Ruvimov et al. (1996) have reported recently that Si doping to levels about

148

R. S. KERN et al.

3 x 1Ol8cm-3 actually results in a GaN layer with improved crystal quality without cracking. In this report, the dislocation density was reduced from 5 x lo9 cm-' in undoped samples to 7 x 10' cm-' in doped samples with an accompanying decrease in the XRC FWHM by almost a factor of two. In addition to Si, 0, and Ge have also been used to produce lowresistivity n-GaN films. However, the relative ease of doping with Si has made the pursuit of alternate donor dopants a minor consideration. Although oxygen has been studied as one potential candidate, the major interest in studying oxygen has been to determine the exact nature of 0 in GaN and its assumed role as a potential residual donor present in unintentionally doped GaN (Seifert et al., 1983). Implanted 0 in GaN has been shown to be a shallow donor with an ionization energy of 29 meV (Zolper et al., 1996). Layers of GaN:O have been grown using N,O as the dopant source (Niebuhr et al., 1997). Controlled doping using 0 has been achieved from 4 x lo" to 4 x 10'8cm-3 with an approximately 50% activation efficiency for the 0 donors. Nakamura et a!. (1992~)have also demonstrated that electron concentrations up to 10'' cm-3 with mobilities similar to those in GaN:Si films can be achieved in GaN:Ge. Germanium doping, like Si, is linear with source flow rate (GeH,); however, morphology degradation and low activation efficiency (one order of magnitude lower than Si) make Ge less attractive than Si.

6. ALLOYGROWTH For nitride LED and LD applications, the growth and conductivity control of GaN and its alloys, AlGaN and GalnN, are vital to the successful production of operational devices. In LEDs, p-type AlGaN is used as the upper confining layer in double-heterostructure (DH) and single-quantum well (SQW) commercial structures. Both n- and p-type AlGaN are necessary as cladding layers surrounding the multiple-quantum-well (MQW) active region in violet LDs. Wavelength control in these optoelectronic devices is achieved by controlling the composition of the GaInN active layers. u.

Aluminum Gallium Nitride

High-quality films of aluminum gallium nitride have been achieved in both low pressure and atmospheric pressure OMVPE systems. However, at near atmospheric pressure, AlGaN growth is plagued by (parasitic) prereactions that occur in the gas phase between TMAl and NH, forming an involatile adduct (Chen et a/., 1996). Parasitic reactions between TMAl and

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 149

NH,, which are more likely than those between TMGa and NH,, effectively reduce the gas phase concentration of both reactants making deposition of A1N and AlGaN difficult, especially at higher reactor operating pressures. Two-flow channel reactors (Schmitz et al., 1996), which keep the sources separate up to the substrate, and low-pressure operation (Chen et al., 1996) have been demonstrated as successful methods for reducing parasitic reactions. Although AlGaN is a vital component in both nitride LEDs and LDs, considerably less effort has been expended on it relative to the other layers that make up the device structure(s). Alloys of AlN and GaN are typically grown at temperatures similar to those used for GaN alone. Reducing the reactor operating pressure (Chen et al., 1996), increasing the growth temperature (Hirasawa et al., 1993), and increasing the TMAl:(TMAl + TMGa) ratio (Bremser et al., 1996) have all been shown to increase the AlN mole fraction in AlGaN. In general, p-type doping is more difficult to achieve in AlGaN layers than in GaN, especially with higher A1N mole fraction in the solid phase. This is due to the higher bandgap of the AlGaN semiconductor and the increased depth of the acceptor level, as is described in Section 111, Subsection 3. As a result, the highest A1N mole fraction AlGaN layer reportedly doped p-type is 30% (Nakamura et al., 1995b);however, most researchers indicate difficulty in achieving hole conductivity above about 15% since the acceptor level continues to move deeper into the bandgap above the AlGaN valence band as the AlN mole fraction increases (Suzuki et al., 1998). Donor doping (with Si) is considerably easier, perhaps due to the fact that Si remains a shallow dopant even as the semiconductor bandgap increases (Murakami et al., 1991).

b.

Gallium Indium Nitride

Growing GaInN involves decreasing the deposition temperature substantially from that used in GaN and AlGaN deposition. Due to the low dissociation temperature of InN (Koukitu et al., 1996), the growth of highquality GaInN must be performed at temperatures of 850°C or lower. Complicating matters, the cracking efficiency of NH, decreases dramatically with decreasing temperature due to the high kinetic barrier for breaking the N-H bonds in the molecule. Since temperatures used in the growth of GaInN are higher than the decomposition temperature of the TMIn source (Buchan et al., 1988), the result can be the formation of In droplets if a suitable amount of activated N is not present on the surface of the growing layer (Bedair et al., 1997; Lu et al., 1997; Shimizu et al., 1994; Sun et al., 1996).

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Other matters related to the materials system also make high-quality GaInN layers difficult to produce. Lattice mismatch between both GaInNGaN and GaInN-AlGaN can be as high as 10% or more, generating misfit dislocations and promoting three-dimensional growth and mosaicity even at very low layer thicknesses (Sato et al., 1997). Poor interfacial areas between different compositional layers are also likely due to poor nucleation of the mismatched layer due to low deposition temperature, lattice parameter differences, and low surface mobility of adatoms. Phase separation (Romano et al., 1998; Singh et al., 1997) and spinodal decomposition (El-Masry et al., 1998) due to solid phase immiscibility (Ho and Stringfellow, 1996; Wakahara et al., 1997) have also been reported for GaInN. To combat these problems, several techniques have been simultaneously applied to achieve the highest quality GaInN layers. First, GaInN films are grown at temperatures that are much lower than those used for G a N to reduce the dissociation of In-N bonds and the reevaporation of In. It is found that significant In does not begin t o incorporate until growth is performed below 850°C (Nakamura and Mukai, 1992). Second, GaInN films are often grown using V:III ratios significantly higher than those used in GaN growth to offset the reduction in NH, cracking efficiency and to partially suppress the decomposition of the In-N bonds (Grudowski et al., 1997; Koukitu et al., 1997). Third, due to etching of In by H, even at very low concentrations, N, is usually used as the carrier gas for GaInN growth (Koukitu et al., 1996; Nakamura and Mukai, 1992; Piner et ul., 1997). Fourth, the brightest SQW LEDs and MQW LDs employ quantum wells 3nm or less in thickness (Nakamura et al., 1995a, 1995b, 1995c), thus reducing the effects of the lattice and phase incompatibility that exists. Additional methods such as flow modulation epitaxy (Keller et al., 1997), the use of compositional grading (Keller et a!., 1996a), and N, (Wakahara et al., 1997) and NH, (Sato, 1997a, 1997b) plasma-assisted OMVPE have also been used with success to grow high-quality GaInN and InN layers. Despite these advances in the deposition of high-quality GaInN, differences in the growth conditions between GaN or AlGaN and GaInN force the crystal grower to successfully implement at least two growth interrupts to achieve the gradients in temperature, growth ambient, and NH, concentration (as well as pressure, which exists as an additional potential parameter) necessary to deposit nitride optoelectronic devices.

AREAEPITAXY 7. SELECTIVE Selective area epitaxy is a technique that enables the crystal grower to engineer the shapes, facets, and dislocation densities present in G a N epilayers. Selective area growth requires growing thin films of G a N or

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AlGaN, masking regions of the film surface with a dielectric (such as SO,, although Si,N, has also been used; Beaumont et al., 1998), and resuming growth on the patterned substrate. Growth will proceed on the unpatterned areas of the surface, maintaining the crystal and dislocation structure of the underlying layer. However, nucleation will not occur on the patterned areas. By properly selecting the pattern, its orientation, and the GaN regrowth conditions, a variety of different results can occur. When a continuous dielectric mask with round, square, or hexagonal openings is deposited on the GaN layer, GaN hexagonal pyramids with { 1TOl) sidewalls (Hiramatsu et al., 1996; Kapolnek et al., 1997b; Nam et al., 1997a; Zheleva et al., 1997) or hexagonal microprisms with { 1TOO) sidewalls and a (0001) top face (Akasaka et al., 1997) result during the deposition of the regrowth layer, depending on the growth conditions. Arrays of undoped and Si-doped GaN hexagonal pyramidal structures have been fabricated into electron field emission devices (Kapolnek et al., 1997b; Nam et al., 1997a). As relates to optoelectronics, however, the most interesting and promising selective area growth experiments are those that involve growth between stripe openings in oxide masks. Interestingly, anisotropic lateral growth rates have been observed in GaN selective area epitaxy. Lateral growth rates have been studied using masks with various SiO, linear stripe openings spaced in rotational increments of a few degrees (Kapolnek et al., 1997a; Park et aE., 1998). When the oxide lines were oriented parallel to the (1 120) GaN directions, overgrown layers had triangular cross sections and showed a very low ratio between the lateral and vertical growth rates. GaN overgrown layers on substrates with lines oriented parallel to (lTOO), however, displayed lateral to vertical growth rate ratios of up to 4.1 (Kapolnek et al., 1997a), resulting in significant lateral mask overgrowth and oxide coverage. These behaviors exhibit hexagonal symmetry with the maxima in lateral growth rate occurring at separations of 60". Further experiments have demonstrated that the growth anisotropy is strongly related to the growth temperature and ammonia flow (Kapolnek et al., 1997a; Park et al., 1998). Lateral growth rates are enhanced with increases in both ammonia flow and growth temperature. Selective area growth has been performed using thin stripes of dielectric that run from one side of the wafer to the other. It is found that selective growth occurs for stripe orientations running along (1 120) directions (Akasaka et al., 1997; Nam et al., 1997a; Zheleva et al., 1997). For stripes oriented in these directions, growth occurs essentially between the dielectric stripes only with a small amount of lateral overgrowth of the dielectric stripes relative to the vertical growth between the stripes. Prismatic stripes of both GaN and AlGaN with (1TOl) side facets and ridge walls (i.e., no truncation present) have been grown. It has been demonstrated that reductions in the V:III ratio (by increasing the Ga organometallic molar

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flow rate) result in a decrease in area of the top (0001) facet, supporting the model that the morphology of selectively grown GaN is dependent on the balance between incoming source flux on the (0001) facet and diffusion to the {liOl} side facets. Thus, for increased alkyl flow, the side facets are stabilized preferentially. By using linear SiO, mask features oriented in the (li00) directions, the overgrown GaN will actually laterally overgrow the oxide mask and form a continuous overgrown layer (shown in Fig. 8) (Marchand et al., 1998; Nakamura et al., 1997; Nam et al., 1997b). This technique, referred to as “epitaxial laterally overgrown G a N or ELOG, has recently been the focus of much study because of the greatly improved structural and optical properties in the GaN that grows laterally above the oxide mask. While the GaN grown homoepitaxially between the oxide stripes replicates the “seed layer” on which it is grown, the threading dislocation density is reduced dramatically in the GaN which has laterally overgrown the oxide stripes. Studies by TEM (Zheleva et al., 1997) indicate that this reduction is due to both dislocation blocking of the vertically propagating dislocations by the oxide layer as well as an abrupt change in the propagation direction (from vertical to lateral) of some dislocations at the oxide edge. Substantial reductions in the density of these vertical threading dislocations (reportedly by at least three to four orders of magnitude; Marchand et al., 1998; Nam et al., 1997b) as well as etch pit densities (Marchand et al., 1998), biaxial strain (Kapolnek et al., 1998), and defect concentration (Freitas et al., 1998) directly above the masked areas make these areas particularly interesting for both materials characterization and device fabrication. Despite the improved crystalline quality, the use of ELOG substrates has not resulted in more efficient LEDs (Mukai et al., 1998~).This is perhaps due to the fact that the device area for an LED is sufficiently large to cover multiple oxide pattern and regrowth stripes, thus resulting in only a minor improvement in overall material quality. However, Nichia Chemical Industries, Ltd., has announced the achievement of violet LDs with lifetimes surpassing 10,000 hours grown on ELOG substrates (Nakamura et al., 1997). After ELOG growth, the Nichia LD stripes are made directly above the oxide stripe, thereby realizing the benefits of increased GaN quality. As a result of their progress in this area, Nichia commercialized their LDs in January 1999. 111. Electronic Properties and Characterization of III-V Nitride Films

At present, the fabrication and optimization of devices based on III-V nitrides is still limited by advancements in materials technology. Control and optimization of the doping is one of the keys to improve device

u l W

Low dislocation ocatioo GaN

FIG. 8. Epitaxial lateral overgrowth of GaN layers. (a) Schematic diagram showing ELOG regions within stripe openings in a mask deposited on a GaN epilayer. (b) Schematic diagram showing the coalescence of the overgrown layer. (c) Transmission electron micrograph of the high and low defect concentrations in an ELOG sample. (Reprinted from Nam et al., 1997b; with the permission of the American Institute of Physics.)

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performance. In most III-V semiconductor materials, extended defects have the potential to introduce nonradiative recombination centers or limit the reliability of optoelectronic devices. However, understanding of these fundamental materials properties for III-V nitrides is still in its infancy. Several areas have been the focus of intense investigation by experiment and theory and numerous important discoveries have been made. For example, as-OMVPE-grown Mg-doped GaN is electrically insulating due to the passivation of the Mg acceptors by hydrogen (Gotz rt ul., 1996b; Neugebauer and Van de Walk, 1995), and the observation by TEM that GaInN quantum wells (QWs) are structurally inhomogeneous (Ponce et al., 1998). This section is devoted to structural and electronic properties of III-V nitride materials. We briefly review the present understanding of the structural properties of heteroepitaxially grown 111-V nitride layers. We summarize results from photoluminescence studies and discuss in more detail the characterization of dopants and common impurities by Hall-effect measurements because this technique has been instrumental for device optimization. Such optimization can only be accomplished when fundamental properties of donors and acceptors in these novel materials are well understood and characterized.

PROPERTIES 1. STRUCTURAL Heteroepitaxial nitride films grown on sapphire substrates by any technique contain high densities of structural line and planar defects. These structural defects are generated due to the large lattice ( - 14%) and thermal expansion coefficient mismatch (- 80%j between the 111-V nitride epilayers and sapphire substrate, as presented in Table I. The structural properties of these layers, especially GaN, have been studied by x-ray diffraction techniques (Akasaki et af., 1989) and crosssectional TEM (see e.g., Liliental-Weber et al., 1996; Ponce, 1997). Typically, on sapphire substrates III-V nitrides grow as wurtzitic crystals since the hexagonal lattice is the equilibrium crystal phase of 111-V nitrides. Under conditions far from thermodynamic equilibrium, such as low temperature, III-V nitrides can crystallize in the metastable zincblende structure, especially when deposited on selected faces of substrates with cubic crystal structures, such as GaAs. For growth of GaN on c-plane (0001) sapphire substrates the films are aligned with the [Tolo] direction parallel to the [lZlO] direction of sapphire (Akasaki et al., 1989). This corresponds to a 30" rotation of the III-V nitride basal plane with respect to the A1,0, substrate. As a consequence, the lattice mismatch Aa between III-V nitride

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heteroepitaxial layers and sapphire substrates can be computed from the a-lattice parameters (Table I) as follows:

Aa

=

J3anitride

- asapphire

asapphire

(3)

The most prominent extended defects observed by TEM are threading dislocations that propagate nearly parallel to the c axis. Their density ranges between lo8 and 101'cm-' even after several micrometers of film growth (Lester et al., 1995). They originate at the GaN-sapphire interface, where the growth of GaN is typically initiated on top of thin GaN or AIN nucleation layers grown at low temperatures (see Section 11). Nucleation layers of AlN and GaN comprise three-dimensional nuclei that wet the sapphire substrate, reducing the interfacial energy differences and the strain between the substrate and growing film. The near interface region is also often dominated by stacking faults and in-plane edge dislocations (Ponce, 1997). Inversion domain boundaries have also been reported to occur in the III-V nitrides (Romano et al., 1996). Generally, the crystal quality dramatically improves with increasing film thickness and some of the dislocations annihilate. Depending on preparation of the nucleation layer, these dislocations can have pure edge, screw, and mixed screw-edge character; the latter two allowing for two-dimensional step-flow growth (Kapolnek et al., 1995). Figure 9 demonstrates a cross-sectional TEM micograph taken for an OMVPE-grown GaN layer under conditions that highlight the dislocation distribution. In planar-view TEM micrographs taken several micrometers away from the GaN-sapphire interface, the threading dislocations tend to appear in cellular arrangements, suggesting a columnar or cell-type structure of the heteroepitaxial GaN films (Cheng and Zheng, 1997; Kapolnek et al., 1995; Ponce, 1997; Wu et at., 1996a). The individual columns are slightly twisted with respect to each other and the c axis of each column is tilted with respect to the growth direction. Thus, their boundaries require threading dislocations. The columnar nature and the tilt and twist components are illustrated in Figs. 10a and lob, respectively. Nitride films of high crystalline quality exhibit small tilt and twist angles, which can be measured from c-oriented XRD reciprocal space maps and XRC using the symmetric (ie., 002) and asymmetric (i.e., 204) reflection, respectively. The diameters of the single crystal domains are typically smaller than 1 ,um and they have a distribution in lattice orientation of about 5 arcmin. Spatially resolved cathodoluminescence (CL) measurements indicate that the optical properties of GaN are inhomogeneous and can be correlated to

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FIG. 9. Cross-sectional TEM micrograph of OMVPE-grown GaN grown on c-planeoriented sapphire. The image demonstrates a defective GaN layer near the GaN-sapphire interface. Above this layer threading dislocations associated with cell boundaries are observed; ’ g indicates the diffraction vector. (TEM micrograph provided by L. T. Romano, Xerox Palo Alto Research Center, Palo Alto, California 94304.)

‘&

11

the dislocation distribution and density, indicating that microstructure and luminescence are related (Rosner et al., 1997). Nucleation layer growth conditions (i.e., growth rate, temperature, and thickness) cause variations in the size and density of nuclei and, consequently, layer properties. As a result, background donor concentration, electron mobility, PL, and CL of GaN films are very sensitive to slight variations in the growth conditions of these layers. To date, several groups (Doverspike et al., 1995; George et al., 1995) have attempted to determine whether the composition (AlN vs GaN) of the buffer layer plays a significant role in determining the properties and defect structure of the nitride epilayers; however, the results are somewhat inconclusive. Contrary to other materials used for optoelectronic devices, such as 111-V arsenides and 111-V phosphides, where dislocation densities above lo4 cm-’ adversely affect the device performance, the realization of highefficiency LEDs has been possible with HI-V nitrides (Lester et al., 1995) despite the fact that these films exhibit dislocation densities 10,000 times higher than this threshold. It has been demonstrated that threading dislocations cross through the active GaInN layer of high-performance blue LEDs; yet these devices still exhibit external quantum efficiencies at least as

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FIG. 10. Illustration of the columnar structure of 111-V nitrides. The films consist of coloumns with a slight distribution in the crystalline orientation with two components: (a) tilt and (b) twist. (Reprinted from Ponce, 1997; with the permission of the Materials Research Society.)

high as -4% at 20 mA (Lester et al., 1995). It is still a question whether the dislocations are associated with electrically active defects that would act as effective nonradiative recombination centers and effectively limit the lifetime of minority carriers.

2. LUMINESCENCE PROPERTIES All wurtzite 111-V nitride semiconductors possess a direct bandgap (Lambrecht et al., 1996). As a consequence, absorption spectra of 111-V nitrides exhibit a sharp absorption edge near the bandgap energy (E,) with an absorption coefficient of lo5 cm-' for photon energies (hv)larger than E,. For GaN, the optical properties hv > E, have been examined by reflectance measurements. The optical response functions were used to provide experimental verification for theoretical band structures (Lambrecht et al., 1996). For direct, wide-bandgap semiconductors at room temperature, the intrinsic optical properties in the near band edge are dominated by excitons. Three fundamental exciton states exist in GaN, which are typically labeled

-

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as A , B, and C free excitons (FE,, FE,, and FE,). These three excitons are composed of electron-hole pairs where the holes are associated with the three non-degenerate valence band (VB) maxima that exist at the r symmetry point of hexagonal GaN. The near-bandgap electronic structure near the symmetry point for wurtzite GaN (C6”symmetry) is depicted in Fig. 11.Due to the symmetry of the valance band maxima T,(FE,), T,(FE, and FE,), and the conduction band (CB) minima r7,the exciton recombination follows certain selection rules. For example, all three excitons can be observed in the so-called “a-polarization” for which the propagation vector of the excitation light is parallel and the electrical field vector is perpendicular to the c axis of the crystal. The most powerful technique to study excitons is PL, and there is a rich literature on the investigation of FEs in GaN (for a summary, see Monemar, 1998).The energy position of these excitons is very sensitive to the strain in the C a N crystal and, hence, depends on the choice of substrate and growth conditions, As a consequence, the energy positions of FEs vary in the different studies. For high-quality samples investigated far away from the substrate interface or for homoepitaxial samples, the following energy positions seem to be consistent (2 K, 2 meV accuracy): 3.478 eV (FE,), 3.484 eV (FE,), and 3.502 eV (FE,) (Monemar, 1998). A binding energy of 26 i-2 meV for the lowest energy exciton (FE,) is quoted in the literature (Eckey et at., 1996) which places the lowest bandgap of GaN (I‘, (CB) to r,(VB)) at 3.504eV (2 K). Due to the effect of strain on the bandgap energy the recombination energy of excitons shifts to higher energy values for growth on c-plane E

FIG. 11. Electronic band diagram for wurtzite GaN near the direct bandgap at the symmetry point.

r

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 159

sapphire (compressive strain in the basal plane) and to lower energies for growth on 6H-Sic substrates (tensile strain in the basal plane) (Monemar et al., 1996). The measurement of the decay time for the recombination of FEs after an excitation pulse provides an important parameter to assess the quality of a direct bandgap semiconductor. The presence of nonradiative recombination centers that introduce electronic states into the bandgap near midgap may strongly limit the lifetime of electron-hole pairs injected into the active region of a light emitting diode and degrade the internal efficiency of the device. This decay time has been measured for GaN by transient P L measurements (time-resolved PL) and the longest reported decay time for the A exciton is 200 ps (Monemar et al., 1996). In the presence of electrically active shallow defects, bound excitons (BEs) coexist with free excitons and are observed in low temperature PL spectra. They can be either bound to neutral donors (donor bound excitons, DBEs) or neutral acceptors (acceptor bound excitons, ABEs). An example of a DBE as observed in a low-temperature PL spectrum is depicted in Fig. 12 (BX). The recombination energy of BEs is lowered from the recombination energy of the lowest energy FE (FE,) by the binding energy of the exciton

Photon Energy (ev)

“Yellow PL Band”

Wavelength (nm) FIG. 12. PL spectrum taken at 2 K for Si-doped, n-type GaN. PL emission lines are labeled “ B X (bound exciton), DAP (donor-acceptor pair recombination), and “yellow PL band.”

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to the donor or acceptor. This binding energy is characteristic of the defect and typically only a fraction of the ionization energy of the donor or acceptor. PL spectra of high-quality, n-type GaN films are typically dominated by sharp DBE lines. Binding energies of DBEs to the respective donors in the range of 3 to 6 meV are reported in the literature (Monemar, 1998). No clear identification of the defects that are associated with the DBEs is available to date, but it is likely that they are associated with Si and 0, which are shallow donors in GaN (see Section 111, Subsection 3). The intensity of a DBE line decreases as the sample temperature is increased due to ionization of the shallow donors. As a consequence, at room temperature, the near band edge signal of n-type GaN is composed of FE lines. ABEs have also been reported for GaN but again a clear identification of the related acceptors has not yet been accomplished. For 1 to 3-pm-thick OMVPE-grown GaN epilayers doped with Mg and Zn, the ABE binding energy for Mg and Zn acceptors is 19 and 34 meV, respectively (Kaufmann et al., 1996). However, the binding energies are strain dependent and ideally should be measured in strain-free material. Electronic transitions between donor and acceptor states (donor-acceptor pair recombination, DAP) and between the conduction band and acceptor states have also been observed by PL. A common example for a DAP spectrum (Grimmeiss and Monemar, 1970) taken at 2, 50, and 110 K is shown in Fig. 13. At 2K, the zero-phonon line is dominant in the spectrum and appears at 3.27 eV for this sample. At the low-energy side of the zero phonon line, there are several LO phonon replicas resolved. As the sample temperature increases, the zero-phonon line becomes weaker and a high-energy shoulder appears that is labeled “CB.” The reduction of the intensity of the DAP zero-phonon line is due to the ionization of the shallow donor and the appearance of the shoulder is most likely due to an electronic transition from the conduction band minimum to the same acceptor state. The energetic distance between CB and DAP of 23 Ifr 2meV yields the position of the shallow donor level below the conduction band minimum. Using the GaN bandgap energy of 3.50.3 & 0.0005 eV (Monemar, 1974), the position of the acceptor level is equal to 208 4meV. The chemical identities of the donor and acceptor involved are not clearly established, but it has been suggested that Si and Mg are involved (Gotz et al., 1996a). For thick GaN films grown by HVPE (see Section 11),infrared absorption measurements were successfully employed to resolve the ls-2p electronic transitions from a shallow donor (Meyer et al., 1995). The absorption line appeared at an energy of 26.6 meV in the absorption spectrum. Under the assumption of a hydrogen-like shallow donor the position of the donor level is estimated to be located -35 meV below the conduction band minimum.

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES161 MOnEnergy (ev)

3.4

3.2

3.0

2.8

FIG. 13. PL spectra for a Si-doped, n-type GaN sample taken at (a) 2 K, (b) 50 K, and (c) 110 K. Emission peaks related to shallow donors are labeled “BX,” “CB,” “DAP1” and indicated by vertical lines. (Reprinted from Grimmeiss and Monemar, 1970, with the permission of the Electrochemical Society.)

It is possible that this donor level is due to the presence of oxygen in the HVPE GaN material (see Section 111, Subsection 3 ). A prominent feature in PL spectra of n-type GaN is the so-called “yellow luminescence.” The yellow luminescence is a broad emission band centered at -2.2eV and is also depicted in Fig. 12. Due to its emission wavelength, the recombination process that gives rise to the “yellow luminescence” must contain at least one deep level. Neugebauer and Van de Walle (1996a) have suggested that the Ga vacancy point defect, which introduces a deep acceptor level into the bandgap of GaN, is likely to be involved. One of the most studied topics in the 111-V nitride materials system at present is the exact nature of the emission mechanism from GaInN active layers. Some experimental data seem to indicate that emission from the GaInN quantum wells occurs due to localized excitons from potential minima that arise from compositional inhomogeneities. To date, conventional cross-sectional TEM (Chichibu et a/., 1997; Narukawa et al., 1997; Ponce et al., 1998; Taki et a/., 1997), quantitative analytical scanning TEM

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(Lakner et al., 1997), convergent beam electron diffraction (CBED) (Lakner et al., 1997), spatially resolved CL (Chichibu et al., 1997; Ponce et al., 1998), time-integrated PL (Allegre et al., 1997), energy dispersive x-ray microanalysis (Narukawa et al., 1997), and chemical imaging maps of the local electron scattering potential (Kisielowski et al., 1997) have been used to advance this theory. According to the theory, GaInN nucleates or phase separates into InN-rich and InN-poor regions within the quantum well due to immiscibility in the InN-GaN system, localized strain relaxation due to dislocations, the mosaic structure of the underlying GaN, or the textured surface on which the quantum wells are grown. LED and LD light emission is then produced from the lower bandgap (higher InN mole fraction) regions that shows very high efficiency at low current. Evidence of a “blue shift” in GaInN LEDs when stressed to higher and higher drive currents also apparently supports these assertions (Eliseev et al., 1997). I n an attempt to confirm this model, GaInN has been intentionally forced to grow on pyramidal features (Kapolnek et al., 1998), which exhibit the same { 1071) side planes as the “V-defects” associated with localized excitons formed during the growth of multiple quantum wells (Chen et al., 1998b; Wu et al., 1998a). The side planes are created by selective area growth (see Section 11) on lattice planes other than (0001). Forced three-dimensional nucleation of GaInN “quantum dots” by preexposing the GaN or AlGaN underlying layer with Si (Hirayama et al., 1997; Keller et al., 1998; Tanaka et ul., 1997a) has also been proposed as a method of achieving localized excitonic transitions in the GaInN layers. In both cases, the emission wavelength is red shifted and the luminescence intensity is increased dramatically. Although the evidence presented here appears to indicate that unity has been achieved on this issue, it must be noted that an alternative model is needed to explain the GaInN blue shift with increasing excitation energy. That model uses piezoelectric considerations and the quantumconfined Stark effect to explain these observances (Takeuchi et al., 1997a; Takeuchi et al., 1998).

3. PROPERTIES AND HALL-EFFECT CHARACTERIZATION OF DOPED III-V NITRIDES The optimization of the efficiency of LEDs requires precise control of the doping. The distribution of both donors and acceptors near the active region of the device must be suitably engineered to maximize the injection efficiency.Also, the dopant concentrations in the semiconductor at the metal contacts must be optimized to achieve ohmic behavior for the contacts to the n- and p-type portions of the diode to produce diodes that operate at low forward voltages. To accomplish such optimizations, the electronic

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LICHT-EMITTING DIODES 163

properties of intentional and unintentional dopants and their incorporation during OMVPE growth must be understood. Both impurities and native defects may act as donors or acceptors and affect the electrical conductivity of semiconductors. For III-V nitrides, Si and Mg, which substitute for Ga in the GaN lattice (SiGa,Mg,,), are the donor and acceptor impurities of choice for doping III-V nitrides nor p-type, respectively (Amano et al., 1989; Nakamura et af., 1992a; Nakamura et al., 1992~).0 is suspected to be a shallow donor (Chung and Gershenzon, 1992) and may be incorporated unintentionally during growth. Zn is a deep acceptor and was utilized in the first commercial nitride LEDs to tune the light emission (Nakamura et al., 1994a). Among native point defects, nitrogen vacancies (a donor) and Ga vacancies (an acceptor) have the lowest formation energies but are expected to play a role only as compensating centers (Neugebauer and Van de Walle, 1994a). Prior to the development of the nucleation layer routines on sapphire and Sic, GaN films were generally n-type as-grown and very conductive. With the implementation of buffer layers, most of the GaN community is currently reporting background electron concentrations of loi7cm-3 or lower. A remaining mystery is the origin of the donor responsible for this relatively high doping level. For years, the speculation was that a deficiency of nitrogen (i.e., nitrogen vacancies) were responsible for the background electron concentration (Pankove et al., 1974); however, other candidates have emerged including impurities such as 0 and Si (Neugebauer and Van de Walle, 1994b; Seifert et al., 1983) (presumably from the source precursors or quartz reactor walls). It has also been suggested that donors are inhomogeneously distributed in III-V nitride heteroepitaxial films and that the high n-type background doping arises from donors situated at the structurally defective region near the GaN-sapphire interface (Fig. 9) (Gotz et al., 1998b). These donors may be of intrinsic nature, but the incorporation of 0,available from the sapphire substrates, is also plausible. The situation for acceptors is complicated by the fact that as-grown, acceptor-doped material produced by OMVPE is semi-insulating. The presence of electrically inactive acceptor-hydrogen complexes has been attributed to the observed phenomena (Nakamura et al., 1992a). In the following sections, the electronic properties of Si and 0 donors, Mg acceptors, and the role of hydrogen for doping of III-V nitrides are reviewed.

a.

N-Type Gallium Nitride

Flowing SiH, during OMVPE growth to introduce Si commonly achieves n-type doping (Gotz et al., 1996c; Hacke et al., 1994; Kaneda et al.,

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1996; Wickenden et al., 1995). Si atoms efficiently incorporate onto G a sites in the GaN lattice. However, possibly due to the small ionic radius of Si donors (Si': 0.41A, Ga' : 0.6281), the c-lattice parameter shrinks and the basal plane of the GaN layer is put into lateral tensile strain eventually exceeding the degree of compression imposed on the GaN by the sapphire substrates. For high Si concentrations, the tensile strain leads to the formation of cracks at Si-doping levels exceeding -2 x 1019cm-3 for thicker films ( > 1 pm). For 3-pm-thick films, as utilized in light emitting diodes, the Si-doping levels are limited to -5 x 1018 cm-3 to avoid cracking. For thin films, however, Si-doping levels of -1020cm-3 can be achieved. The highest reported mobility for thick heteroepitaxial films of GaN is 900 cm2 V s-' at 300 K at a relatively low background electron concentration of 2 x 1016cm-3 (Nakamura et al., 1992b). The electronic properties of shallow dopants are best characterized by variable temperature Hall-effect measurements. The analysis of the temperature dependence of the carrier concentration enables determination of important dopant parameters, such as the concentration, the thermal activation energy for dopant ionization, and the concentration of compensating acceptors and donors in n- and p-type semiconductors, respectively (Gotz et al., 1996~). The Hall-effect measurement yields the sheet resistance (rSq) and the Hall coefficient (RH(n,p)) of a GaN thin film. Under the assumptions of film thickness uniformity and of uniform transport properties within the films, the resistivity (p), the carrier concentration (n,p), and the carrier Hall mobility (pn,11,) can be determined from the measured quantities as follows:

-

p

= rSq.d

(4)

and

In Eq. (5) rH is the Hall scattering factor, which is unknown for GaN. It is typically assumed to be isotropic and of unity value. In case of inhomogeneities, Eqs. (4)-(6) are inapplicable (Gotz et al., 1998b). The temperature dependence of the electron and hole concentrations of the GaN films can be employed to derive parameters for shallow donors and acceptors, respectively. For n-type semiconductors, in the case where the

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 165

temperature dependence of the charge transport is dominated by ionization of shallow donors, the electron concentration is given by (charge neutrality condition)

In Eq. (7), m is summed over the number of distinct donor species; N i and AE, are the concentrations and the activation (thermal ionization) energies is the concentration of compensating of the donors, respectively; Ncomp acceptors, Ncveff(7') is the effective density of states for the conduction band; gi is degeneracy of the donor states, and k is the Boltzmann constant. A similar expression applies for acceptors in p-type semiconductors. A leastsquares fit of Eq. (7) to the experimental n ( T ) or p ( T ) data yields concentrations N i and activation energies AEi for donors or acceptors, respectively. For the data presented in this section, the degeneracy factors for both donors and acceptors were taken as two and the effective density of states for the conduction band and the valence band was computed with an effective mass of 0.2m0 for electrons and of m, for holes (mo = free electron mass). Figure 14 demonstrates results from variable-temperature Hall-effect measurements for five n-type GaN samples. The experimental results for each sample are represented by unique symbols. Sample 1 was unintentionally doped and samples 2-5 were Si-doped. The electron concentrations are shown as functions of reciprocal temperature in Fig. 14a and the electron mobilities are shown as functions of temperature in Fig. 14b. The unintentionally doped GaN film (sample 1) exhibits the lowest electron concentration ( - 5 x 1016cm-3 at 300K) and the highest electron mobility ( 650 cm2 V - l sC1 at 300 K). The Si-doped samples show a progression of higher electron concentrations and lower electron mobilities as the SiH, flux was gradually increased from samples 2 to 5. Sample 5 exhibits a temperature independent electron concentration of -2 x 10'' ~ m - Such ~ . behavior is typically observed for degenerately doped semiconductors. The peak mobilities for the n-type GaN samples shown in Fig. 14 range from 135 (sample 5, 300 K) to 1450 cmz V-' sC1 (sample 1, 130 K) and decrease with increasing electron Concentration (Fig. 14b). The solid lines in Fig. 14a result from least-squares fits of the charge neutrality condition as already described. The fits yield parameters for a shallow donor level that is dominant in the GaN films. The modeling yields activation energies for the shallow level of 16meV (sample 1), 14meV (sample 2), 12meV (sample 3), and 7 meV (sample 4). The thermal ionization energy as deduced from Hall-effect measurements is expected to decrease with increasing donor concentration

-

166

R. S. KERNet al. (a)

Temperature (K)

1000/T (VK)

(b)

Temperature (K)

FIG. 14. (a) Electron concentration vs reciprocal temperature and (b) Hall electron mobility vs temperature for GaN doped with increasing Si concentration. The symbols refer to experimental data while the solid lines in (a) results from least-squares fits to the experimental data.

(Schubert, 1993). The concentration of the shallow donor level is 6 x 10l6 cm-3 in sample 1,2.3 x lOI7 cm-3 in sample 2,7.4 x 1017cm-3 in sample 3, and 7.9 x 10'8cm-3 in sample 4. At higher donor concentrations the modeling cannot reproduce the experimental electron concentrations at lower sample temperatures (samples 3 and 4). This is due to the onset of hopping conduction between closely spaced donor levels and the presence of compensation. The measured activation energy for Si-donors is consistent with a donor level positioned -23meV below the conduction band edge as was predicted by PL (see Fig. 13), because high doping effects, such as coulomb interaction between neighboring donors, lead to a lowering of the activation energy as estimated by variable temperature Hall-effect measurements. The presence of Si in samples 2-5 was verified with secondary ion mass spectrometry. SIMS enables chemical identification and yields quantitative Si or Mg profiles when calibrated with implantation standards. The elemental Si concentration was found to agree within experimental uncertainties with the concentration of the shallowest donor for samples 2-4; for sample 5, the Si concentration was found to be -2 x 1019cm-3. For sample 1, Si

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 167

could not be detected with SIMS above the background level of the SIMS apparatus (- 1 x 10l7cm-3); however, its is likely that Si is the dominant donor because the activation energy (16 meV) is consistent with SiGa. The presence of a second donor level was assumed to model the temperature dependence of the electron concentrations for samples 1-3. The activation energy for this donor level ranges from 32 to 37eV and its concentration from 3 to 6 x 10'6cm-3, independent of the Si-doping. It is possible that this second donor level is due to oxygen substituting for nitrogen in the GaN lattice acting as a donor. In fact, samples intentionally doped with 0 show donor levels in this range. Results from Hall-effect measurements for 0-doped GaN (sample 6) are shown in Fig. 15. The electron concentration at 300 K is 4.4 x 1017cm-3 and the mobility is 314cm2V-'s-'. The dominant donor level in this sample has an activation energy of -30meV and a concentration of -9.8 x 1017cmW3.However, the 0 concentration, as determined by SIMS, is -8 x 1018cm-3. For 0-doped GaN, a one-to-one correlation between the atomic and the donor concentration has not been observed suggesting that not all the 0 atoms act as donors. As a consequence, to date, a clear identification of the donor level for oxygen in GaN has not been achieved.

loia

P I ' I

'

I

'

I

----I

OMVPE-grown GaN

sample#:

"7-

+ 6

E

s 0 c

5 8

5

Pw

++

10"

I

2

.

I

3

.

I

4

.

l

5

IOOOIT (1IK)

,

I

6

100 300 Temperature (K)

600

FIG. 15. (a) Electron concentration vs reciprocal temperature and (b) Hall electron mobility vs temperature for GaN contaminated with oxygen. The symbols refer to experimental data while the solid line in (a) is the result of a least-squares fit to the experimental data.

168

R. S. KERXet al.

b. P-Type Gallium Nitride The growth by OMVPE of highly conductive p-type GaN is difficult due predominantly to three factors. First, the relatively high background donor concentration means that all GaN.Mg films are compensated to a high degree (Chadi, 1997; Kennedy et al., 1995). Although advancements in the understanding of GaN crystal growth have led to parallel advances in p-type doping, very high compensation levels (generally > lo1*cm-,) are not uncommon. Second, the deep acceptor level exhibited by even the shallowest acceptor dopant, Mg, is still several kT above the valence band edge of GaN, meaning that only a small portion (calculated to be less than 1%) of the dopant atoms incorporated into the film are actually ionized at room temperature. The result is that Mg atoms must be incorporated at concentration levels one to two orders of magnitude higher than the required hole concentration, seriously challenging the solubility limits in the material. Finally, a Mg-H complex, which passivates the acceptors, forms due to the incorporation of interstitial H into the crystalline lattice (Gotz et ul., 1995a, 1996b; Nakamura et al., 1992a; Neugebauer and Van de Van de Walle, 1996b; Okamoto et al., 1996). This problem is very severe when OMVPE is used as the growth technique due to the source and carrier gases (particularly NH, and H2) which provide considerable activated hydrogen, but is not evident in plasma-assisted MBE due to the absence of these gases (Wang and Davis, 1993). Results from variable-temperature Hall-effect measurements for three Mg-doped GaN samples (nos. 7-9) are shown in Fig. 16. The OMVPE growth parameters for each of the samples were selected to achieve a range of Mg concentrations. The SIMS profiles revealed Mg concentrations of 1x ~ r n (sample - ~ 7), -5 x l O ” ~ m -(sample ~ 8), and 1020cm-3 (sample 9). For each of the Mg-doped samples, the p-type conductivity was achieved by thermal annealing. Figure 16a shows the hole concentration as a function of reciprocal temperature, and in Fig. 16b the hole mobility is shown as a function of temperature. Sample 9 exhibits the highest hole concentrations and the lowest hole mobilities. At 300 K, the hole concentration is -8 x 1 0 ” ~ m - ~and , the mobility is only -3cm2V-‘s-’ . The analysis of the temperature dependence of the hole concentration at temperatures above 280K using the charge neutrality condition (Gotz et al., 1996a) is consistent with the presence of a single acceptor with an activation energy of 160meV and a concentration of - 8 x 10’9cm-3 (solid line). Due to this high activation energy, only 1% of the acceptors are ionized. A good agreement exists between the concentration of Mg and the concentration of the dominant acceptor level; the latter determined by analysis of the Hall-effect data. This indicates that most of the Mg atoms are incorpor-

-

-

-

-

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 169 (a) 1oa

Temperature

500 1 . 1 .

300 I

(K)

200 '

150

I

OMVPE-gmwn GaN:Mg

1oT4

A 0

iot3

A

,

I ....

300 lOOOlT (1IK)

600

Temperature (K)

FIG. 16. (a) Hole concentration vs reciprocal temperature and (b) Hall hole mobility vs temperature for GaN doped with increasing Mg concentration. The symbols refer to experimental data while the solid lines in (a) results from least-squares fits to the experimental data.

ated on Ga lattice sites and act as acceptors. This is typically observed for Mg-doped GaN as long as the Mg concentration is less than 1 x 1020cm-3 (Gotz et al., 1995a). For higher Mg concentrations, the acceptor concentration decreases and the nitride materials become more resistive (Bour et al., 1997). The formation of Mg3N2 precipitates has been suggested to be responsible for the reduction of acceptor concentration when the solubility limit for Mg in GaN is exceeded (Neugebauer and Van de Walle, 1996~). For sample 9 at temperatures below 280 K, the temperature dependence of the hole concentration is not consistent with acceptor ionization acceptor (Fig. 16a). In this temperature range the experimental hole concentrations increase with decreasing temperature. This behavior, which is accompanied by a strong decrease of the hole mobility (Fig. 16b), is indicative of hopping conduction via closely spaced acceptor states. A similar behavior is observed for sample 8. The temperature dependence of the hole concentration can be modeled only for temperatures above -250 K. For sample 7, however, the temperature dependence of the hole concentration can be described over the entire temperature range of the Hall-effect data (Fig. 16a). The analysis

-

170

R. S. KERNet

ul.

reveals that the acceptor concentrations of samples 7 and 8 are lower than in sample 9. For samples 7 and 8 the acceptor concentrations were estimated to be -2 x 1019cm-3 and -4 x 10’9cmp3, respectively. The acceptor activation energies for acceptors in samples 7 and 8 were found to be rv 175 and 170meV, respectively. As in the case for Si-doped GaN higher activation energies are usually obtained from samples with lower dopant concentration (Schubert, 1993). To take the effect of high acceptor concentration into account, we assume that the lowering of the activation energy is proportional to the distance between ionized acceptors ( -(N;)ll3). Such models have been suggested for donors and acceptors in Si (Pearson and Bardeen, 1949), Ge (Debye and Conwell, 1954), and GaAs (Stillman et ul., 1982). Lowering of the activation energy may occur due to overlap of the coulomb potentials of the dopants and screening of the coulomb potentials by a high concentration of free carriers. For the case of Mg-doped GaN, Eq. (7) must be modified according to

[ + ypp (

p =N, 1

NV,e‘f

exp

)]

AEA(NA) kT

p1

- Ncomp

where

As compared to Eq. (7), Eqs. (8) and (9) are written for a p-type semiconductor (hole concentration, p), a single acceptor (concentration, NA), and the activation energy is a function of the ionized acceptors concentration, AE,(N, ); AEA.0 is the activation energy for a very low acceptor concentration, E, is the dielectric constant for GaN ( = 9c0), and f is a geometric factor and describes the probability of finding another ionized acceptor in the neighborhood of an ionized acceptor (given by f(2/3)(4n/3)”3; Schoner, 1994). With Eq. (8), the temperature dependence of the hole concentrations for a large number of Mg-doped GaN samples with acceptor concentrations in the range between 10l8c m V 3and lo2’ was analyzed. A concentration-independent activation energy of 208 k 6 meV was consistently found. This activation energy, AE,,, is in excellent agreement with results from low-temperature photoluminescence measurements for the position of the Mg acceptor level in the bandgap of GaN (208 & 4 meV, Fig. 13). A good agreement between the “thermal” activation energy as determined from electrical measurements and the position of the energy level in the bandgap is expected for the case of very low doping concentrations.

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES171

Also, the observation of hopping conduction in samples 8 and 9 at lower temperatures is consistent with the presence of higher acceptor concentrations in these samples. Hopping conduction is also indicative of the presence of compensation. For samples 7, 8, and 9 the concentration of donors was determined to be -2 x 101'cm-3, -2 x lo1*cm-,, and -5 x 10" ern-,, respectively. Possible candidates for compensating donors are the nitrogen vacancy (V,) and substitutional oxygen ( O N ) ,the latter exhibiting a strong affinity to Mg, as well as isolated hydrogen which acts as a donor in p-type GaN (Neugebauer and Van de Walle, 1995). However, the hole concentration is only marginally affected by the presence of compensation since the acceptor concentration in Mg-doped GaN is typically more than one order of magnitude higher than that of compensating donors. The hole mobility is adversely affected by both the presence of compensating donors and high concentrations of acceptors. The 300-K hole 15 cm2 V - l s - ' and 6 cmz V-'s-', mobility for samples 7 and 8 are respectively. The 300-K resistivity for all three Mg-doped GaN films depicted in Fig. 16 was measured to be 3 Qcm. Due to this relatively high resistivity the players in a nitride LED are kept thin (< 1 pm) to avoid impact on the forward voltage of GaN light emitters. Similar to other elemental or compound semiconductors, hydrogen forms electrically inactive complexes with acceptors in GaN (Pankove and Johnson, 1991). Hydrogen is abundant in the OMVPE reactor chamber from H,, which is used as a carrier gas, and also as a cracking product from the N precursor, NH,. It has been shown that H acts as a donor (H') in Mg-doped GaN (Neugebauer and Van de Walle, 1995) and is highly mobile (Gotz et al., 1995a). The situation is different for n-type GaN where H is an acceptor (H-) and immobile (Gotz et al., 1995a). Due to coulombic attraction, H and Mg;, form complexes after growth during cool-down and, as a result, as-grown Mg-doped GaN is highly resistive (Neugebauer and Van de Walle, 1996b). However, annealing of GaN:Mg in a hydrogenfree ambient ( e g , N,) will activate the p-type conductivity by dissociating the Mg-H complexes (Gotz et al., 1995b; Nakamura et al., 1992a). This phenomenon was first discovered by Amano et al. (1989), who irradiated as-grown Mg-doped GaN samples with low-energy electrons (low-energy electron-beam irradiation, LEEBI treatment) and discovered that the material became p-type conductive. Experimental evidence for the existence of Mg-H complexes in as-grown Mg-doped GaN is presented in Fig. 17. Figure 17a shows low temperature infrared absorption spectra for (1) a 4-pm-thick Mg-doped GaN film ([Mg] 5 x lo1' cm-,) in the as-grown state, (2) after thermal activation, and ( 3 ) after exposure to elemental deuterium in a deuterium plasma at 600°C for 2 h (Gotz et al., 1996b).The infrared spectra exhibit an absorption

-

+

-

-

172

R. S. KERNet ul. (a)

Energy (mew

(b) Temperature (K) 500 300 200 150

285 290 295 380 385 390

m

7-1

100

10"

Ion

y; (1) asgrown

E

I0'

G

f '0'

i!

103

10'

Ld 2280 2320 2360

3080 3120 3160 Wavenumber (cm.')

,o-l

(2)activated (3) (3)deuterated

u 2 3 4 5 6 7 8 9 10 IOOO/T (1IK)

FIG. 17. (a) Infrared absorption spectra and (b) electrical resistivity vs reciprocal temperature for Mg-doped GaN. Data are shown for as-grown material (l), after RTA activation of the p-type conductivity (2), and after deuteration (3). The vertical bar in (a) indicates the magnitude of the absorbance scale.

line at 3125cm-' (right portion of Fig. 17a) which is due to a local vibrational (stretch) mode (LVM) indicative of Mg-H complexes in GaN (1). The peak height in the absorption spectra is proportional to the concentration of the complex in the sample and is highest for the as-grown sample. After activation, the peak height is reduced to 1/3 of the original peak height indicating the dissociation of a large quantity of Mg-H complexes (2). After deuteration, a peak at 2321cm-' appears in the infrared spectrum (left portion of Fig. 17a) that is due to a stretch mode of the Mg-D complex (3). The observation of the isotopic shift of the vibrational frequency clearly establishes the involvement of H. The isotopic ratio of the two absorption line frequencies is -1.346, close to $,as observed for H-related LVMs in other semiconductors (Chevallier et al., 1991; Stavola and Pearton, 1991). Figure 17b shows the resistivity as a function of reciprocal temperature for a Mg-doped GaN sample that underwent treatments similar to those for the sample used to record the infrared absorption spectra in Fig. 17a. Initially, in the as-grown state, the sample is highly resistive and the resistivity can only be measured at high sample temperatures (1). After

-

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES173

thermal activation and the achievement of p-type conductivity, the resistivity of the sample decreases by more than seven orders of magnitude (2). After deuteration, the resistivity of the sample increases by about five orders of magnitude (3). The data presented in Fig. 17b is consistent with the model that the presence of Mg-H complexes is responsible for the electrically insulating nature of as-grown Mg-doped GaN. The intensity of the absorption line at 3125 cm- decreases significantly on activation of the p-type dopant. This observation is consistent with the model that after growth most of the Mg acceptors are passivated by hydrogen. The Mg-H complexes are electrically inactive and give rise to the absorption line observed in the infrared spectrum. In the as-grown material, only a small fraction of the Mg atoms is required to act as acceptors to compensate the residual donors (Neugebauer and Van de Walle, 1996b). As a consequence, as-grown Mg-doped GaN is highly resistive. The postgrowth activation treatment results in the dissociation of the Mg-H complexes. As a consequence, the intensity of the LVM line at 3125 cm-' decreases (Fig. 17a). From variabletemperature Hall-effect measurements, the concentration of acceptors is estimated to be -2 x 10'9cm-3 at this stage. During the activation process, the hydrogen atoms are likely to either evolve out of the sample or to be immobilized at internal sites, such as grain boundaries or dislocations, where they cease to influence the p-type conductivity. It was suggested that the presence of hydrogen during OMVPE-growth is beneficial for Mg-doping of GaN (Neugebauer and Van de Walle, 1996b). A combination of acceptors and donors is an energetically favorable condition for crystal formation in wide-bandgap semiconductors under thermodynamic equilibrium (Vechten et al., 1992). Among potential donors that could form during growth, such as the nitrogen vacancy (V,) and interstitial or substitutional Mg (Mg, or Mg,), hydrogen (H') has the lowest formation energy. The formation energies for these donors and for Mg acceptors in Mg-doped GaN at OMVPE-growth temperature were computed from first principles calculations (Neugebauer and Van de Walle, 1996b) and are plotted as functions of the Fermi level in Fig. 18. The formation energies relate to the equilibrium concentrations of each defect. A higher (lower) formation energy will translate into a lower (higher) equilibrium concentration. The formation energy for donors increases whereas the formation energy for Mg,, decreases with larger distance of the Fermi level from the valence band edge. Figure 18 demonstrates that the formation energies for Mgi or Mg, are high and are not likely to play an important role as compensating donors in Mg-doped GaN. At growth temperature, the Fermi level is positioned near the crossover point between the line for Mg& and the donor with the lowest formation energy. At this

174

R.S. KERNet al.

FIG. 18. Formation energies for electrically active point defects vs position of the Fermi level as derived from first principles total energy computations. The Fermi level is normalized to the valence band maximum. Results are shown for Mg substituting for Ga (MgGa),for N (Mg,), interstitial Mg (Mg,), the nitrogen vacancy (VN), and isolated hydrogen. Charged donors and acceptors exhibit a positive and a negative, respectively. A change in slope indicates a change in charge state of the respective defect. (Data provided by C. G . Van de Walle, Xerox Palo Alto Research Center, Palo Alto, California 94304.)

point, the formation energies for Mg,, and Hi are equal and, as a consequence, their concentrations are equal. In the presence of hydrogen, the crossover point between the lines for MgG, and H is located 1.2 eV above the valence band edge. In the absence of hydrogen, the donor with the lowest formation energy is the nitrogen vacancy. Under this condition the Fermi level is positioned closer to the valence band (- 0.8 eV above the valence band edge). As a consequence, the presence of hydrogen leads to an increased concentration of the acceptor dopant (Mg,,) and to a reduced concentration of compensating donors, such as V, (Neugebauer and Van de Walle, 1996b). After growth, the H atoms bond with Mg,, forming Mg-H complexes that are subsequently dissociated by postgrowth activation treatments. +

N

4. GALLIUM INDIUM NITRIDE At the heart of all visible AlGaTnN optoelectronic devices is GaInN. The addition of In to GaN to form GaInN alloys lowers the bandgap energy

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 175

-

and makes Ga,In, - x N (x < 0.8) suitable as the active region in LEDs or LDs to achieve light emission in the visible range of the electromagnetic spectrum. As a consequence, the role of native defect or impurities that influence the recombination of electrons and holes injected into the GaInN region from the n-and p-side of the diodes is critical for their performance. Composition and bandgap are related through the simple relation

where E,(x) represents the bandgap energy of Ga,In,-,N, Eg,InN, and Eg,GaNrepresent the bandgap energies of InN (1.95 eV) and GaN (3.4 eV), respectively, and b is a bowing parameter. The value of b has been the source of some controversy with values of 1.0eV (Nakamura et al., 1993c), 2.6eV (Wright and Nelson, 1995), and 3.2eV (Takeuchi et al., 1997b) reported. It was suggested that for strained GaInN layers the bowing parameter of Eq. (10) actually is a function of the composition (i.e., b(x)) and for x = 0.1 is equal to 3.8 eV (McCluskey et al., 1998). The first commercial nitride LEDs (Nakamura et al., 1994a) utilized a codoped In,,,Ga,,,N:Si,Zn region in a DH structure to shift the light emission to the blue range because the growth of high-quality GaInN with an In composition larger than 10% was not possible at that time. Similar to the situation for GaN, Si is a shallow donor in GaInN; however, its electronic properties have not been characterized by Hall-effect measurements. Nakamura et al. (1993b) reported that Si-doping significantly enhances the near bandgap luminescence efficiency of GaInN at room temperature. Zn is a deep acceptor in GaInN and dominates the radiative recombination under low-injection conditions. Similar to GaN (Pankove et al., 1974), Zn substituting for the group I11 metal is believed to introduce an acceptor level -0.4 eV above the valence band edge (Nakamura, 1994b). Cd-doping was reported to shift the emission wavelength -0.5 eV below the bandgap of GaInN; however, no recombination model has been proposed to explain the experimental observations (Nakamura et al., 1993a). Yamasaki et al., (1995) achieved p-type conductive GaInN by doping with Mg during OMVPE growth. After thermal activation, a room temperature hole concentration of - 7 x 1017cm-3 was measured.

5. ALUMINUMGALLIUM NITRIDE

-

Doped AlGaN layers with A1 compositions of up to 15% are utilized in both nitride LEDs and LDs. For commercial SQW LEDs (Nakamura et al., 1995a), Mg-doped AlGaN is employed as an electron confining layer on

176

R. S. KERNet al.

top of the GaInN active region. As a consequence, control of the doping is critical in this layer due to its proximity to the active region. In LDs, both n- and p-type AlGaN layers are incorporated as cladding layers (Nakamura et al., 1996). Low resistivity is required for these layers. The n-type conductivity is accomplished by doping with Si, which, as in the case of GaN and GaInN, is a shallow donor. Undoped AlGaN is also n-type most likely due to the incorporation of oxygen. However, AlGaN with an A1 composition > -40% is highly resistive (Bremser et al., 1996). It has been suggested that oxygen exhibits a DX-like behavior (Stampfl and Van de Walle, 1998) and changes its nature from a shallow to a deep donor at x = 0.4 leading to a freeze-out of free electrons. In the case of Si-doping, compensation by aluminum vacancies (VAl), an acceptor in AlGaN, was suggested to be responsible for the high resistivity of AlGaN with high A1 compositions ( >40%). P-type doping with Mg has been reported in the literature for A1 compositions up to -13% (Bremser et al., 1996). In Fig. 19 we present results from variable temperature Hall-effect measurements for p-type A1GaN:Mg containing 17%. The p-type conductivity was activated by thermal annealing. The hole concentration as a function of inverse tempera-

-

(a)

500 lom

Temperature (K) 300 200 I

1 . 1 .

.

150 10’

1

OMVPE-grown AIa,,7Ga,,,N:Mg

lo‘* “7-

5

1ol8

h

2

Y

5

6

10”

Y

2

zP

2.

; s d 10l5 IOl6

I

loa

10“

3

4

5

1000/T (l/K)

6

7

10

100

300

600

Temperature (K)

FIG. 19. (a) Hole concentration vs reciprocal temperature and (b) Hall hole mobility vs temperature for Mg-doped Al,,,,Ga,,,N. The symbols refer to experimental data while the solid lines in (a) results from least-squares fits to the experimental data.

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 177

ture is shown in Fig. 19a and the hole mobility as function of temperature in Fig. 19b. For this sample, the A1 composition was determined by roomtemperature cathodoluminescence measurements (Khan et al., 1983; Pankove et al., 1970). After thermal activation, the Mg-doped sample exhibits p-type conductivity (300 K: p z 21 R . cm). The 300-K hole concentration and mobility are 1 x lo’’ cm-3 and -3 cm2 V-’ sC1, respectively. The analysis of the temperature dependence of the hole concentration for sample 10 reveals the presence of a single acceptor. The low-concentration activation energy for ionization (AEA,O) was estimated to be 260 meV and the concentration of acceptors was found to be 1.1 x 1020cm-3.The concentration of donors in this sample was estimated to be 1 x lOI9 ~ m - ~ . These results indicate that the acceptor level for Mg-doped AlGaN shifts deeper into the bandgap with increasing A1 composition. In fact, from a study of a series of A1GaN:Mg samples with a range of AlN compositions (0 to 17%) (Gotz et al., to be published), an increase of the acceptor level of 3.2meV per percent added A1N composition was observed. As a consequence higher acceptor concentrations are needed for AlGaN than for GaN to accomplish the same hole concentration at room temperature. In the literature, Hall-effect measurements for p-type Mg-doped AlGaN are reported for A1 compositions up to 8% (Tanaka et al., 1994) where a room-temperature hole concentration of 3 x 10l6cm-3 was achieved. The activation energy for the Mg acceptors in this material was found to be 192 meV for an acceptor concentration of 1.5 x lo1’ ~ m - ~ .

-

-

-

-

-

IV. Nitride LED Device Structures and Performance Since the first demonstration of high-efficiency nitride LEDs by Nichia Chemical Industries, Ltd. (Nakamura et al., 1994a, 1994b), other groups have also demonstrated bright III-V nitride devices. In the following sections, a general discussion of the various device structures is presented, followed by a review of design considerations for some of the individual device layers and the resulting light output performance. This section concludes with a discussion of the reliability characteristics of III-V nitride devices and the use of nitride devices for the production of white LEDs. 1. LED STRUCTURES There are numerous descriptions of III-V nitride device structures in the literature. The original device structures reported by Nichia have evolved considerably, and other groups have developed variations of their own. The

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device structure shown in Fig. 20a has been reported to produce superbright LEDs when the active layer is a single quantum well (Nakamura et al., 199%). A 3-D schematic of a typical LED chip with an electrically insulating sapphire substrate is shown in Fig. 20b. The schematic in Fig. 20 is not meant to include every layer in an LED device structure but rather only the most important layers necessary for creating the high-performance active region. Additional layers can also be important, but they are related more to obtaining high material quality. For example, the low-temperature buffer layer between the GaN:Si and the substrate is critical for alleviating the large lattice and thermal mismatches between the GaN epilayers and the sapphire substrate (Akasaki et al., 1989), as discussed in detail in Section I1 of this chapter. Another example is the low-temperature cap layer following the growth of GaInN active layer, important for preventing decomposition of the fragile GaInN active layer (Shmagin et al., 1997). The structure of Fig. 20 consists of the basic layers in a typical double heterostructure LED: the Si-doped n-type GaN lower confining layer, the active layer, the Mg-doped p-type AlGaN upper confining layer, and finally the Mg-doped GaN contact layer.

GaN:Mg AlGaN:Mg M e layer GaN:Si

substrate

FIG. 20. (a) Schematic diagram of a nitride device structure. The schematic shows only the most important layers necessary for device operation. (b) Schematic of a device structure after fabrication showing the mesa etching required to expose the n-GaN layer for n-electrode formation.

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 179

Of the variations in device structure reported in the literature, the most notable is the presence of an AlGaNSi layer below the active layer, reported by Nakamura et al. (1994a) in their early devices. In later LED structures, the A1GaN:Si layer was omitted without detriment to device performance (Nakamura et al., 1995~).The reason for this is perhaps that the hole effective mass in 111-V nitride materials is high enough to prevent the injected minority carriers (holes in this case) from escaping the active layer. Thus, GaN:Si appears to be adequate as a confinement layer for holes in 111-V nitride LED devices. Another notable modification of the device structure concerns a 500-A Ino,,,Ga0~,,N shoulder layer deposited before the GaInN active layer and intended to reduce the strain produced by the difference in lattice spacing between GaN and GaInN. This layer is first reported in the SQW Nichia devices (Nakamura et al., 1995b); however, it was later found that brighter devices could be obtained without including this shoulder layer (Nakamura et al., 1995~).

SELECTION 2. SUBSTRATE There are several substrate choices for 111-V nitride epitaxial growth, including sapphire (A1203), silicon carbide (Sic), spinel (MgAl,O,), LiAlO,, LiGaO,, and silicon. The impact on crystal growth by the substrate is discussed in Section I1 of this chapter. For device applications, the important considerations are conducting vs nonconducting substrates and transparent vs nontransparent substrates. An electrically conductive substrate, such as Sic, enables the construction of a vertical device geometry, similar to other conventional 111-V LEDs grown on GaAs and GaP substrates, where the current flows predominantly vertically through the chip. Figure 21a shows a schematic diagram for a vertical device mounted on a leadframe. The n-contact is a sheet of metal (shown in black) covering the bottom of the device; the p-contact is typically a circular dot (in black) on the surface of the chip. The processing steps needed to fabricate such a chip are straightforward with only one masking step required to define the top contact. Since the bottom n-contact is in electrical contact with the metal cup of the leadframe through a conductive die-attach epoxy, only one wire bond is needed to the top p-contact. For a nonconducting substrate, such as sapphire, the device structure is illustrated in Fig. 21b, and here the current flow is predominantly horizontal between the a- and p-contacts. (The vertical scale of the step between the contacts is greatly exaggerated in the figure.) Because both n- and pcontacts are made to the frontside of the chip, it is necessary to etch through the top p-layer to deposit the n-contact metal, and the processing required

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FIG. 21. Schematic diagrams of two types of nitride chips die-attached and wire-bonded in kddframes. A “traditional” vertical device geometry is shown in (a) where the substrate is electrically conducting. (h) The substrate is electrically insulating and the device geometry is “horizontal,” because the direction of current flow is mainly horizontal. Note that two wire bonds are required in the configuration in (b).

to accomplish this is more complicated than for the vertical chip structure. In addition, two wire bonds are needed to make the electrical connection to the chip. Light emission from the active region of an LED is directionally isotropic and can escape from the chip if the angle of incidence to the chip wall is less than the critical angle, 8,, defined by the index of refraction of the chip and the surrounding medium (air, epoxy, etc.). In the case of a typical cubeshaped chip, the light coming from any point in the active region that satisfies this condition defines six cone-shaped regions, each with an apex of 2 8, and directed toward a cube face. When the substrate is nontransparent, the light in the downward-directed cone is completely absorbed. Similarly, the bottom half of each of the side cones is also absorbed by the substrate (Huang et al., 1992). In comparison, when the substrate is transparent, the light within all six cones has a chance to escape from the chip, essentially doubling the light output. Thus, transparent substrates are preferred for high-efficiency devices.

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 181

Sapphire is the most commonly used substrate for 111-V nitride growth. The orientation of the sapphire is usually (OOOl), or c-plane, although (1 120) a-plane sapphire is also used. From a growth point of view, the advantages of using sapphire are that sapphire has high crystalline quality, is widely available in large quantities, and is fairly inexpensive. Sapphire is also transparent to U V and visible light, making it ideal for high efficiency LEDs. There are several disadvantages to using sapphire as a substrate, however. First, sapphire is electrically insulating, thus making it necessary to fabricate devices in the horizontal configuration, as shown in Fig. 21b. Second, sapphire has relatively poor thermal conductivity, slightly worse than GaAs. Thus, heat dissipation could be a problem for high-power devices such as lasers. In typical LED applications, however, heat dissipation has not been a significant problem because 111-V nitride devices have a high To,a result of the large bandgap offset between active layer and the carrier confining layers. Third, sapphire lacks the orthogonal natural cleavage planes necessary to form a cavity for edge-emitting lasers. Fortunately, resonant cavities are not needed for LED applications. Finally, the extreme hardness of sapphire can make die fabrication difficult. Sic has also been used as a substrate for 111-V nitride growth (see Section 11). First, from a device point of view, the major advantage of S i c as compared to sapphire is that it can be easily doped either n- or p-type to make a conductive substrate. This enables the device structure to be vertical, as shown in Fig. 21a. Second, Sic has a high thermal conductivityapproximately 10 times better than sapphire. (In fact, Sic exhibits a higher room-temperature thermal conductivity than any metal.) This property helps reduce heating effects in an LED or LD device. Finally, Sic possesses the natural orthogonal cleavage planes necessary for making facets to form the resonant cavity in a laser structure. Sic is more expensive than sapphire, however; and the use of Sic is much less common. This situation may change when more Sic vendors become available. Another disadvantage is that Sic may not be entirely transparent. Although the bandgap energy for 6H-Sic is high enough for it to be transparent to visible light, it contains midgap levels; and doped Sic can even be absorbing at the wavelengths of blue light. Finally, and most important, the LEDs produced on conductive 6H-Sic to date are less efficient than those made on sapphire substrates. Similar to the growth of GaN on sapphire, the growth of GaN on S i c requires a buffer layer (see Section I1 for details). An AlN buffer layer works well, but since AlN is an insulator and is impossible to dope n- or p-type, the advantage of the conductive Sic is almost lost. To circumvent this difficulty, Cree Research developed a chip fabrication process in which they deposited a metal shorting band around the edges of the AlN buffer layer

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to connect the conductive Sic substrate to the conductive GaN:Si layer (Compound-semiconductor, 1998). More recently, Cree has replaced the insulating A1N with a conducting A1GaN:Si alloy as the buffer layer. This configuration has recently become available in commercial products (‘“itride News,” 1998).

3. DEVICE DESIGNCONSIDERATIONS FOR TYPE LAYERS To date, all III-V nitride devices reported in the literature are n-sidedown devices where the initial layer deposited is n-GaN, as shown in Fig. 20. The doping characteristics in OMVPE have been discussed in Section 11, and the electronic behavior of n-GaN has been discussed in Section I11 of this chapter. For the horizontal device shown in Fig. 21b, the conductivity of the n-GaN layer is very important. Figure 22 shows a simplified view of the current paths for a horizontal device. The vertical scale in the figure is greatly enhanced compared to the horizontal scale to show the details of the current streamlines. Figure 22 shows that the current streamlines are primarily vertical in the players and horizontal in the n-GaN layer. The player is more resistive than the semitransparent metal contact on top and the n-layer underneath; thus, the holes will seek the path of least resistance Thin semi-transparent metal Anode bond wire

/

GaN:Mg /

Cathode bond wire

Sapphire substrate

FIG. 22. Schematic diagram of the current flow streamlines for a horizontal device with a conducting p-metal contact layer on the top surface of the device. The scale of the layer thicknesses is exaggerated in this figure.

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by traveling vertically through the p-layer. Furthermore, since the horizontal dimension of the epitaxial layers is of the order of several hundred micrometers, whereas the vertical thickness is only a few micrometers, the electrons have to travel a long distance along the horizontal direction in the n-GaN before making a turn to move upward. Thus, the current streamlines are mainly horizontal in the n-GaN layer. Figure 22 shows the ideal situation when the GaN:Si sheet resistance is comparable to that of the top semitransparent p-metal, making the vertical current streamlines uniformly spaced across the device. If the GaNSi layer conductivity is much larger than that of the semitransparent metal layer, current in the n-layer will crowd toward the p-contact pad before traveling vertically upward, leading to current crowding near the p-layer bonding pad. The opposite situation occurs if the GaN:Si layer conductivity is much less than the p-contact metal layer. In this case, current will flow across the p-metal contact layer before moving downward to the n-GaN layer, leading to current crowding near the n-side bonding pad. The resistance contribution from the GaN:Si layer is important in 111-V nitride devices because the electrons have to travel a long distance along the horizontal direction in a thin layer. The situation is analogous to that encountered in the base region of a transistor where the electrons also have to flow horizontally through a thin layer. For the purpose of lowering the resistance and the forward voltage of devices, it is necessary to grow a thick GaN:Si layer with a high Si doping. However, it is well known that cracking occurs when a thick GaN:Si layer is doped to high concentrations (Murakami et al., 1991). Typically, the Si doping level is maintained from 1 to 5 x lo’* cm-3 to avoid cracking in the epilayers. At these doping levels, the GaN:Si series resistance can be a significant contributor to the device operating voltage. In addition, the low Si doping level can also lead to nonohmic behavior at the n-metal contact. There are several ways to resolve the conflict between cracking and the need for thick and highly conducting n-GaN. One solution is to use hydride vapor phase epitaxy to grow thick GaN:Si without cracking (Naniwae et al., 1990). This thick n-type layer acts as a conducting “substrate” for subsequent OMVPE growth. Another engineered solution is to use a three-layer structure on the n-side of the device (Rudaz, 1998), schematically shown in Fig. 23. In this design, the doping in the 3rd layer, n3, is optimized for current injection into the active layer. The doping in the second layer, n2, is kept high to allow the formation of a good n-contact (during processing, layer n2 is exposed by etching to form the n-contact). Finally, the doping and thickness of the first layer, n,, are maximized to achieve high conductivity without the occurrence of cracking.

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Active layer 3rd GaN:Si with doping n, 2nd GaN:Si with doping n2 1st GaN:Si with doping

Sapphire substrate

FIG.23. Design of a GaN structure using three layers of GaN:Si to achieve maximum doping for high-conductivity and low-resistivity n-contact formation while avoiding cracking of the epitaxial layers due to stress.

DESIGN CONSIDERATIONS FOR p-TYPELAYERS 4. DEVICE High p-type doping has been a vital key for obtaining a high-efficiency pn-junction device and has been one of the major stumbling blocks in the development of high brightness GaN devices. After the development of the low temperature buffer layer, which greatly improved the material quality, Amano et al. (Amano et al., 1988) were still forced to make GaN light emitting devices using metal-insulator semiconductor (MIS) device structures before the techniques for achieving high-p-type doping were discovered. The performance in the MIS devices was understandably poor due to the inherent low carrier injection efficiency. High performance LEDs were not realized until p-type doping was demonstrated (see Sections TI and I l l of this chapter). Typically, the GaInN active layer is n-type in the range of 1 x loL7 cm-3 in 111-V nitride devices (Nakamura et al., 1993~).Minority carriers (holes) must be injected into the active layer from the AlGaN layer. It is extremely important that this layer be p-type, otherwise the holes from the p-type GaN:Mg must tunnel through an n-type AlGaN to get into the GaInN active layer. Tunneling leads to poor hole injection efficiency and low light output efficiency. Because it is difficult to dope AlGaN alloys p-type when the A1N component is high (see Sections I1 and III), one must select the AlN composition carefully. In most cases, the A1N solid composition is about 15 to 20%. Also, AlGaN layers are prone to cracking above a certain thickness (Nakamura et al., 1998). Thus, most device structures use a relatively thin AlGaN layer, typically about 200 to 1000 A in thickness. It is instructive to consider the influence of the players on the device forward voltage (Vfj. According to Fig. 22 and the discussion in Section IV, iubsection 3, current moves vertically in the players. Thus, the calculation

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES185

of the resistance R of the player follows the expression

where p is the resistivity, d is the thickness, and A is the area of the player. For p-GaN, typical values are 1 R . cm, 0.5 pm, and 300 x 300 pm, respectively. These values yield a resistance of 0.0560, giving rise to a miniscule voltage drop of approximately 0.001 V at 20 mA forward current. For p-AlGaN, typical values are 10 R * cm, 0.1 pm, and 300 x 300 pm, respectively, leading to a resistance of 0.1 1 R and a correspondingly low contribution of approximately 0.002V to the V, at 20mA. Thus, the effect of series resistance from the player on device V’ can be ignored in most cases. However, the player can still have an impact on V, through its influence on the p-contact resistance and possible formation of a metal-semiconductor potential barrier. Device structures typically incorporate a top p-type contact layer to improve ohmic contact behavior (Nakamura et al., 1994a), usually achieved through high Mg doping (Sassa et al., 1998).

5. DESIGN OF ACTIVELAYERS The design of the active layer where light is generated is of paramount importance for obtaining high-quantum efficiency. It is well known that the most efficient LED and LD devices require double heterostructures to obtain both high current injection efficiency and high carrier confinement (Kressel and Butler, 1977). Thus, researchers designed the first 111-V nitride devices with the heterostructure AIGaN-GaN-AlGaN (Akasaki and Amano, 1992), similar to the heterostructures in the AlGaAs and AlGaInP material systems. However, GaN has an energy bandgap of 3.39 eV at room temperature, corresponding to a UV emission wavelength of 366 nm. For visible light emission, it is necessary to use a GaInN ternary alloy that has a lower bandgap energy than GaN as the active region. Figure 24 shows three designs for the active layer of a 111-V nitride device. The codoped active layer of Fig. 24a was used in the first generation of commercial 111-V nitride devices (Nakamura et al., 1994a, 1994b). At the time when these devices were produced, it was difficult to grow high-quality GaInN with the high-In composition required for blue light emission. However, by doping a low-In-composition active layer with a deep-level impurity a longer emission wavelength was achieved. Zn incorporated into the active layer accomplished this purpose (Nakamura et al., 1994a, 1994b). As a deep acceptor impurity with an ionization energy of approximately 0.4-0.5 eV in GaInN, Zn effectively turned violet-emitting GaInN into

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Upper confaring layer

1

Lower confining layer FIG. 24. Schematics of active layers in nitride device structures: (a) a codoped structure. (b) a single-quantum-well structure, and (c) a multi-quantum-well structure.

blue-emitting GaInN without the difficulty of increasing the In composition. Furthermore, by codoping with Si (Nakamura et al., 1994a, 1994b), the luminescence is enhanced by as much as 36 times as compared to undoped GaInN layers (Nakamura et at., 1993b). The optimal doping level is approximately 1 x 10’’ ~ m - This ~ . high doping level is necessary to enable the formation of a sufficient quantity of donor-acceptor (D-A) pairs for bright photoluminescence and efficient device light output (Nakamura et al., 1994b). The light emission occurs when an electron in a neutral donor recombines with a hole in a neutral, but spatially separated, acceptor. It is well known that the D-A recombination depends on the distance of the donor and acceptor impurities. When taking the coulombic interaction into account (Dean, 1982), the emitted photon energy is approximated by hv = E, - E, - Ed

+ e2/(.d)

(12)

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where E, is the bandgap energy of the GaInN, E , is the acceptor ionization energy (about 400-500meV for Zn), Ed is the donor ionization energy (approximately 20meV for Si), e is the electron charge, E is the dielectric constant of GaInN, and r is the distance between the donor and the acceptor impurities. Using the codoping method, the first “candela-class” high-efficiency blue LEDs were released in late 1993 by Nichia Chemical Industries, Ltd., and reported in early 1994 by Nakamura et al. (1994a). Output power as high as 1.5 mW was obtained, corresponding to an external quantum efficiency of 2.7%. This work was extended to blue-green 500-nm LEDs (Nakamura et al., 1994b), the longer wavelength achieved by reducing the growth temperature to incorporate more In into the active layer. At this wavelength, the output power was as high as 1.0 mW, corresponding to an external quantum efficiency of 2.4%. One problem with codoped devices is that the emission wavelength shifts dramatically with increasing forward current. In this situation, the electron and hole densities increase, leading to a reduction in the distance between the electron and hole pairs and a higher emission energy, as described by Eq. (12). This shift is graphically illustrated by the series of electroluminescence spectra shown in Fig. 25a. The magnitude of the shift is plotted in Fig. 25b. For this codoped device, the emission color shifts from green at 0.5 mA to blue-green at 20 mA. Another problem characteristic of codoped devices at high current is the saturation of the D-A pair luminescence and the appearance of higher energy band-to-band luminescence. When the drive current is sufficient for the electron concentration to exceed the donor concentration (- lo1’ cmT3),the lack of available donor states keeps the extra electrons in the conduction band. The similar situation obtains for holes in the valence band. As a result, D-A pair luminescence saturates at high drive current (Nakamura et al., 1994b), and conduction band to valence band recombination increases (Egawa et al., 1997). This phenomenon is clearly seen in the 50-mA luminescence spectrum in Fig. 25a. The appearance of the shorter wavelength peak is dependent on the saturation of doping levels in the active layer; the peak position is dependent on the solid composition of the GaInN alloy. This UV peak is detrimental for the reliability of the III-V nitride devices and will be discussed later. A further disadvantage of codoped devices is the broad emission peak. There is a distribution in the distance between the D-A pairs in Eq. 12, which leads to a distribution of recombination energies. The electroluminescence spectrum is a composite of this D-A pair energy distribution. The resulting color of the emission is not pure and appears whitish-blue to the human eye. Due to these shortcomings, most institutions have abandoned the codoped active layer approach and have instead focused on SQW or MQW

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300

-5

400

700

500 600 Wavelength (nm)

470

460

I a d 450 440 I

10

I1111111

100

I

I1111111

I

I1111111

I

1000 10000 Drive Current (PA)

I

llLu

100000

FIG.25. (a) Electroluminescence spectra for a codoped LED at five different levels of forward drive current. The drive currents are, starting with the bottom curve, SOpA, 100pA, 1 niA, lOmA, and SOmA. (b) Electroluminescence peak wavelength as a function of drive current for a typical codoped LED.

active layers. The SQW device is schematically shown in Fig. 24b. The key parameter for quantum well devices is to keep the GaInN layers thin since thick GaInN layers have poor material quality due to the large lattice mismatch between GaN and GaInN and because GaInN is thermodynamically unstable. Quantum well layers of the order of 30A in thickness lead to quantum-size effects and a reduction in the emission wavelength. Typical well thickness is 20-30A for an GaInN active layer with high indium solid composition (Nakamura et al., 1995~). The first SQW high-brightness blue and green LEDs were reported in 1995 (Nakamura et a!., 1995a, 1995b, 199%) with much better efficiency and

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES 189

color purity than codoped devices. Reported power output for 450-nm blue and 520-nm green LEDs are 4.8 and 3.0 mW, respectively, corresponding to external quantum efficiencies of 9.1 and 6.3%, respectively (Nakamura et al., 1995b, 199%). Nichia (Nakamura, 1999) further improved the LED efficiency significantly and reported 14% external quantum efficiency for blue devices and 11.6% efficiency for 530-nm green devices. In addition to improved light output, SQW LEDs also possess other favorable characteristics as compared to codoped LEDs. For codoped 111-V nitride LEDs, the FWHM of the electroluminescence peak is approximately 70 nm for the blue and 90 nm for the blue-green wavelengths. The FWHM for SQW LEDs is approximately 20nm for blue and 30nm for green, which, as shown in Fig. 26, is a more saturated pure-color emission than that of codoped LEDs (Nakamura et al., 1995b). However, the

loo 100

I +- co-doped

80

-

n

* SQW

/

Ir

MQW

v

400

440

480

520

560

600

Wavelength (nm) FIG. 26. Electroluminescence spectrum FWHM measurements for different LED active layer designs. The forward drive current is 20mA for all the data. The codoped and SQW data were previously published by Nakamura et nl. [7-111. (The multiple quantum well data were obtained from devices produced at Hewlett-Packard Company, San Jose, California 95131.)

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emission spectrum FWHM broadens somewhat as the indium content in the crystal active layer is increased (Nakamura, 1994a; Nakamura et at., 1995a). A possible reason for this behavior is phase separation in the GaInN layers due to increased lattice mismatch, which occurs when the growth temperature is reduced to obtain the higher indium content needed for longer wavelength devices. A further advantage of SQW LEDs is that the light output does not saturate at high drive currents. SQW LEDs rely on intrinsic band-to-band recombination for light emission and not on the D-A pair recombination in codoped devices. This factor is especially important for achieving population inversion in GaInN lasers. Nakamura et al. (1995b) reported that the shift in peak wavelength with drive current in SQW devices is greatly reduced compared to codoped devices for drive currents up to 100mA, although some color shift still occurs (Chichibu et al., 1996). The reasons for color shift are discussed in Sections I1 and 111 of this chapter and the references therein. The MQW heterostructure has also been used as the active layer in 111-V nitride LED devices (Koike et al., 1996). Toyoda Gosei in Japan has already released LED products with this type of structure. Figure 24c shows the schematic of a typical MQW active layer. MQW devices have many of the spectral characteristics of SQW devices, as shown in Fig. 26. One main reason for using the MQW structure is the possibility of reduced carrier overflow from the thin active layers, especially at high operating current densities (Kozodoy et al., 1997). Typical barrier thicknesses for MQW devices are 30- 150A, and typical well thicknesses are 20-30 8,(Kozodoy et ul., 1997). As mentioned earlier, the well thickness is kept thin to maintain good GaInN quality. The preferred number of wells varies in different laboratories from a few to tens of wells. The reported performance of MQW LEDs is as good as SQW LEDs. Reported power output values for 470-nm blue, 500-nm blue-green, and 522-nm green LEDs are 6.0, 6.1, and 3.7 mW, respectively, corresponding to external quantum efficiencies of 11.2, 12.3, and 7.8%, respectively (Chen et al., 1998a).

DISCHARGE SENSITIVITY AND RELIABILITY 6. ELECTROSTATIC CHARACTERISTICS A commonly encountered problem with 111-V nitride LEDs is their sensitivity to electrostatic discharge (ESD) damage as compared to conventional 111-V devices based on AlGaAs and AlGaInP. A GaInN LED that has been damaged by ESD will exhibit leaky current-voltage behavior under both forward and reverse bias as a result of the shunt path created in

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the p n junction by the discharge. Low light output behavior is also typical, and the region where the ESD damage occurs will appear as a dark spot. Multiple areas of damage on a single device can also be observed. The semiconductor industry relies on two testing models for determining the ESD sensitivity of a device. Both methods involve discharging a capacitor charged to a specific voltage through the device under test and increasing the voltage up to the point where the device fails. In the so-called machine model, the test circuit consists of a 200-pF capacitor, a 500-nH inductance, and a 25-0 series resistor to simulate the abrupt discharge caused by manufacturing or test equipment. In the human body model, the circuit consists of a 100-pF capacitor and a 1500-0 resistor, which simulates a charged “human.” The discharge test is performed at incrementally increasing voltage until the device under test fails. Because of the higher resistance in the human body model circuit, the voltage at which a device fails is typically much higher than for the machine model. AlGaAs and AlGaInP III-V LEDs do not exhibit ESD damage at any reasonable voltages and therefore do not require special ESD precautions in handling. These devices are considered ESD insensitive. On the other hand, commercial III-V nitride devices typically fail when subjected to voltages of the order of 200-400V when machine model testing is used. Voltage levels at which damage occurs using the human body model are about an order of magnitude higher. In any case, III-V nitride LEDs are considered “class 1” ESD sensitive. Antielectrostatic wristbands or gloves are recommended when handling III-V nitride LEDs, and all handling and test equipment and machinery should be electrically grounded (Nichia, 1996). It has been reported that yield loss is high if ESD handling procedures are not followed (Doverspike, 1997). Nevertheless, demonstrated improvements in the OMVPE epilayer structure can enhance the ESD damage resistance significantly (Nakamura, 1997). Another important LED device characteristic is reliability. This has been one of the important considerations for outdoor applications using highbrightness LEDs such as AlGaAs and AlGaInP LEDs (Stringfellow and Craford, 1997). Since III-V nitride LEDs have high brightness, they also must satisfy the same lifetime requirements for III-V nitride LEDs to be used in stringent applications such as traffic signals and outdoor full-color display panels. Reliability of III-V semiconductors typically relates to defects present in the material and their migration and multiplication as the device is operated. For III-V nitride materials, the lack of bulk GaN substrates makes it necessary to grow III-V nitride materials on a foreign substrate using heteroepitaxy. This invariably leads to the generation of dislocations in high

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density (Lester et al., 1995). The presence of this high dislocation density raises concerns for the long-term reliability of III-V nitride devices. Steigerwald (1997) studied the reliability of both codoped and SQW III-V nitride devices from Nichia Chemical Industries, Ltd., and Hewlett Packard Company. The devices were subjected to stress testing at various DC currents and temperatures for more than 10,000h. For traditional III-V devices, the light output (LOP) decrease is linear with the logarithm of the stress time. For codoped III-V nitride devices, however, the degradation can be modeled by an exponential function of the operation time t ; that is, the LOP is proportional to exp(-t/z), where z is the characteristic time constant. After chemically removing the encapsulation epoxy, however, the device light output is recovered. In addition, nonencapsulated III-V nitride LEDs show no sign of the gradual degradation seen in the encapsulated packages. Thus, Steigerwald concluded that the degradation is due to the packaging epoxy material, not the semiconductor itself. The problem is related to the epoxy becoming yellow over time. This “yellowing process” is accelerated at elevated temperatures and may lead to a shorter lifetime for III-V nitride LEDs. Packaged codoped devices also exhibit shorter lifetime when the operating current is increased. In this case, the UV peak that emerges at higher operating current after the donor-acceptor pairs are saturated may accelerate the degradation. Steigerwald (1997) observed a similar effect when comparing green and blue SQW LEDs. Because blue light is absorbed more readily than green light by the yellowed encapsulation epoxy, blue LEDs degrade more rapidly. Using current epoxy technology, the green SQW LED lamp lifetime may be two to five times longer than the lifetime for blue lamps. Osinski et al. (1996a) also studied the reliability of III-V nitride devices. Eighteen Nichia NLPB-500 LEDs were placed under stress test at various currents for up to 1OOOh. These codoped LEDs were observed to have minimal degradation within the testing time of 1000 h at 20 mA forward current. The LEDs showed accelerated degradation at elevated stress current, consistent with the model mentioned earlier (Osinski et al., 1996a). At extremely high stress currents of several amperes, degradation of the current-voltage characteristics and junction shorting have been observed (Osinski et al., 1996b). Interestingly, the shorts could be removed by applying a reverse bias. The damage and recovery can be repeated, although recovery may not always be complete. Even under these severe stress conditions, the damage has been found to be related to encapsulation materials, not the semiconductor material (Osinski et al., 1997). These studies show that the reliability of III-V nitride LEDs is excellent, providing an insight into the overall robustness of the III-V nitride system despite the poor crystalline quality of the layers. Some degradation is

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observed, but that is related to the encapsulation materials. Generally, the reliability is acceptable for outdoor commercial applications. These independent test results are consistent with the data sheets of commercial products provided by the LED vendors.

7.

WHITE

LEDS

With the realization of bright blue and green GaInN LEDs, coupled with the already available AlGaInP red LEDs, there is now enormous potential for creating highly efficient white illumination sources. By mixing together the emission from red, green, and blue LEDs in the proper proportions, white light, or almost any color hue, can be created. This technique is currently used in the production of large-screen LED video displays commonly used in sports arenas and stadiums. The development of whitelight LED sources as a replacement for conventional incandescent and fluorescent bulbs is also under intense investigation. Besides combining the light from discrete red, green, and blue devices, however, another way to obtain white light from LEDs is to use a blue GaInN chip to pump a phosphor. (This technique is analogous to a conventional fluorescent light bulb where a phosphor coating is pumped by the UV emission from a mercury vapor discharge.) The combination of the blue light from the LED chip and the broad yellow emission from the phosphor appears white to the human eye. Figure 27 shows the luminescence spectrum from a phosphor-pumped white LED. The peak at 465 nm is emitted by the LED, but some of this blue emission is absorbed by a phosphor coating over the LED chip and is reemitted as a broadband peak centered around 560nm. The spectral quality of the white light can be adjusted by changing the amount of phosphor applied to the chip or by using phosphors with different spectral emission characteristics. The subject of white LEDs is covered in much greater detail in Volume 65 of this series.

V. Summary and Conclusions The progress of the development in III-V nitride materials during the 1990s has been spectacular. The demonstration of highly efficient UV (Mukai et al., 1998a; Sat0 et al., 1998), blue, green, amber (Mukai et al., 1998b), and white (Nichia, 1996) LEDs as well as the recent achievement of violet injection laser diodes from Nichia has astounded the research com-

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40 I

-t .-

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&

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FIG.27. Electroluminescence spectrum of a "white-light" LED. The spectrum is a combination of the direct electroluminescence from a blue GaInN LED chip (the 465-nm peak) and fluorescence from a phosphor coating over the chip (the broad peak centered at 560-nm).

munity. The successful commercialization of these devices by numerous companies is the result of the pioneering work of Professor Isamu Akasaki and Dr. Shuji Nakamura in the late 1980s and early 1990s. Their realization of epilayer improvements by using low-temperature nucleation layers and the understanding and achievement of p-type doping were the enabling events for the advancement in this materials system. To date, blue and green LEDs with external quantum efficiencies exceeding 11% have been reported, values that challenge and, in some cases, exceed those achieved in AlGaInPand AlGaAs-based LED devices. Nitride-based amber LEDs with reported external quantum efficiencies comparable to AlGaInP amber LEDs have also been demonstrated (Mukai et al., 1998b). Even with these breakthroughs, however, the currently achievable material quality is still far from optimal, and the understanding of this material system is still quite limited. There are clearly many remaining issues that, when resolved, could lead to further advancement in the current device technology. Perhaps the most striking characteristics that distinguish the III-V nitride materials from other III-V compounds are the luminous properties, which appear to be unaffected by the high density of dislocations in the crystal (lo8 to 10'' dislocations cm -'). The dislocation density in all commercial nitride LEDs is typically on the order of lo9 cm-2 as compared with AlGaInP and AlGaAs where dislocation densities below lo4 cm-' are

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routinely obtained. Considering that in these other 111-V materials systems the luminous performance is inversely related to the dislocation density present in the epilayers (Lester et al., 1995), the achievement of highefficiency nitride LEDs is remarkable. In fact, the nitride materials grown are not even well-defined single crystals. Rather, they possess a mosaic-like crystalline structure. In addition to, and perhaps as a consequence of, the highly defected, imperfect nature of 111-V nitride crystals, their electronic properties are challenging to produce and control. Undoped GaN usually exhibits a high background electron concentration (typically 10' cm-3 or higher) with low mobility, which arises from either native defects (e.g., nitrogen vacancies) or contaminants (such as oxygen and silicon). Although low lo1*cm-3 hole concentration is achievable in OMVPE-grown GaN, typical concentrations are only in the 1 to 5 x 1017cm-3 range, which is much lower than what can be achieved in the GaAs or InP-related materials. This low p-type doping efficiency may not be a significant issue in most LEDs, but it is a concern for diode laser fabrication where thicker players are employed and low series resistance is crucial for high-power operation. These doping issues become even more severe for ternary compounds, such as GaInN (which usually exhibits a background electron concentration) and AlGaN (which usually exhibits high resistivity and low p-doping efficiency). For n-type doping using Si, cracking occurs at high doping levels for films only 3 to 4 pm in thickness. In an LED structure, the series resistance could also limit high-power performance, and usually, highly doped thick films are more desirable. This is particularly important for nitride LEDs grown on sapphire, where lateral current conduction in the n-layer is necessary. Another issue of importance for 111-V nitride LEDs is the internal quantum efficiency (i.e., the number of photons generated per electron crossing the pn junction), which involves the quality of the GaInN active layer as well as the confining layers. It has been observed that the quantum efficiency of blue, green, and amber nitride LEDs typically decreases with increasing In composition in the active region. A probable reason for the drop in efficiency is that the material quality degrades due to increasing lattice mismatch at the heterojunction when the In composition is increased. Efficiency of short wavelength devices, <450 nm, on the other hand, decreases as wavelength decreases. This is probably due to insufficient carrier confinement as the bandgap of the active layer increases. Although this is an issue that can be normally overcome by increasing the A1 composition in the AlGaN confining layer, the material quality and p-type doping efficiency may suffer. Similar to the development of other 111-V material technologies, the future improvement in the nitride materials will come from developing better crystal growth techniques.

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The final important issue for 111-V nitride LEDs involves the development of a high-volume production process for the epitaxial materials and LED chips. The efficiency of the blue and green nitride LEDs has become high enough that many new commercial applications are currently possible. Production volume has been increasing from less than 1 million units per month in 1994 to several tens of millions a month in 1998. At the same time, the unit price per lamp has dropped from well over a dollar to about $0.60 apiece. Even at this price, the cost of manufacturing these nitride devices, however, remains high when compared to other more mature LED and laser material technologies, such as AlGaAs and AlGaInP. The problem lies principally with low yields obtained from high-volume OMVPE processes. As has been described in the previous sections, the requirements for high temperature and high ammonia flow rate used in the current OMVPE process have made it difficult to achieve a large-area batch process with high uniformity and reproducibility. As a result, extensive light output and wavelength binning of the chips that go into the LED products is necessary, increasing the manufacturing cost. Currently, more efficient reactor designs are actively being pursued. The OMVPE process for nitride devices uses less toxic materials than those associated with production of the 111-V arsenides and phosphides, which inherently makes it a cheaper process once the yield issues are overcome. As with the development of other 111-V material technologies, low-cost manufacturing technology will eventually emerge. The prospects for 111-V nitride-based LEDs and LDs are virtually limitless. It is expected that with sufficient cost reduction through improved manufacturing technology, more applications and market opportunities will become possible for both nitride-based and AlGaInP-based high-brightness visible LEDs as well as lasers.

ACKNOWLEDGMENTS The authors acknowledge the full support of the Hewlett-Packard Optoelectronics Division and Hewlett-Packard Laboratories. Credit also goes to Dr. George Craford for his coaching and encouragement. Special thanks go to Dr. Daniel Steigerwald for providing valuable references and figures. We are also grateful to Dr. L. T. Romano and Dr. C. G. van de Van de Walle at Xerox, Palo Alto Research Center, for contribution of data, figures, and invaluable discussions. REFERENCES Abernathy, C. R.,MacKenzie, J. D., Pearton, S. J., and Hobson, W. S. (1995). “CCI, doping of GaN grown by metalorganic molecular beam epitaxy,” Appl. Phys. Lett. 66, 1969.

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Stringfellow, G. B., and Craford, M. G., eds. (1997). High Brightness Light Emitting Diodes, Semiconductors and Semimetals, Vol. 48, Series edited by Willardson, R. K., and Weber, E. R. (Academic Press, San Diego). Strite, S., and Morkoq, H. (1992). “GaN, AIN, and InN: a review.” J . Vac. Sci. Technol. B 10, 1237. Sun, C. J., Yang, J . W., Chen, Q., Lim, B. W., Anwar, M. Z., Khan, M. A., Temkin, H., Weismann, D., and Brenner, I. (1996). “Photoluminescence study of high quality InGaN/ GaN single heterojunctions.” Appl. Phys. Lett. 69, 668. Suzuki, M., Nishino, J., Onomura, M., and Hongo, C. (1998). “Doping characteristics and electrical properties of Mg-doped AlGaN grown by atmospheric-pressure MOCVD.” .I. Cryst. Growth 189-190, 511. Takeuchi, T., Sota, S., Katsuragawa, M., Komori, M., Takeuchi, H., Amano, H., and Akasaki, I. (1997a). “Quantum-confined Stark effect due to piezoelectric fields in GaInN strained quantum wells.” Jpn. J . Appl. Phys. 36, L382. Takeuchi, T., Takeuchi, H., Sota, S., Sakai, H., Amano, H., and Akasaki, I. (1997b). “Optical properties of strained AlGaN and GaInN on GaN.” Jpn. J . Appl. Phys. 36, L177. Takeuchi, T., Sota, S., Sakai, H., Amanoa, H., Akasaki, I., Kaneko, Y., Nakagawa, S., Yamaoka, Y., and Yamada, N. (1998). “Quantum-confined Stark effect in strained GaInN quantum wells on sapphire.” J . Cryst. Growth 189-190, 616. Taki, W., Tsuda, H., Tsukamoto, K., Kuwano, .N., Oki, K., Kawaguchi, Y., Shimizu, M., and Hiramatsu, K. (1997). “TEM analysis of compositional inhomogeneity in InGaN layers grown by OMVPE.” In The Proceedings of the Second International Conference on Nitride Semiconductors, Tokushima, Japan, p. 102. Tanaka, T., Watanabe, A,, Amano, H., Kobayashi, Y., Akasaki, I., Yamazaki, S., and Koike, M. (1994). “P-type conduction in Mg-doped GaN and Al,,,,Ga,,,,N grown by metalorganic vapor phase epitaxy.” Appl. Phys. Lett. 65, 593. Tanaka, S., Kern, R. S., and Davis, R. F. (1995). “Initial stage of aluminum nitride growth on 6H-silicon carbide by plasma-assisted, gas-source molecular beam epitaxy.” Appl. Phys. Lett. 66, 37. Tanaka, S., Kern, R. S., Bentley, J., and Davis, R. F. (1996). “Defect formation during heteroepitaxial growth of aluminum nitride thin films on 6H-silicon carbide by gas-source molecular beam epitaxy.” Jpn. J . Appl. Phys. 35A, 1641. Tanaka, S., Hirayama, H., Ramvall, P., and Aoyagi, Y. (1997a). “Artificially controlled formation of TnGaN quantum dots on AI,Ga, -,N surfaces.” In The Proceedings of the Second International Conference on Nitride Semiconductors, Tokushima, Japan, p. 48. Tanaka, S., Iwai, S., and Aoyagi, Y. (1997b). “Reduction in the defect density in GaN using ultra-thin AIN buffer layers on 6H-Sic.” J . Cryst. Growth 170, 329. Uchida, K., Watanabe, A., Yano, F., Kouguchi, M., Tanaka, T., and Minagawa, S. (1996). “Nitridation process of sapphire substrate and its etTect on the growth of GaN.” J . A p p l . Phys. 19, 3487. Uchidd, K., Tokunaga, H., Inaisbi, Y., Akntsu, N., and Matsumoto, K. (1997). “111-V nitrides growth by atmospheric-pressure MOVPE with a three layered flow channel.” Muter. Res. Soc. Symp. Proc. 449, 129. Vechten, J. A. V., Zook, J. D., Homing, R. D., and Goldenberg, 9. (1992). “Defeating compensation in wide gap semiconductors by growing in H that is removed by low temperature de-ionizing radiation.” Jpn. J. Appl. Phys. 31, 3662. Vodakov, Y. A,, Mokhov, E. N., Roenkov, A. D., Boiko, M. E., and Baranov, P. G. (1998). “High rate GaN epitaxial growth by sublimation sandwich method.” J. Cryst. Growth 183, I a.

3 HIGH-BRIGHTNESS NITRIDE-BASED VISIBLE-LIGHT-EMITTING DIODES207 Wakahara, A,, Tokuda, T., Dang, X. Z., Noda, S., and Sasaki, A. (1997). “Compositional inhomogeneity and immiscibility of a GaInN ternary alloy.” Appl. Phys. Lett. 71, 906. Wang, C., and Davis, R. F. (1993). “Deposition of highly resistive, undoped, and p-type, magnesium-doped gallium nitride films by modified gas source molecular beam epitaxy.” Appl. Phys. Lett. 63, 990. Weeks, T. W., Jr., Kum, D., Carlson, E., Perry, W. G., Ailey, K. S., and Davis, R. F. (1994). “Growth and characterization of undoped and doped GaN thin films by OMVPE.” In Trunsactions of the Second International High Temperature Electronics Conference, Charlotte, NC, USA, p. P-141. Weeks, T. W., Jr., Bremser, M. D., Ailey, K. S., Carlson, E., Perry, W. G., and Davis, R. F. (1995). “GaN thin films deposited via organometallic vapor phase epitaxy on a(6H)SiC(OOO1)using high-temperature monocrystalline A1N buffer layers.” Appl. Phys. Lett. 67, 401. Wickenden, A. E., Rowland, L. B., Doverspike, K., Gaskill, D. K., Freitas, J. A., Jr., Simons, D. S., and Chi, P. H. (1995). “Doping of gallium nitride using disilane.” J . Electron. Mater. 24, 1547. Wright, A. F., and Nelson, J. S. (1995). “Bowing parameters for zinc-blende Al,-,Ga,N and Gal _,In,N.” Appl. Phys. Lett. 66, 3051. Wu, X. H., Brown, L. M., Kapolnek, D., Keller, S., Keller, B., DenBaars, S. P., and Speck, J. S. (1996a). “Defect structure of metal-organic chemical vapor deposition-grown epitaxial (0001) GaN/Al,O,.” J . Appl. Phys. 80, 3228. Wu, X. H., Fini, P., Keller, S., Tarsa, E. J., Heying, B., Mishra, U. K., DenBaars, S. P., and Speck, J. S. (1996b). “Morphological and structural transitions in GaN films grown on sapphire by metal-organic chemical vapor deposition.” Jpn. J . App/. Phys. 35B, L1648. Wu, X. H., Kapolnek, D., Tarsa, E. J., Heying, B., Keller, S., Keller, B. P., Mishra, U. K., DenBaars, S. P., and Speck, J. S. (1996~).“Nucleation layer evolution in metal-organic chemical vapor deposition grown GaN.” Appl. Phys. Lett. 68, 1371. Wu, X. H., Elsass, C. R., Abare, A,, Mack, M., Keller, S., Petroff, P. M., DenBaars, S. P., Speck, J. S., and Rosner, S. J. (1998a). “Structural origin of V-defects and correlation with localized excitonic centers in InGaN/GaN multiple quantum wells.” Appl. Phys. Lett. 72, 692. Wu, X. H., Fini, P., Tarsa, E. J., Heying, B., Keller, S., Mishra, U. K., DenBaars, S. P., and Speck, J. S. (1998b). “Dislocation generation in GaN heteroepitaxy.” J . Cryst. Growth 189-190, 231. Yamasaki, S., Asami, S., Shibata, N., Koike, M., Manabe, K., Tanaka, T., Amano, H., and Akasaki, I. (1995). “P-type conduction in Mg-doped Ga,,,,In,,,,N grown by metalorganic vapor-phase epitaxy.” Appl. Phys. Lett. 66, 1112. Yoshida, S., Misawa, S., and Gonda, S. (1983). “Epitaxial growth of GaN/AlN heterostructures.” J . Vac. Sci. Techno/. B 1, 250. Yuan, C., Salagaj, T., Gurary, A., Thompson, A. G., Kroll, W., Stall, R. A,, Hwang, C.-Y., Schurman, M., Li, Y., Mayo, W. E., Lu, Y., Krishnankutty, S., Shmagin, I. K., Kolbas, R. M., and Pearton, S. J. (1995). “Investigation of n- and p-type doping of GaN during epitaxial growth in a mass production scale multiwafer-rotating-disk reactor.” J . Vac. Sci. Techno/.B 13, 2075. Zheleva, T. S., Nam, 0. H., Bremser, M. D., and Davis, R. F. (1997). “Dislocation density reduction via lateral epitaxy in selectively grown GaN structures.” Appl. Phys. Lett. 71, 2412. Zolper, J. C., Wilson, R. G., Pearton, S. J., and Stall, R. A. (1996). “Ca and 0 ion implantation doping of GaN.” Appl. Phys. Lett. 68, 1945.

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SEMICONDUCTORS AND SEMIMETALS, VOL. 64

CHAPTER 4

Organic LED System Considerations Yoshiharu Sato YOKOHAMA RESEARCH CENTER MITSUBISHI CHEMICAL CORWRATION YOKOHAMA, JAPAN

I. INTRODUCTION . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11. RELIABILITY ISSUES FOR PRACTICAL APPLICATION . . . . . . . . . . . . .

1. Degradation Mechanism . . . . . . . . . . 2. Operation Stability . . . . . . . . . . . . . 3. Temperature Dependence . . . . . . . . . . 111. DRIVING SCHEME . . . . . . . . . . . . . . 1. Driving Method . . . . . . . . . . . . . . 2. Passive us Active Matrix . . . . . . . . . . IV. OLED PANELS.. . . . . . . . . . . . . . . 1. Material Issues . . . . . . . . . . . . . . 2. Process Considerations . . . . . . . . . . . 3. Monochrome Panels . . . . . . . . . . . . 4. Multicolor and Full-Color Panels . . . . . . . V. SUMMARY . . . . . . . . . . . . . . . . .. REFERENCES. . . . . . . . . . . . . . . . .

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209 214 214 217 228 231 231 236 239 239 244 246 241 248 252

I. Introduction The advent of organic electroluminescent devices advanced from a long history of development of organic photocondoctors (OPCs) for xerographic copy machines. While the organic solar cell has a similar mechanism as the organic photoreceptor, the organic light-emitting device (OLED) or organic electroluminescent (EL) device has a different mechanism, being almost the reverse process. Organic photoconductor technology has been well established after 30 yr of continuous efforts. In contrast to OPCs, OLED research looks back only 10 yr to the first development. Since the first double-layer device using organic molecular materials was published by Tang and VanSlyke (1987), great progress in luminous efficiency and color tunability has been achieved. OLED technology is now 209 Copyright 02030 by Academic Press All rights of reproduction in any form reserved. ISBN 0-12-752173-9 ISSN 0080-8784/00 $30.00

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attracting much attention in the field of flat panel displays (FPDs). In the application of organic EL devices to practical displays, there are three fundamental issues to be considered: (1) stability (lifetime), (2) power efficiency, and (3) full-color capability. During the 10 yr under consideration some technologies have greatly contributed to the first commercial product made of organic active materials, which are listed below chronologically: 1987 Double-layer device (Kodak) 1988 Doping technology (Kodak) 1989 PPV (poly p-phenylene vinylene) (Cambridge University) 1991 Quinacridone doping (Pioneer) 1991 Doped polymer (Asashi Chemical) 1992 AlLi cathode (Pioneer) 1993 Starburst molecule (Osaka University) 1995 Application of poly-Si TFT to OLED (TDK) 1995 CCM (color-changing medium) method (Idemitsu) 1997 Ultrathin LiF layer at cathode interface (Kodak) 1997 Commercial product for car audio display (Pioneer) 1998 Demonstration of full-color panel by shadow masking (Pioneer). Each technological step improved the efficiency or stability of the device, contributing to the success of organic devices in the display field. Features of OLEDs are

.

Solid-state device of total thickness less than 2 mm Emissive device with wide viewing angle Full-color capability Low-voltage operation below 10 V DC Low power consumption High resolution with pixel pitch less than 0.3 mm Fast response time of the order of lops, compatible with video rate operation Polymer substrate capable of being used to fabricate a flexible display of light weight

Based on these features, applications of OLED have been investigated or proposed in the fields shown in Fig. 1. The first commercial product by Pioneer Electronic Corp. is a small character display based o n a passive dot matrix for car audio system. Pioneer, NEC, and Idemitsu have demon-

-

4 ORGANIC LED SYSTEM CONSIDERATIONS

I

211

1

I

I

multi

1car navigation I small character display mobile phone, pager camera finder, game car stereo

1 A T F I

segmented

light source

LSI optical link

7-segment

backlight for LCD,

copy machine, facsimile LCD projection

FIG. 1. Applications of OLED.

strated full-color QVGA panels. Besides display applications, light sources such as backlights have also been considered for watches and mobile phones. The four steps of research in OLED technology -material, device, panel, and system-are illustrated in Fig. 2. We are able to utilize a variety of organic materials, which can be controlled by proper molecular design to our desired properties such as emission color or carrier transport polarity. The potential of organic materials seems infinite; however, only a limited number of them has been used for practical devices; the reason is described in the next section. The second step is how to use organic materials in OLEDs. The most important point is that we must use organic materials in the form of thin films. The first proposed OLED has a bilayer structure consisting of a hole transport and an emitting layer. With the development of new materials, variations of this bilayer structure have been investigated. Device performance such as efficiency or lifetime is strongly subject to this layer structure. Simultaneously, we have learned much about injection and emission mechanisms by comparing different materials and device structures. Employing a proper material in a proper layer could lead to OLEDs that can compete with other display devices.

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RESEARCH STEPS

FIG. 2. OLED research from materials to panels.

To fabricate a display panel, we must face a patterning problem that is inherently difficult, since organic layers can not stand a conventional photolithographic process or a wet process. Special techniques have been introduced to avoid this problem; that is, the use of shadowing effects in vacuum deposition. We must also investigate driving method for the OLED. Electronic circuit technology is also required to obtain the best performance from OLEDs. The final step is designing the OL,ED into the system or application. Depending on the field of application, requirements for panel performance might be different. Since an OLED panel must compete with the already existing flat panels such as VFDs, STN-LCDs, inorganic EL, and LEDs, the competitiveness of OLEDs is the most important point of view for the future. OLED technology has several issues in each step of research. The present status of OLEDs is summarized in Table I. Details of each issue will be described and discussed in the following sections.

4 ORGANIC LED SYSTEM CONSIDERATIONS

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TABLE I ISSUESAND PRESENT STATUS OF OLED Issue

Efficiency

Stability

Full color

Target

Present status

Future issues

B 1 (Im/W) G: 6 R: 3 Lifetime> 10,000h at 100cd/m2

6 (Im/W) (Idemitsu) 10 (Pioneer) 0.5 (Kodak) 20,000h (Idemitsu) G: 50,000 (Kodak) R: 20,000 (Kodak) Y:20,000 (Sanyo) Blue/CCM (Idemitsu) White/CF (Mitsubishi, TDK) Patterned RGB (Pioneer, NEC) Commercial product (Pioneer) 5-in. QVGA (Idemitsu, NEC, Pioneer) 5-in. QVGA (TDK poly-Si)

Overall efficiency Red: efficiency and chromaticity Low-voltage operation Heat resistivity Good white-emitting device High-duty operation Differential aging New approach?

Patterning method

Passive matrix

Active matrix Panel fabrication

Cathode patterning; encapsulation; production

B: G R: B:

All dry process; drying agent; multichamber system

Gray scale Current-driving 1C Low-voltage operation Low-temperature poly-Si Cost consideration Encapsulation Mass-production system Repair of pixels

In Section I1 on the reliability of OLEDs, the most important issue from a practical point of view is reviewed. Progress in both materials and device structures has led to a great improvement of operation lifetime. Section I11 provides an understanding of another aspect of OLED panels -the driving scheme. It builds on detailed information on the basic characteristics of OLEDs. Here, a passive vs an active matrix is a very interesting issue. All the subjects considered in this section are related to the field of application. In Section IV, various type of OLED panels are introduced, starting with the simplest one, the monochrome passive matrix panel. Multi or full-color capability is a next step for this technology to spread out widely in the display market. Some approaches to achieve full-color panels are discussed. Finally, two “cultures” in OLED technology, molecular OLEDs and polymer LEDs, are compared. Future trends are stressed, and we argue that OLED technology has much potential for flat panel displays.

214

YOSHIHARIJ SATO 11. Reliability Issues for Practical Application

Among the three key issues pointed out in the first section, stability is most critical from a practical point of view. Several researchers have reported good stability data that demonstrated the long-term durability of OLED devices (Hamada et al., 1995; Hosokawa et al., 1997; Nakada and Thoma, 1996; Sat0 et al., 1996; Shi and Tang, 1997; VanSlyke et al., 1996b). However, the degradation mechanism of the device due to operation is still not fully understood. Although there are a few preceding technologies of organic devices such as OPCs, solar cells, and photodiodes, OLEDs reside in a different world; they are subject to a stress of much higher current density, compared with the other organic devices. The organic photoconductor, which is widely used in copy machines, is a successful example of an organic device. It undergoes electrostatic charging and photodischarge processes. The amount of charge passing through an organic layer of 40 pm in a cycle is typically of the order of C/cm2. This number corresponds to lo-’ C/cm2 after 100,000 copies. On the other hand, a typical drive current density in an OLED is about 1 mA/cm2; this means that OLED consumes the same amount of charge in only 10 s! When we drive an OLED in a passive matrix scheme, the drive current density will be much higher, of the order of 100mA/cmZ in a short period of time. The operation of organic devices at such a high current density has never been experienced; both the thermal and the electrochemical stability of the device must be investigated. Another difficulty for OLEDs, particularly for molecular OLEDs made of small molecules, is crystallization of amorphous organic thin films. in O P C technology, a binder polymer is usually employed to give the photoreceptor the mechanical strength and surface morphology that are required for the copy process such as charging, developing, and cleaning. Generally 50 wt% binder polymer is used so that the crystallization problem of charge transporting and generating materials does not seem to occur explicitly. This crystallization problem is considered in the following subsections in detail.

1. DEGRADATION MECHANISM

The degradation of an OLED during operation appears in four modes: (1) the decay of luminance, (2) a voltage increase in the constant-current mode, (3) the growth of nonemissive areas (dark spots), and (4) an abrupt

4 ORGANIC LED SYSTEM CONSIDERATIONS

215

short circuit or electrical breakdown (Sato et at., 1998). The origin of these failure modes has been investigated and the following degradation mechanisms have been proposed by several groups (as summarized in Fig. 3): Morphological instability or crystallization of the amorphous organic films (Han et at., 1996) Interdiffusion between the organic layers (Fujihira et al., 1996) Photooxidation of the organic layer by self-emitting light (Stairing et al., 1997; Zyung and Kim, 1995) Electrochemical decomposition of the organic materials (Sato and Kanai, 1994) Oxidative polymerization of Alq in the presence of water (Papadimitrakopoulos and Zhang, 1997) Poor contact between the organic layer and the electrodes (Do et al., 1997; McElvain et al., 1996; Sat0 and Kanai, 1994) Oxidation of the organic layer by IT0 (Scott et at., 1996) An energy barrier at the anode interface (Adachi et al., 1995) The importance of encapsulation and protective layer (Burrows et al., 1994) These mechanisms can be classified into three categories: (1) thermally induced degradation, (2) electrochemically or charge-induced degradation, and (3) interfacial degradation. For molecular OLEDs, the first degradation

#

FIG. 3. Bilayer structure of device and degradation mechanisms,

216

YOSHIHARUSATO

path is related to the morphological instability of amorphous organic films. This type of instability is observed as crystallization or aggregation in heat treatment such as storage at an elevated temperature and during operation via joule heating. Amorphous organic films tend to crystallize or aggregate on heating. It has been quite often observed that a uniform organic film of small molecules from its as-deposited state gradually crystallized even at room temperature. As it crystallizes, the film morphology has changed into a rough structure, which gave a diffraction pattern assigned to the organic components (Sato et a/., 1998). In the center of the crystallized region, protrusions have grown out and the peripheral area forms a kind of groove. The diameter of this mountain or valley structure was about 10pm. The second path of degradation comes from chemical instability of the materials used. This kind of degradation can be ascribed to electrochemical decomposition of molecules. This means that the device failure mode is coulombic and depends on the total amount of charge injected into the device. The electrochemical instability of organic materials could be assessed with cyclic voltammetric measurements. One such measurement was carried out for a few hydorazone compounds used as an organic photoconductor that does not show any problem in the copy process. During the oxidation process from the neutral state to the cation radical state, these hydrazone compounds decomposed electrochemically. This electrochemical instability corresponds to a quite short lifetime of an OLED with a hole transport layer of the hydrazone compound (Sato and Kanai, 1994). It was further confirmed by the experiment that cooling the device to -8°C did not improve the operational lifetime. Concerning the emitting materials, another class of unstable compounds includes oxadiazole derivatives and metal azole complexes (Tamoto et al., 1997). The third possibility is interfacial deterioration caused by migration of molecules or a chemical reaction with oxygen or water. This degradation could be avoided by considering device structures such as an I T 0 buffer layer or a cathode interface layer, or by doping in some cases. Some of these degradation mechanisms are common for molecular OLEDs and polymer LEDs. Interdiffusion processes at the metal-organic interface and organicorganic interfaces are important since they may alter the EL characteristics of the device. In some cases, thermal processes cause the interdiffusion. In the bilayer structure, the glass transition temperature (ZJ affects such a molecular migration behavior. The conventional hole transport material TPD shown in Fig. 3 has a relatively low T, value of 63°C. Annealing of this bilayer sample resulted in an interdiffusion between hole transport and emitting layers, leading to an increase of the driving voltage. The metal contact at both cathode and anode (ITO) interfaces is most important in terms of the stability and the injection-emission mechanism.

4 ORGANIC LED SYSTEM CONSIDERATIONS

217

Using I T 0 buffer layers and cathode interface layers has been reported to improve the stability and efficiency of OLEDs. In the next part of this section, the preceding three kinds of degradation modes are examined in terms of the operation characteristics of OLEDs.

2. OPERATION STABILITY In this part, the stability of OLEDs that are small molecules is discussed mainly in conjunction with their operation characteristics. The operational stability was investigated from the viewpoints of materials and device structures in an attempt to elucidate the factors that dominate the operational lifetime or degradation of the device. A bilayer structure made of hole transport and emitting layers was chosen as a starting point to achieve an understanding of the operation characteristics. As already described, T, of the hole transport material is critical because it influences the thermal stability of amorphous organic films. Since the T, of Alq has been reported to be sufficiently high (Naito and Miura, 1993), the most thermally unstable part of the device is the TPD layer. High-T, hole transport materials have been developed to achieve device thermal stability. Some of these materials are shown in Fig. 4; the T, values and the ionization potential, HOMO-LUMO gap are summarized in Table 11. When we suppose a mobile or vehicle application for an OLED display, a heat resistance up to 85°C is required. Therefore, a hole transport material (HTM) with T, above 85°C is desirable. Instead of TPD, a-NPD having a T, of 96°C was used as a hole transport layer to examine the operation characteristics. OLED cells having a-NPD (60 nm)-Alq (75 nm) layers with MgAg cathode were operated in the constant current mode. Relative decay curves of luminance were demonstrated in Fig. 5, where that of TPD cells was included for comparison. The a-NPD cell exhibited an improved decay, demonstrating the T, effect. When the lifetime, defined as half-decay time of luminance, of the a-NPD devices is plotted as a function of drive current density, we find it to be inversely proportional to the square of current density, as shown in Fig. 6, which suggests that the operational lifetime was dominated by a thermal process such as joule heating. To make the preceding point clear, the temperature of the device during operation was measured using a radiation thermometer at the glass surface of an OLED cell, which has an I T 0 (120 nm)-TPD (60 nm)-Alq (75 nm)MgAg (100 nm) layer structure. The result is shown as a function of input power in Fig. 7. An increase of the temperature was found to be proportional to the input power up to 2 W/cm2. Since the input power is of the

218

YOSHIHARU SATO

DN0

"8 % ?? TCTA

1-TNATA

FIG. 4. Hole transport materials with high T,.

TABLE I1 PROPERTIES OF HOLETRANSPORT MATERIALS

HTM

TPD M-NPD P-NPD TPTE 1-TNATA TCTA

mp ("C)

T, ("C)

E, (ev)

IP (eV)

180

63 96 88 130 113

3.20 2.91 2.91

5.15 5.20 5.08 5.05 5.8

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100

1000

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Operation Time [hour] FIG. 5. Decay curves of luminance for the devices with a-NPD and T P D as the hole transport layers.

order of 0.1 W/cm2 in the operation conditions used in Fig. 5, the increase of the device temperature can be estimated within a few degrees from this measurement. This fact may lead to a possibility of local heating or concentration of current in the device; poor contact between organic layers and electrodes is described to support this idea in the subsequent section. In a simple one-dimensional model, the increase of temperature of OLED device, A T can be formulated as follows: AT = Q/(h

0

+ k/t)

0.01 0.02 0.03 0.04 1/(Current Density [mA/cm*] )2

0.05

FIG. 6 . Dependence of half-decay time on the square of driving current density

220

YOSHIHARIJ SATO

20 0

2

4

6

8

10

Input Power [W/cmz] FIG.7. Measurement of the device temperature during operation together with theoretical values.

where Q is the input power, h is the heat dissipation through air, k is the thermal conductivity of glass substrate, and t is the thickness of the glass substrate. Putting typical values of h, k, and t into the preceding equation results in the dashed line in Fig. 7. Actually, the term h is negligible compared with heat dissipation through the substrate. We clearly see that the deviation of the calculated temperature from observed temperature becomes larger at higher input powers. This deviation can be attributed to the limitation of a one-dimensional model; lateral heat dissipation or a two-dimensional model must be taken into account. The preceding results, however, show that the one-dimensional analysis gives the device temperature in the practical range of the input power, below 2 W/cmZ, and also indicates the need for a good thermal conductor for the substrate and heat sink structure. We often observed an initial rapid decay of the luminance followed by slow decay in OLED operation. Although the cause of this initial instability is not clear, we must stabilize the device performance from two approaches: (1) to improve the decay characteristics by optimizing the materials and the device structure and (2) to carry out aging treatment prior to operation at the sacrifice of loss of initial luminance. For the latter purpose, the OLED was driven at a higher current density of 90mA/cmZ for 4 h before the luminance degradation was measured. The same devices as evaluated in Fig. 5 were operated at 15 mA/cm2 after this aging operation, and, as shown in Fig. 8, the decay curves are quite different from those in Fig. 5 in that the luminance remains at the initial level during the first 100 h for the a-NPD

221

4 ORGANIC LED SYSTEM CONSIDERATIONS I 1

0

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0 0

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00 0

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o

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aging conditon: 9 0 d c m z - 4 hours

J=15 d c n 9

FIG. 8. Luminance decay after aging at a high current density.

device. Compared with the decay curve of the TPD device, such an aging treatment is effective for a device with a high-T, hole transport material. If an initial luminance loss of 50% is allowed, this approach is valid for low-luminance but high-stability applications such as the backlighting of watches. Essential improvement of the operational stability is desired for applications in information displays. OLED research has focused on both materials and device structures to improve the operational stability. The evolution of the OLED structure is schematically depicted in Figs. 9a, 9b, and 9c. As already described, operation stability was insufficient with the fundamental bilayer structure (Fig. 9a). A contact problem between hole transport layer and I T 0 anode can be considered as one of the causes of degradation.

B.dter Layer

(a)

@I

(4

FIG. 9. Evolution of the layer structure in molecular OLEDs

222

YOSHIHARUSATO

To improve this contact, an I T 0 buffer layer was introduced, as shown in Fig. 9b. As an I T 0 buffer layer, copper phthalocyanine was used to lower the drive voltage because its low ionization potential ensures injection of holes from IT0 into the hole transport layer. After a cleaning process, such as ultrasonic agitation in an organic solvent or UV-ozone treatment, the surface of the I T 0 substrate becomes quite wettable with a contact angle of water of On, while the contact angle on the surface of the a-NPD film was observed from 70 to 80". This fact indicates that there is a mismatch between the surface energies of I T 0 and a-NPD layers, leading to crystallization or restructuring of the hole transport layer. Copper phthalocyanine (CuPc) is believed to improve this contact. Other buffer materials such as amorphous carbon, vanadium oxide, and conductive polymers have been reported to be functional for the same purpose. The device shown in Fig. 9b with a CuPc-a-NPD-Alq trilayer structure exhibited an improved lifetime of about 4000-5000 h at 20 mA/cm2 operation (VanSlyke et al., 1996b). The increase of the drive voltage during operation was also suppressed to an allowed level. This concept of an I T 0 buffer layer also worked for polymer LEDs, because oxidation of polymer layers by the action of I T 0 has been reported (Staring et al., 1997; Zyung and Kim, 1995). Both the initial drive voltage and the voltage increase due to operation were definitely improved by the insertion of the CuPc layer. The thickness of CuPc layer was only 20nm so that it was not enough to planarize the I T 0 surface. Therefore, the role of CuPc layer can be considered to be electrical and interfacial. Yamamori et al. (1998) reported using a thicker conductive buffer layer using a hole-transporting polymer doped with a strong acceptor to demonstrate a comparable device performance. In this technology, the polymer buffer layer could be as thick as 650 nm, which makes it probable that the device characteristics become independent of I T 0 properties such as surface roughness, inhomogenity, or work function. In Fig. 9b, the cathode is also improved using AlLi alloy with Li content below 1 wt%. The work function was further lowered to around 3.0eV compared to 3.6 eV of MgAg. A reduction of the driving voltage, an increase of luminous efficiency, and stability enhancement were all attained with the AlLi cathode. However, these low-work-function cathodes are so sensitive to ambient moisture that some protective layer or encapsulation technique is inevitably required to suppress dark spot formation. The cathode is further improved by inserting an ultrathin layer (0.5 nm) of insulating material such as Li,O or LiF (Hung et al., 1997; Wakimoto et al., 1997a). This nano-layer improved electron injection from the aluminum cathode into the organic layer. The mechanism of electron injection is now

223

4 ORGANIC LED SYSTEM CONSIDERATIONS

I

8o

1

’ ‘0001

.B

cl 2000

1

10

Applied Voltage [V]

oo

c

0 0

0

5

Cz-i/Ag

1

0

10

15

20

Applied Voltage [V]

FIG. 10. Effect of the Cz-TPD cathode interface layer on I-V and L-V characteristics.

under investigation; the formation of electric double layer at the interface is one of the possibilities (Lee et al., 1998). A cathode such as the LiF-A1 system is preferable in terms of device stability and processing. Instead of alkali halides, organic materials have been also tried at this interface. As shown in Fig. 10, Cz-TPD from the TPD family was deposited between the cathode and emitting layer with a thickness of 15 nm (Kanai et al., 1997). The interesting effect of this cathode interface layer (CIL) was that high-work-function metals such as Ag could be used. The I-V characteristics of the devices with and without CIL are shown in Fig. 10a. The Ag cathode device without CIL exhibited almost ohmic behavior. This suggests that the Ag atom migrates into organic layers, causing leaky current. CIL is supposed to avoid this penetration of cathode metal atom as a barrier layer. The L-V characteristics shown in Fig. 10b indicate that the Cz-TPD layer gave a comparable level of luminance. This effect can be interpreted by a rough morphology of the Cz-TPD layer, which enabled electron injection from the Ag cathode with the help of a strong local field. CIL combined with the Ag cathode resulted in a storage stability under a nitrogen atmosphere. Without CIL, a rapid growth of dark spots was observed. Further improvement of dark-spot suppression was observed for the cell with Ag cathode (Sato et al., 1998). This long-term-stability device derived

224

YOSHIHARUSATO

from the use of an inactive Ag cathode was made possible by Cz-TPD CIL. Operation tests of Cz-TPD-Ag cells showed a rapid increase of voltage in the constant DC current mode, which seems to indicate, that Cz-TPD is electron-transporting but with a low mobility. In the structures of Figs. 9b and 9c, the emitting layers (EMLsj are doped with efficient organic dyes. Since the first introduction of doping technology in OLEDs (Tang et al., 19891, dopant dyes have been intensively investigated to improve efficiency and stability. Laser dyes such as coumarin and DCM have been found to be effective due to their high quantum yield of fluorescence. In the meantime, quinacridone as a green dopant and rubrene as a yellow dopant have been demonstrated to be practical materials in terms of both efficiency and stability enhancement. Rubrene has been intensively investigated as a dopant. This dye is most interesting from the viewpoint of concentration quenching and because of its bipolar character. Two different types of cell structures have been reported (see Fig. 11) (Hamada et a!., 1995; Sat0 et al., 1996). In the type I structure already described, rubrene was doped into the Alq layer, while it was doped into TPD as an emitting layer in type 11. The type I device demonstrated a lifetime of longer than 3000 h at an initial luminance of 535 cd/m2 with a current density of 15 mA/cm2. The type I1 device also exhibited a lifetime of 3500 h at an initial luminance of 530cd/m2 with a current density of 5 mA/cmZ.Morphology of amorphous organic thin films has been observed to be stabilized by doping (Sato and Kanai, 1994). TPD has a low T, and easily crystallizes; however, rubrene doping is supposed to delay such a crystallization, resulting in a long-term durability. In addition to the morphological stability given by rubrene doping, a different mechanism of emission between the rubrene-doped and the undoped cells is supposed to reflect the stability. Two emission mechanisms have been proposed for organic EL devices: energy transfer from host to dopant and trapping at a dopant site. Since HOMO and LUMO of the

FIG. 11. Two types of cell structure having rubrene as a dopant.

4 ORGANIC LED SYSTEM CONSIDERATIONS

225

rubrene molecule are located within the energy gap of Alq, rubrene may act as both a hole and an electron trap, suggesting recombination in the rubrene molecule as a possible emission mechanism. To clarify this point, a device that has its Alq emitting layer doped with rubrene only near the cathode interface was fabricated. An EL spectrum was obtained consisting mainly of the rubrene contribution, which supports the latter emission mechanism for the rubrene-doped device. In terms of stability, the device with the entirely doped Alq layer gave the best result compared with the partially doped device, which suggests that the delocalized emission sites favor operational stability. A general tendency for a longer lifetime with increasing of rubrene concentration was also observed; rubrene dye, unlike other dyes such as coumarin or quinacridone, does not show strong concentration quenching, which enables heavy doping up to 7 wt% without a significant degradation of luminous efficiency (Sato et al., 1997). The rubrene-doped device also revealed the dependence of the lifetime on the cathode metals. The work function of the cathode materials was measured as 3.4 eV (AlLi), 3.6 eV (MgAg), and 4.0 eV (Al), respectively. AlLi exhibited a further improved stability compared with MgAg device, while the A1 device had a quite short lifetime of only 2 min under the same operation conditions. This cathode dependence shows that a lower work function cathode favored stable operation, which suggests the change of carrier injection balance with cathode metals affects the operation lifetime strongly even in the doped device. Next, the dependence of operation lifetime on dopants was examined, as shown in Fig. 12; these doping materials were tested in the Alq host layer. Luminance decay curves are also shown in Fig. 12 for each dopant. Four dopants except quinacridone improved the operation lifetime compared with the undoped cell. The reason quinacridone did not improve the device stability can be ascribed to its chemical structure; it has both amino and carbonyl groups that easily form a dimer through intermolecular hydrogen bond (see Fig. 13). This dimer formation generally provides a nonradiative decay path for the excited singlet. Quinacridone can be chemically modified by replacing its amino group with a methylated amino group to achieve a device with a greatly improved lifetime (Shi and Tang, 1997). Although the best data were obtained with rubrene and diphenyltetracene (DPT), the other condensed aromatic dopants were also effective in improving the operational stability. 9,lO-Diphenylanthracene and perylene did not change EL spectra; an increase of drive voltage was observed, which indicated these three dopants acted as a trap in the Alq layer, which modified the emission mechanism. Another possibility for improving stability using the preceding condensed aromatic dopants is the morphological effect observed for rubrene-doping,

226

YOSHIHARU SATO

&meridone

perylene

9,1Cdiphe1&mthracwe (9,lO-DPA)

1

8

0.8

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0.6

cl

P *z

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rd

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0.2

-

A

perylene

'I 0

DPT

A

quinacridone

9,10-DPA I

0 0.1

1

10

100

I

1000

10000

Operation Time [ Hour 3 Fic;. 12. Dependence of the operation stability on different dopants.

FIG. 13. Dimer structure of quinacridone via intermolecular hydrogen bond.

4 ORGANIC LED SYSTEM CONSIDERATIONS

227

and a singlet oxygen quenching effect (Stevens and Algar, 1968). The latter possibility was not well understood or experimentally supported; we just speculate residual oxygen in the organic layers might be captured by some gettering effect of these molecules. To summarize the doping technology, requirements for dopant dyes are discussed in the following. Although we see fewer restrictions concerning dopant dyes compared with host materials, we must consider a general phenomenon of fluorescent dye or pigment -concentration quenching. Concentration quenching is thought to be mainly due to dimer formation between dopants via intermolecular or van der Waals forces. If dopants form a dimer conformation, a triplet nonradiative path is generated, leading to quenching; a decrease of fluorescent yield and/or a red shift of fluorescent spectra. Quinacridone exhibits a strong molecular interaction via hydrogen bonding, which can be considered as a main reason for the short lifetime. Another example is perylene. Perylene tends to form dimers, which shift the fluorescent spectrum from blue to green. VanSlyke et al. (1996a) reported that a perylene-doped blue-emitting cell showed spectral change during operation; these phenomena may be explained by dimer formation of perylene. Since blue dopant dyes are generally small in size and molecular weight, they are likely to aggregate after being doped. Distyrylbiphenyl derivatives with relatively large molecular weight have been successful as a blue dopant. Concerning red dopants, the situation is further complicated by the fact that extended n-conjugated systems interact with each other to cause severe concentration quenching. Therefore, a special design is required for red dopants. As already mentioned, HOMO and LUMO levels of dopants relative to those of host materials are important in the case of trapping mechanisms. Since the doping mechanism has not been well understood and depends on materials, we must collect more information on dopants in search for more efficient and stable materials. Figure 14 shows some selected lifetime data measured consistently on the same devices. The filled circles indicate the bilayer device, and the open circles correspond to the trilayer device with a doped emitting layer (type I device in Fig. 11). These two lines show thermal degradation. The type I1 device of Fig. 11 (represented by open triangles), which behaves coulombically in the low-current-density region, shows thermal decay in the higher current region. The same tendency is also observed for polymer LEDs. It is tempting to conclude that charge-dependent degradation proceeds at low current density, while thermal degradation dominates in the high-currentdensity region. Since this conclusion is based on degradation phenomena, however, it still requires direct evidence.

228

YOSHIHARU SATO 1o5 1o4

10’ 1 1

10

100

1000

Drive Current Density [mA/cm2] FIG. 14. Lifetime data plotted as a function of the current density.

3. TEMPERATURE DEPENDENCE The previous subsection presented operational stability data at room temperature. When we consider practical applications for cars or mobile displays, we must consider the device characteristics at an elevated temperature and high humidity. Because the development of OLEDs for commercial products has just begun, few data are published or available. In this section, the thermal behavior of OLED devices is discussed based on a rather limited amount of data. First, the temperature dependence of EL characteristics was examined. The luminance and voltage of the CuPc-a-NPD- Alq trilayer device were measured at a current density of 15mA/cm2 at various temperatures. Luminance decreased as temperatures increased for the undoped cell. This result indicates temperature quenching of Alq itself. The device with a rubrene-doped Alq layer, on the other hand, exhibited almost constant luminance as the temperature changed (see Fig. 15a). It is known that rubrene does not show temperature quenching in some organic solvents. Quinacridone has also demonstrated a constant luminance, not affected by temperature. In doped OLEDs, the thermal properties of the fluorescent dopant determine the temperature dependence of EL emission. Therefore, OLED doping with fluorescent dyes that do not show temperature quenching is preferable for practical applications. Drive voltage, in both cases,

229

4 ORGANIC LED SYSTEM CONSIDERATIONS 1.2

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FIG. 15. Temperature dependence of luminance and voltage for the rubrene-doped and undoped devices.

decreased with increasing temperature, as shown in Fig. 15b, which suggests that some temperature-activating process determines carrier injection or transport. The activation energy in both devices showed almost the same values: 0.19 eV for the undoped and 0.14 eV for the rubrene-doped cells, which indicates there is no difference between temperature dependencies of charge injection or transport in the two devices. When a temperatureindependent dopant dye is used, an extra circuit to compensate temperaturedependent luminance is not necessary. The environmental test for display use requires storage stability at 85°C for 500 h. Under this condition, the most critical failure mode is the growth

230

YOSHIHARU SATO

of dark spots. Dark spots larger than 20 pm in diameter can be observed by the human eye. The Idemitsu group has reported that the acceleration coefficient of dark spots growth at 60°C compared with room temperature was estimated to be 40 times. Basically, the heat resistance of an OLED is governed by the T, of the organic materials of which it consists. In the device structure shown in Fig. 9b, the T, of the hole transport material is the lowest, compared with CuPc buffer and the Alq-emitting material; the T, of CuPc can be estimated at higher than 200°C using a relationship between T, and melting point (Tm), T,/T, = 213, and that of Alq was reported to be 175°C (Naito and Miura, 1993). Therefore, the development of a high- T, hole transport material is quite essential to obtaining OLEDs that can operate at an elevated temperature. The T, value is thought to dominate the migration-diffusion in organic films formed by small molecules during and after deposition. To elucidate this molecular migration, two kinds of bilayer samples were vacuum deposited: TPD (60 nm)-rubrene (15 nm) and ENPD (60 nm)-rubrene (15 nm) (Sato et al., 1998). Rubrene, as shown in Fig. 2, exhibits very weak fluorescence in the solid state; strongly fluorescent only in a dispersed state. These samples were fabricated keeping the substrate temperature at 20, 40, and 73°C to examine the thermal behavior of the organic films. At 20°C substrate temperature, both samples indicated blue photoluminescence peaks, which were assigned t o the hole transport materials, TPD and a-NPD, respectively. At 73"C, while a-NPD-rubrene sample exhibited almost the same PL spectra as that of 20°C sample, the TPD-rubrene sample showed an enhancement of rubrene PL peaks at 565nm and the TPD-related peaks disappeared. This complete change of the PL spectra can be attributed to a thermally induced migration of rubrene molecules into the TPD layer to form a dispersed state, because the substrate temperature of 73°C is higher than T, of TPD. These experimental facts show that T, is critical to obtaining a thermally stable bilayer structure. From the preceding experimental results, the OLED using a-NPD as a hole transport layer can be expected to have good thermal stability. To prove this thermal stability, the bilayer EL device with a-NPD was annealed at 100°C for 3 h (Sato et at., 1998). This annealing condition reflects process requirements such as encapsulation or wire bonding. I-V and L-V characteristics were observed as unchanged by this annealing. When TPD was used instead of a-NPD, complete crystallization took place, causing a growth of dark spots or electrical short in some cases. Operation of OLED at an elevated temperature (60°C) was reported to result in a half-decay time of 4000 h, compared to 7000 h at room temperature (Wakimoto et al., 1997b). An acceleration law for heat or current in

4 ORGANIC LED SYSTEM CONSIDERATIONS

231

operation has not yet been determined. The acceleration ratio can be used only to predict a shelf lifetime in terms of storage temperature. This experimental result suggests that thermally induced degradation was observed during operation; however, the degradation mechanism at an elevated temperature must be investigated more profoundly in the future.

111. Driving Scheme

1. DRIVING METHOD

As already seen, OLED luminance is proportional to current density. In the previous section, all the lifetime tests were carried out in the constantcurrent mode since this operation mode gives the prolonged half-decay time of the luminance. In the first part of this section, issues concerning constantcurrent vs constant-voltage operation are addressed. In terms of lifetime, usual operation is performed under constant-current conditions. In Section 11, the device with a buffer layer, high-T, hole transport layer and doped emitting layers such as ITO-CuPc-a-NPDA1q:rubrene-MgAg exhibited a dramatically improved operation stability in the constant-current mode. Figure 16 shows decay of both luminance and current density of this device in the constant-voltage mode: 8-V, 50-Hz pulse operation with duty ratio of 5% at an initial luminance of 100cd/m2. The

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1000 2000 3000 Operation Time [hour]

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FIG. 16. Operation characteristics of the trilayer OLED in the constant-Volga pulse mode.

232

YOSHIHARU SATO

half-decay time was 3500 h. Although most of the degradation results from the decrease of current density, indicating an increase of serial resistance of the device (see Fig. 16), this lifetime value was sufficiently improved to enable constant-voltage operation. In the constant-voltage operation, the voltage drop along the I T 0 lines must be considered. If we assume the following conditions: Whole display area driven in the four divided areas: upper and lower and right and left I T 0 sheet resistance of 15 O/square Fill factor 70% Percentage of ON pixels, 50% the size of the panel is plotted against current density, with AK the voltage drop between both ends of the I T 0 data lines, as a parameter (see Fig. 17). This plot shows that a 5-in. diagonal panel can be driven at 1 mA/cm2 with a voltage drop of 0.05 V between both ends. The margin of A V depends on the L-V characteristics of the device. Usually, luminance depends on the seventh to ninth power of voltage, a small change of voltage at I T 0 line gives rise to a great difference in luminance. Therefore, constant-voltage operation is limited to small size displays. One of the advantages of constant-voltage operation, however, is its compatibility with conventional driver IC chips. To realize an OLED panel with good luminance uniformity, we must consider constant-current operation. The situation is depicted in Fig. 18.

-5 25 20

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3

% 10 $ 5

0 0.1

1

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100

Cunent Density [mA/cm2 J FIG. 17. Estimation of voltage drop at both ends of I T 0 lines.

4 ORGANIC LED SYSTEM CONSIDERATIONS DATA lines (ITO)

233

equivalent circuit

SCAN

lines (cathode)

OLED

eR-

FIG. 18. Schematic illustration of constant-current driving taking IT0 resistance into account.

The problem is whether two pixels can be driven at the same current density or not. The simple driving circuit that ensures constant-current operation is comprised of an operational amplifier and a field-effect transistor (FET). The equivalent circuit involving serial resistance of I T 0 (RITo) is also included in Fig. 18. How large R,,, can be made is determined by the characteristics of the FET. A rough estimation shows that a 10-in. panel can be driven in the constant-current mode. Crosstalk, another issue for OLEDs is addressed in the passive matrix scheme (Braun, 1998). Figure 19 illustrates a diode-matrix made of OLEDs. When one data line ( I T 0 line) and scan line (cathode line) are selected, two other diodes (shown in gray) may be operated unintentionally due to the reverse current passed through the adjacent diodes. Forward and reverse I-V characteristics are very important in this sense. A typical J-V characteristics in both directions are shown in Fig. 20; a rectification ratio of 10’ to lo6 was observed. This rectification ratio can be improved using AlLi or LiF-A1 cathodes. However, leakage current due to pinholes in the organic layers increases reverse current, and may cause crosstalk. For practical applications, in addition to improvement of the fundamental I-V characteristics of OLEDs, the driving circuit must be considered. One possible scheme is shown in Fig. 21 where the capacitance associated with organic layers is taken into account (Okuda et al., 1997). When one scan line is selected, the other lines are reverse biased to suppress emission. In the

234

YOSHIHARU SATO

FIG. 19. Crosstalk in a passive matrix panel

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4 ORGANIC LED SYSTEM CONSIDERATIONS

I

235

\

act as capacitance suppress cross talk

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FIG. 21. Example of driving method of passive matrix panel for suppressing crosstalk.

second sequence, the unselected cell is first reset, since it has been reverse charged, inhibiting forward biasing. In the earlier lifetime tests, DC continuous operation has usually been adopted. It has been reported that a unipolar drive causes polarization of the device due to charge trapping, which affects operation stability. Negative or reverse bias effects have been investigated in conjunction with stability. An increase of drive voltage in the course of operation has been suppressed with constant current mode in the forward bias and reverse biasing in the constant-voltage mode (VanSlyke et al., 1996b). In Fig. 22, DC and AC

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100

150

200

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Operation Time [hour] FIG. 22. Effect of AC operation on the decay of luminance and voltage increase.

236

YOSHIHARU SATO

operation are compared in terms of both luminance decay and voltage increase. The waveform used in this comparison is shown in the inset in Fig. 22; constant-current pulses of 87 mA/cm2 followed by constant-voltage pulses of - l O V of 2ms each result in a duty ratio of 20% at a frequency of 100Hz. The experimental result shows that in AC operation, luminance decay was slower and voltage increase relaxed, although the starting voltage was higher. Please note that this sort of AC operation can be performed using the drive sequence shown in Fig. 21 to suppress crosstalk. 2. PASSIVE vs ACTIVEMATRIX

Although the first product is a passive matrix panel, active matrix operation must be considered for advanced applications. Both types of panels are shown schematically in Fig. 23. As already mentioned in the previous subsection, passive matrix operation requires dynamic operation; that is, high-current pulse operation. An active matrix drive ensures 100% duty operation of each pixel, leading to low-current and low-voltage operation. In the active matrix panel, the pixel thin film transistor (TFT) configuration must be considered at first. When we assume a capacitor connected to a source or drain of the TFT under the following conditions

passive matrix

active matrix

-DATA+

I

SCAN

i 1 frame diode matrix:multiplex operation

ply-Si TFTZ SCAN

1 fiame

continuous DC operation

FIG. 23. Schematic illustration of passive and active matrix panels.

4 ORGANIC LED SYSTEM CONSIDERATIONS

237

drive current density = 1 mA/cm2 drive voltage = 1OV frame frequency = 60 Hz SiNx capacitor of 100 nm the area of capacitor would be 25 times as large as the pixel area. Consequently, the capacitor should be connected to the gate of the driving TFT, and one more TFT is necessary to switch the first TFT. This fundamental configuration of pixel TFTs is shown in Fig. 23, which is the same as in active matrix driven inorganic EL (AMEL) (Arbuthnot et al., 1996). An example of pixel pattern using two a-Si TFTs is depicted in Fig. 24. There have been several efforts to fabricate an active matrix OLED by means of amorphous Si and poly-Si technology. Single TFT pixel structures have been fabricated without capacitive memory (Wu et al., 1997). TwoTFT structures have been reported to be applied to a 320- x 240-dot panel using poly-Si, which indicated a hole mobility from 16 to 22cm2/Vs (Hatalis et al., 1997). Four-TFT pixel structure was also proposed to obtain the uniformity of luminance introducing autozero operation (Dawson et al., 1998). Green monochrome and full color QVGA panels were fabricated by the TDK group for demonstration. In comparison with active matrix driving, passive matrix panels have the drawback that there is a decrease of power efficiency and an increase of operation voltage under high-duty-cycle operation. This situation is shown in Fig. 25, where typical L-V and luminous efficiency vs voltage curves are plotted. A practical level of luminance (100 cd/m2), was attained at 5 V. In COM

siNx+isi 1

-c SCAN

FIG. 24. Example of pixel structure with two TFTs and a capacitance.

238

YOSHIHARIJ SATO

A

1

t 0

*

A

0

2

4 6 8 Applied Voltage [v]

10

12

FIG. 25. Deterioration of EL characteristics in high-duty operation.

the case of 100 scan lines, each line must be driven at 10,000cd/m2, taking the duty ratio into account. The latter luminance was reached at 10 V, while luminous efficiency decreased from 5 to 21m/W. This decrease of power efficiency is fundamental to dynamic operation in passive matrix panels. The comparison between both driving schemes is summarized in Table 111. From the preceding arguments we must consider the number of scan

TABLE 111 PASSIVE vs A c n w MATRIX Passive matrix

Item Fabrication Technology Precision Cathode Performance Crosstalk Current Voltage Efficiency Stability Fill factor Size Cost

Active matrix

Conventional photolithography 5 Pm Patterned

poly-Si TFT process 1 Pm Uniform

Reverse bias 240mA/cm2 (duty = 1/240) >10v 1/2-1/3 of maximum efficiency Degradation due to high-duty operation

None 1 mA/cm2 (duty

70-80%

60 Yo poly-Si process limit Medium; system on panel

< 10 in. Low; driver IC?

=

100%)

5v Maximum efficiency Same as single device

4 ORGANIC LED SYSTEM CONSIDERATIONS

239

lines in practical application to decide which driving scheme will be suitable. Actually, this decision strongly depends on which application field we are aiming at. In terms of power consumption and lifetime, active matrix driving seems advantageous; low temperature poly-Si technology is the future trend for active matrix OLED panels.

IV. OLED Panels

Since the first announcement of the commercialization of an OLED display from Pioneer Corp. in the autumn of 1995, demonstration or test panel fabrication has been actively pursued, mainly in the Japanese industry. Some of the panels are listed in Table IV. Pioneer’s green monochrome panel has already been commercialized for car audio display dedicated to FM broadcast for traffic information. The pixel size is comparable to that of a 10-in. diagonal VGA display. TDK has demonstrated an active-matrixtype 5-in. QVGA full-color panel, using white-emitting OLEDs combined with color filters. Idemitsu and NEC have exhibited passive matrix QVGA full-color panels; the Idemitsu panel employed the CCM (color changing media) method and NEC has claimed to use the patterned organic dye approach. Stanley also has demonstrated segmented display intended for car use. In this section, OLED panel fabrication technology is reviewed from the materials, processes, and panel construction points of view.

1. MATERIAL ISSUES In molecular OLED technology, as already described, emission color is controlled by a combination of host and guest or dopant. Considering stability of the device, buffer materials and charge transport materials have been developed simultaneously. In Table V, materials used for OLED panels are listed in three columns: past, present, and future. IT0 has been not intensively investigated in the research of OLEDs. Sputtered or electron-beam-deposited I T 0 films on glass substrate (Corning 7059, 1737) have been conventionally used. These ITOs are not specially developed for OLEDs but for LCD applications. It is well known that the surface roughness, work function, and grain size of crystallites are quite different from sputtering to electron-beam methods. Since OLED characteristics are strongly affected by these IT0 surface properties, the first step of fabrication is to select a good (!) I T 0 substrate. Idemitsu’s commercial

TABLE IV OF OLED PANELS EXAMPLE

Item

h,

P

a

Application Panel size Number of dots Pixel pitch (pm) Drive scheme Duty ratio Driver Luminance (cd/m2) Emission color Voltage (V) Power (W) Lifetime (h) Other

Pioneer Car stereo 95 x 30mm 256 x 64 340 x 300 Passive 1/64 Six ICs 100 Green 9 0.5 (all ON) 5000

Commercial

Car navigation 5.2-in. diagonal 320 x 240 330 x 330 Passive 11120

150 Full color -

6 (all ON) 2000 Shadow mask

TDK Car navigation 5-in. diagonal 320 x 240 250 x 250 poly-Si TFT 1 poly-Si 150 Green 10 1 Aperture 57%

Idemitsu

NEC

320 ( x 3) x 240 90 x 300 (sub) Passive 1/240

Demonstration 5.7-in. diagonal 320 x 240 360 x 360 Passive i/240

-

-

80 Full color 20 5

20 Full color

CCM

Shadow mask

TV

5-in.diagonal

Stanley Demonstration 40 x 15mm 128 x 64 300 x 300 Passive 1/64 30-50 W, B, Or

241

4 ORGANIC LED SYSTEM CONSIDERATIONS TABLE V MATERIALS FOR OLED: PAST,PRESENT, AND FUTURE Item Anode Buffer Hole transport Electron transport Cathode Host: G-R Host: B G-dopan t R-dopant B-dopant

Past

Present

I T 0 (SP, EB) CUPC TPD Oxadiazole MgAg, Ca

LCD-grade I T 0 CuPc, PEDOT U-NPD

Ah AILi, LiF-A1

Ah

Ah

BAlq Coumarin DCM Perylene

DPVBi Quinacridone DCJTB DPVBi derivative

Future InZnO? Polymer? %> 120°C; starburst? New materials Moisture-insensitive New material? Wide-gap material Suppression of concentration quenching

sputtering target, InZnO, has been announced as providing a smoother surface than conventional ITO. The importance of an I T 0 buffer layer in obtaining low-voltage and stable devices was stressed in the previous section. The recent trend is to use a conductive polymer such as polyaniline (PANI) or 3,4-polyethylene-dioxythiophene (PEDOT), because the wet coating process is expected to mask the surface defects or inhomogenity of ITO. For full-color applications, transparent buffer material is preferable. Many possibilities for buffer materials remain. High-T, material is a key word for the hole transport layer. The current standard material, a-NPD, has a T, of 96°C. For higher thermal stability, a T, higher than 120°C is a target value. Starburst molecules are such candidates for future applications. If the thermal stability of the device is further improved, a wider range of applications or markets appears feasible. The family of oxadiazole compounds were studied as the electron transport material in the initial stage of OLED research. However, the efficiency of electron injection was not as good as with Alq. Therefore, Alq is still used also as an electron transport layer. In the field of OPCs, electron transport materials have long been searched for a positive charging photoreceptor. Thus, chemists in both fields are able to share interests. A nitrogencontaining electron-deficient material such as quinoxaline is an example. The cathode concept has changed drastically since the discovery of the effect of Li compounds at the cathode interface. With increasing comprehension of the electron injection mechanism at the cathode interface, more efficient and stable materials can be developed in the near future. While an electric double-layer model has been proposed (Hung et al., 1997), photo-

242

YOSHIHAR~J SATO

electron spectroscopy measurement revealed the shift of the vacuum level at the interface, which eases electron injection from the cathode into Alq (Ishii et al., 1998). As an emitting host material, Alq can cover the spectral range from green to red. Many materials have been developed in an attempt to replace Alq without success. Almq with ligand of 4-methyl-8-hidroxyquinoline is a modified Alq and has a higher quantum efficiency than Alq (Kido and Iizumi, 1998). Although there exist potentially many molecules, we must learn about Alq for the time being. On the other hand, a variety of materials have been used as a dopant. As a green dopant, coumarin is the first successful example for efficiency improvement. The second generation is quinacridone derivatives, which have demonstrated quite an improved durability. However, these derivatives exhibit strong concentration quenching, which required accurate control of deposition rate relative to that of the host, Alq. Therefore, easier-to-use dopants with less concentration quenching tendencies should be developed. In terms of concentration quenching, rubrene is an unusual molecule because of its sterically hindered structure. It gives yellow emission, which resembles that of ZnS:Mn. A future target for yellow emission is to increase luminous efficiency to 20 lm/W. For red emission, only a limited number of red fluorescent dyes have been reported. DCM has an orange hue, not red. To modify the DCM molecular structure under the concepts of the introduction of rigid substituents and planarizing of the skeleton has led to DCJTB, and Chen et al. (1997) reported that the luminous efficiency is 1lm/W and CIE color coordinate is (0.62,0.38).The improvement of color purity and luminous efficiency will be pursued for full-color applications, which are mentioned again in the latter part of this section. The situation is different for blue-emitting OLEDs, since Alq cannot be utilized. The first example of blue host material is one of the Alq family, a mixed-ligand type, is bis(2-methyl-8-quinolinolato)(g-phenylphenolato)aluminum (BAlq). BAlq doped with perylene gave blue emission, however, with stability inferior to the Alq system. 1,4-Bis(2,2-diphenylvinyl)biphenyl (DPVBi) was developed by the Idemitsu group and reported to have a high luminous efficiency of 6 lm/W when it was doped with a DPVBi analog. To summarize the present status of emitting materials, the lifetime data of each emission color are plotted as a function of initial luminance in Fig. 26. The green device is the most efficient and stable among the OLEDs. Yellow devices are compatible with the green device. For red, good stability data were published (Chen et al., 1997). The blue device is a bit behind these, however, and when we take into consideration the required luminance for blue in full-color applications, this performance locates within a practical region. Compared with the preceding devices, the performance of white-

4 ORGANIC LED SYSTEM CONSIDERATIONS

P

?

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~

.

243

~ 1 /

0

"

~

whit.

10'

loo0

100

Initial L u m i i e [ d m q FIG. 26. Review of lifetime data as a function of the initial luminance for RGBYW colors.

emitting device has been found to be quite low. White-emitting devices are important because they can be used as a backlight or directly viewable display. A combination of a white-emitting device with a color filter is one of the approaches to achieving a full-color OLED panel, which is described later. To realize white emission, there are mainly two approaches: one is to stack RGB layers and the other is to place RGB dopants in a single layer. Sato et al. (1997) demonstrated one of the stacking approaches using a combination of hole transport layer (HTL) doped with orange-fluorescent dye and a bluish-green emitting layer. The luminance of the white emitting device was 3000 cd/m2 at 250 mA/cm2 with a luminous efficiency of 0.5 lm/W at 100 cd/m2. CIE color coordinates (x = 0.34,y = 0.33) were obtained for this white-emitting device. These results proved that a simple device structure combining an orange-emitting HTL and a blue-emitting layer is an effective approach to achieve pure white emission. Doping of poly(N-vinylcarbazole) with RGB fluorescent dyes such as 1,1,4,4-tetraphenyl-1,3-butadiene, coumarin 6, and DCM was demonstrated to be an effective method for obtaining white emission (Kido et al., 1995). In the latter approach, a wet process based on polymers is more advantageous since it is difficult to control deposition rates of multisources in a vacuum codeposition process. In addition to these two approaches, new types of emitting materials such as zinc benzothiazole complexes were found to show broadening of EL spectra (Hamada et a!., 1996). Unfortunately, since EL spectra of these complexes lack a red portion, they must still be doped with red dye to move

244

Y OSH~HARUSATO

the CIE coordinate to (0.33, 0.33). The luminous efficiency was insufficient in any approach for practical applications. CIE color coordinates of the present status RGB materials are summarized in Fig. 27. For comparison, those of inorganic phosphors currently used for CRT are also indicated. Green and red OLEDs demonstrated compatible performance with inorganic phosphors. For blue emission, OLED gives a whitish-blue color since its spectral linewidth is relatively wide. Then, improvement in color purity and stability remains necessary for blue emission. The ultraviolet-emitting device based on polysilanes was reported, however, as unstable at room temperature with rapid degradation (Suzuki et al., 1998). A near-infrared emission device has been also proposed (Fujii et al., 1997).

2. PROCESS CONSIDERATIONS The fabrication processes of OLE11 passive matrix panels that are based on small molecules are 1. I T 0 patterning with etching 2. Formation of barrier ribs for cathode isolation 3. Cleaning of the I T 0 surface 4. Vacuum deposition of organic layers 5. Vacuum deposition of cathode 6. Encapsulation in a glove box 7. Interconnection with flexible printed cables (FPCs) 8. Assembly of the module

The most unique process in OLED fabrication is pattering of the cathode. When a metal electrode is patterned in an usual photolithographic method, the organic layers are subject to exposure to organic solvents or an inorganic base for developing and removing photoresists. These wet processes cause severe damage in OLED devices with dark spot formation. Therefore, cathode patterning of an OLED requires a dry processing. One of the methods is depicted in Fig. 28. In this method, prior to vacuum deposition of organic layers and cathodes, barrier rib structures are patterned in a inversely tapered form using nega-type photoresist (Nagayama et al., 1997). The width of the barrier ribs can be 20 pm and their height is around 4-5pm. These barrier ribs define the cathode lines rather well. During the deposition of organic layers and cathode, a cathode line can be

4 ORGANIC LED SYSTEM CONSIDERATIONS

0.9

245

c 520

530

0.8 l r Z x

...

..

.

..

Y 1:DCJTB

0.1

0.2 0.3

0.4

0.5

0.6 0.7 0.8

X

FIG.27. CIE color coordinates of typical OLEDs.

isolated from the other lines owing to the shadowing effect. This selfalignment process is quite effective to define the fine OLED pixels. The other fabrication steps are common in display technology. I T 0 etching is nothing special in OLED technology. The cleaning process is most important to reduce the number of dark spots. Vacuum deposition is used for forming organic thin films. In-line or continuous production vacuum systems must be specially designed for OLEDs, since there is no example of an organic deposition system for mass production. Coevaporation or doping control must also be developed. Encapsulation is determining the moisture resistance of the OLED panels. This process is usually performed in a glove box with moisture below 1 ppm. While inorganic EL panels utilized silicone oil for encapsulation, organic LED panels currently employ a drying agent such as barium oxide to protect them from ambient moisture. Other encapsulating schemes will be investigated in terms of both cost and performance. Driver ICs are installed on FPCs to make the whole module compact.

246

Y OSHIHARU SATO

t

FIG.28. Micropatterning process of cathode lines in the fabrication of a passive matrix panel.

3. MONOCHROME PANELS

The first commercial product from Pioneer Corp. is a passive matrix green monochrome display for traffic information systems. Specifications of this panel were already indicated in Table IV. This type of panel is fabricated using the cathode-pattering method, which has been described in the preceding section. It has a polarizer to enhance the contrast ratio up to 1OO:l at an ambient light level of 500 lux. The drive voltage is below 1OV with low power consumption; one-third of a back-lit STN-LCD, even when all the pixels are ON. A prototype of a green monochrome QVGA active matrix panel was demonstrated by TDK (see also Table IV). Other than dot matrix displays, segmented type or fixed pattern displays have also been shown by many companies (see Stanley in Table IV). This type of panel embodies the same concept as inorganic LED displays. In all the monochrome panels, the organic layers do not have to be patterned, but are uniformly deposited on the whole display area.

4 ORGANIC LED SYSTEM CONSIDERATIONS

247

4. MULTICOLOR AND FULL-COLOR PANELS A three-color (green, yellow, and red) display was exhibited at ICEL conference at Kitakyushu in 1997 (Arai et al., 1997). This display is based on a double emitting layer, rubrene-doped HTL and coumarin-6-doped Alq with color filters. It also seems to be developed for a car information system with 256 x 64 dots. This panel is an area-separated multicolor display. To extend a multicolor display means developing a full-color one. To realize a full-color panel is most exciting and important for future applications and the potential market. Requirements for RGB emission are summarized in Table VI, assuming 25% area for each subpixel (overall fill factor is 75%) and considering the balance of white emission. When we request the full-color OLED panel to compete with a TFT-LCD panel in terms of power consumption, operation at 5 V, 3 mA/cm2 is required, giving the required efficiency. Theoretical limits of RGB emission, assuming external quantum efficiencies of 5%0, should be compared with the required values (Tang, 1996). The bottom line of the lists present status of each emission color. Only the red emission has not satisfied the required efficiency. Several approaches to obtain full-color panels are described in the previous section. The patterned RGB method is an ideal way to realize a full-color panel, which has been actually fabricated by NEC group (see Table V). The CCM method, based on the down-conversion technique, utilizes a blue-emitting cell to excite green fluorescent and red fluorescent dyes, which are dispersed in a polymer matrix. In this approach, there is no need to pattern the organic layers; it can be uniformly deposited over the entire display region, The filtered-white method is a strategy similar to the “color by white” approach in inorganic EL devices. The uniform deposition of the organic layer is also the same as in CCM approach. Stacked-RGBcells is a simplifying idea that combines each emitting cell separately;

TABLE VI

FULL-COLOR REQUIREMENTS FOR RGB EMISSION Item

Blue

Green

Red

Luminance (cd/m2) Required efficiency (lm/W)

31 0.64

263

106 2.2

2.2 6.6 (Idemitsu)

12.0 10.0 (Pioneer)

Theoretical limit (lm/W) Attained efficiency (Ww)

5.5

5.7 0.5 (Kodak)

Remarks White balance At 5 V per 3 mA/cm2 1 W for 5-in. QVGA Maximum QE = 5% Blue chromaticity Red efficiency

248

YOSHIHAR~J SATO

patterned RGB

bluelCCM uniform blus w(a

(conmion 27%

average cunent density FIG. 29. Comparison of consumed current density for two full-color approaches.

however, it requires a complicated process to make interconnections to each cell. OLED cells with stacked RGB layers show a change of EL spectra depending on applied voltage, however, the change itself is gradual, not abrupt, and not enough to switch RGB color completely. The microcavity approach has been demonstrated to be “workable”; the spectral shift with viewing angle and the high cost associated with the use of interference filters are weak points. Finally, the photobleaching method is an innovative method that has not been fully demonstrated. The first two methods, which seem most feasible for practical application, are compared from the viewpoint of power efficiency in Fig. 29. The current densities required for each color were estimated using the values shown in Table VI. For both approaches, power consumption for the red subpixel is large. This situation reveals that the development of red fluorescent dyes is primarily important for full-color applications. In the patterned RGB method, improvement of the red portion directly contributes to a lowpower-consumption device; however, in CCM method, the inefficient green portion also remains.

V. Summary Since the preceding sections mainly describe molecular OLEDs, it is noteworthy, here, to compare molecular OLEDs with polymer LEDs. Table VII summarizes the present status of both technologies in terms of panel

4 ORGANIC LED SYSTEM CONSIDERATIONS

249

TABLE VII

SMALLMOLECULES vs POLYMERS Items

Small molecules

Polymer

Features Efficiency Voltage Lifetime Degradation

M W =400-800; T,= 60- 150°C

MW > 10,000; T,= 100-200°C 20 Im/W 3 V at 100 cd/m2 10,000h Photooxidation: by residual 0,, IT0 White-CF, cross-linking, ink-jet

Full-color approach Driving scheme Panel process Application Future Issues

16 Im/W 5 V at 100cd/m2 > 20,000 h Thermal:crystallization, interdiffusion White-CF, blue-CCM, shadow mask

Passive or active matrix Cathode isolation: dry process Car use (VICS, navigation), Backlight, mobile phone PDA, TV Nano-layer: interface control Nanocrystal? Continuous deposition, cost Control of impurity, thickness

performance and fabrication processes. Generally speaking, high T, is a strong merit of polymers, however, many of the high T, small molecules have been synthesized to make it less definite. Although efficiency has been considered as a strong pro for molecular OLEDs, progress in the efficiency of polymer LEDs makes it comparable to that of molecular OLEDs. The case for stability is also the same as efficiency in polymer LEDs has caught up with the performance of molecular OLEDs. But special attention must be paid to the difference of degradation mechanisms: crystallization in molecular OLEDs and photooxidation in polymer LEDs. Concerning the full-color approach, quite an exciting proposal to utilize ink-jet printing technology for polymer LEDs is worth watching in the near future (Bharathan and Yang, 1998). The difference in the fabrication process, that is, dry processing vs wet processing, must be investigated in terms of mass production and cost. The differences between both LED technologies suggest a different field of application. For OLED applications, there is another way to look at the situation: number of pixels vs size of panel (see Fig. 30). As already mentioned, the first commercial product or starting point of OLED panels is a car display (monochrome). The next target is PDA or car navigation displays, which would be full-color or multicolor QVGA panels. Both approaches, passive and active, have aimed at this area. QVGA panels seem to be located at the

250

YOSHIHARU SATO

W-LCD> 1 1 1 color

2

4

6

8

10

20

Panel Size [diagpnoll inch] FIG. 30. Trend in developing OLED displays depending on size and pixel number.

boundary of the two driving schemes. Mobile phone or segmented displays are smaller panels to be aimed at. Light source applications such as backlighting of LCDs could also be feasible. Low-voltage operation is quite attractive for mobile application. In Fig. 30, there are competing displays listed for each application. Table VIII summarizes the comparison of current FPD technologies. Features of OLEDs that are very positive are good efficiency, low-voltage operation, and quick response time. To compete with the already existing displays, it is necessary to reduce power consumption and driving voltage further, facilitating a low-cost full-color emissive display.

ACKNOWLEDGMENTS The author would like to thank Mr. H. Kanai for his collaboration throughout the whole work, and also Ms. S . Ichinosawa for her support of our OLED research by preparing organic materials. The author also appreciates K. Matsuda and M. Ichihara for their substantial contribution to the TFT work and Dr. Y. Yamaguchi for helpful discussion on thermal analysis of the device.

TABLE VIII COMPARISONOF OLEDs WITH THE OTHER FPDs ~

p3

cn -t

Item

LCD

PDP

View angle Efficiency(lm/W) Full color Size (in.) Voltage (V)

Improving 2-3 Excellent < 21 TFT: 2-5 B L lo00 20-60 ms View angle; large area

Excellent 1 Good 40 AC 90- 150 2-20 Efficiency; cost

Response (ps) Issues

~

~~

~

~

VFD

LED

Inorganic EL

FED

OLED

Excellent

Good? 0.2-7 Good Various DC 2-5 1

Excellent 2-4 Limited 2-20 AC 200 50 Blue phosphor

Excellent 7 Limited 5-20 DC lo00 1 Blue phosphor; low voltage

Excellent 5-10 Improving 2-20 DC < 10 1 Reliability; full color

0.8-14

Limited Small

Dc 10-40 10 Full-color resolution

252

YOSHIHARLI SATO

REFERENCES Adachi, C., Nagai, K., and Tamoto, N. (1995). Molecular design of hole transport materials for obtaining high durability in organic electroluminescent diodes. Appl. Phys. Lett. 66, 2679-2681, Arai, M., Nakaya, K., Onitsuka, O., Inoue, T., Codama, M., Tanaka, M., and Tanabe, H. (1997). Passive matrix display of organic LED’S. Synth. Met. 91, 21-25. Arbuthnot, L., Mendes, J., Sproull, W., Aguilera, M., Aitchison, B., King, C.N., Donly, G., Ipri, A,, Hsueh, F.-L., Stewart, R.G., Ahmed, S., Solanki, R., Keyser, T., Schlesinger, S., Becker, G., Kagey, and Spitzer, M. (1996). A 2000-lpi active-matrix EL display. SID 96 Dig., pp. 374-377. Bharathan, J., and Yang, Y. (1998). Polymer electroluminescent devices processed by inkjet printing: I. Polymer light-emitting logo. Appl. Phys. Lett. 72, 2660-2662. Braun, D. (1998). Crosstalk in passive matrix polymer LED displays. Synth. Met. 92, 107-113. Burrows, P. E., Bulovic, V., Forrest, S. R., Sapochak, L. S., McCarty, D. M., and Thompson, M. E. (1994). Reliability and degradation of organic light emitting devices. Appl. Phys. Lett. 65, 2922-2924. Chen, C. H., Tang, C. W., Shi, J., and KIubek, P. (1997). Improved dopants for organic electroluminescent devices. Macromol. Symp. 125, 49-58. Dawson, R. M. A,, Shen, Z., Furst, D. A., Connor, S., Hsu, J., Kane, M. G., Stewart, R. G.,Ipri, A,, King, C. N., Green, P. J., Flegal, R. T., Pearson, S., Barrow, W. A,, Dickey, E., Ping, K., Tang, C. W., VanSlyke, S., Chen, F., Shi, J., Sturm, J. C., and Lu, M.H. (1998). Design of an improved pixel for a polysilicon active-matrix organic LED display. SID 98 Digest, pp. 11-14. Do, L. M., Kim, K., Zyung, T., Shim, H. K., and Kim, J. J. (1997). In situ investigation of degradation in polymeric electroluminescent devices using time-resolved confocal laser scanning microscope. Appl. Phys. Lett. 70, 3470-3472. Fujihira, M., Do, L. M., Koike, A,, and Han, E. M. (1996). Growth of dark spots by interdiffusion across organic layers in organic electroluminescent devices. Appl. Phys. Lett. 68, 1787-1789. Fujii, A,, Ohmori, Y., and Yoshino, K. (1997). An organic infrared electroluminescent diode utilizing a phthalocyanine film. IEEE Electron. Devices 44, 1204- 1207. Hamada, Y., Sano, T., Shibata, K., and Kuroki, K. (1995). Influence of the emission site on the running durability of organic electroluminescent devices. Jpn. J . Appl. Phys. 34, L824-826. Hamada, Y., Sano, T., Fujii, H., Nishio, Y., Takahashi, H., and Shibata, K. (1996).White-lightemitting material for organic electroluminescent devices. Jpn. J . Appl. Phys. 35, L1339L1341. Han, E. M., Do, L. M., Yamamoto, N., and Fujihira, M. (1996). Crystallization of organic thin films for electroluminescent devices. Synth. Met. 273, 202-208. Hatalis, M. K., Stewart, M., Tang, C., and Burtis, J. (1997). Polysilicon TFT active matrix organic EL displays. Proc. S P I E 3057, 277. Hosokawa, C., Eida, M., Matsuura, M., Fukuoka, K., Nakamura, H., and Kusumoto, T. (1997). Organic multi-color electroluminescence display with fine pixels. Synth. Met. 91, 3-7. Hung, L. S., Tang, C . W., and Mason, M. G. (1997). Enhanced electron injection in organic electroluminescent devices using an Al/LiF electrode. Appl. Phys. Lett. 70, 152-154. Ishii, H., Sugiyama, K., Yoshimura, D., Ito, E., Ouchi, Y . , and Seki, K. (1998). Energy-level alignment at model interfaces of organic electroluminescent devices studied by UV photoemission: Trend in the deviation from the traditional way of estimating the interfacial electronic structures. IEEE J . Select. Top. Quantum Electron. 4, 24-33. Kanai, H., Ichinosawa, S., and Sato, Y. (1997). Effect of aromatic diamines as a cathode interface layer. Synth. Met. 91, 195-196.

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Kido, J., and Iizumi, Y. (1998). Fabrication of highly efficient organic electroluminescent devices. Appl. Phys. Lett. 73, 2721-2723. Kido, J., Shionoya, H., and Nagai, K. (1995). Single-layer white light-emitting organic electroluminescent devices based on dye-dispersed poly(N-vinylcarbazole). Appl. Phys. Lett. 67, 2281 -2283. Lee, S. T., Hou, X. Y., Mason, M. G., and Tang, C. W. (1998). Energy level alignment at Alq/metal interfaces. Appl. Phys. Lett. 72, 1593-1595. McElvain, J., Antoniadis, H., Hueschen, M. R., Miller, J. N., Roitman, D. M., Sheats, J. R., and Moon, R. L. (1996). Formation and growth of black spots in organic light-emitting diodes. J . Appl. Phys. 80, 6002-6007. Nagayama, K., Yahagi, T., Nakada, H., Tohma, T., Watanabe, T., Yoshida, K., and Miyaguchi, S. (1997). Jpn. J . Appl. Phys. 36, L1555-Ll557. Naito, K., and Miura, A. (1993). Molecular design for nonpolymeric organic dye glasses with thermal stability: Relations between thermodynamic parameters and amorphous properties. J . Phys. Chem. 97, 6240-6248. Nakada, H., and Thoma, T. (1996).Applications oforganic LEDs. In Inorg. and Org. ELIEL %Berlin, R. H. Mauch and H.-E. Gumlich, Eds., pp. 385-390. (Wissenschaft und Technik Verlag, Berlin.) Okuda, Y., Ochi, H., Tsuchida, M. (1997). Driving circuit for organic LED dot-matrix display. Proc. 10th Annu. Meeting LEOS 2, 362. Papadimitrakopoulos, F., and Zhang, X. M. (1997). Environmental stability of aluminum tris(8-hydroxyquinoline)(Alq3) and its implications in light emitting devices. Synth. M e t . 85, 1221-1224. Sato, Y., and Kanai, H. (1994). Stability of organic electroluminescent diodes. Mol. Cryst. Liq. Cryst. 253, 143-150. Sato, Y., Ichinosawa, S., and Kanai, H. (1996). Improved stability of doped organic electroluminescent diodes. In Inorg. and Org. ELIEL 96 Berlin, R. H. Mauch and H. E. Gumlich, Eds., pp. 255-258. (Wissenschaft und Technik Verlag, Berlin.) Sato, Y., Ogata, T., Ichinosawa, S., and Murata, Y. (1997). Characteristics of organic EL devices with new dopants. Synth. Met. 91, 103-107. Sato, Y., Ichinosawa, S., and Kanai, H. (1998). Operation characteristics and degradation of organic electroluminescent devices. IEEE J. Select. Top. Quantum Electron. 4, 40-48. Scott, J. C., Kaufman, J. H., Brock, P. J., DiPietro, R., Salem, J., and Goitia, J. A. (1996). Degradation and failure of MEH-PPV light-emitting diodes. J. Appl. Phys. 79,2745-2751. Shi, J., and Tang, C. W. (1997). Doped organic electroluminescent devices with improved stability. Appl. Phys. Lett. 70, 1665-1667. Staring, E. G., Berntsen, A. M., Romme, S. T. R., Rikken, G. L. J., and Urbach, P. (1997). On the photochemical stability of dialkoxy-PPV; a quantitative approach. Phil. Trans. R. Soc. London A 355, 695-706. Stevens, B., and Algar, B. E. (1968). The photoperoxidation of unsaturated organic molecules. I. Relaxation and oxygen-quenching parameters of the sensitizer singlet state. J. Phys. Chem. 72, 2582-2587. Suzuki, H., Hoshino, S., Yuan, C-H., Fujiki, M., Toyoda, S., and Matsumoto, N. (1998). Near-ultraviolet light-emitting diodes based on n-conjugated linear silicon-backbone polymers. IEEE J. Select. Top. Quantum Electron. 4, 129-136. Tamoto, N., Adachi, C., and Nagai, K. (1997). Electroluminescence of 1,3,4-oxadiazole and triphenylamine-containing molecules as an emitter in organic multilayer light-emitting diodes. Chem. Mater. 9, 1077-1085. Tang, C. W., (1996). Organic electroluminescent materials and devices. In$ Display 10, 16-19. Tang, C. W., and VanSlyke, S. A. (1987). Organic electroluminescent diodes. Appl. Phys. Lett. 51, 913-915.

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Tang, C. W., VanSlyke, S. A,, and Chen, C. H. (1989). Electroluminescence of doped organic thin films. J . Appl. Phys. 65, 3610-3616. VanSlyke, S. A,, Bryan, P. S., and Tang, C. W. (1996a). Blue organic light emitting devices. In Iizorg. cind Org. ELIEL 96 Berlin, R. H. Mauch and H.-E. Gumlich, Eds., pp. 195-200. (Wissenschaft und Technik Verlag, Berlin.) VanSlyke, S. A,, Chen, C. H., and Tang, C. W. (1996b). Organic electroluminescence devices with improved stability. Appl. Phys. Lett. 69, 2160-2162. Wakimoto, T., Fukuda, Y., Nagayama, K., Yokoi, A,, Nakada, H., and Tsuchida, M. (1997a). I E E E Truns. Electron. Devices 44, 1245- 1248. Wakimoto, T., Ochi, H., Kawami, S., Ohata, H., Nagayama, K., Murayama, R., Okuda, Y., Nakada, H., Tohma, T., Naito, T., Abiko, H. (1997b). Dot-matrix display using organic light-emitting diodes. 1.SID, 5-3, 235-240. Wu, C. C., Thciss, S. D., Gu, G., Lu, M. H., Sturm, J. C., Wagner, S., and Forrest, S. R. (1997). Integration of organic LED'S and amorphous Si TFT's onto flexible and lightweight mctal foil substrates. ZEEE Electron Device Lett. 18, 609. Yamamori, A,, Adachi, C., Koyama, T., and Taniguchi, Y. (1998). Doped organic light emitting diodes having a 650-nm-thick hole transport layer, Appl. Phys. Lett. 72, 2147-2149. Zyung, T., and Kim, J. J. (1995). Photodegradation of poly@-phenylenevinylene)by laser light at the peak wavelength of electroluminescence. Appl. Phys. Lett. 67, 3420-3422.

.

SEMICONDUCTORS AND SEMIMETALS VOL. 64

CHAPTER 5

Molecular Organic Light-Emitting Devices V. Bdovik. P. E. Burrows. and S. R . Forrest CENTER FOR PHOTONICS AND OPIOELECTRONIC MATERIALS (POEM) OF ELECTRICAL ENGINEERING AND THE PRINCETON MATERIALS INSTITUTE DEPARTMENT

~JNIVHRSITY PRINCETIIN

PRINCETON. NEWJERSEY

1. OVERVIEW . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1 . Historical Perspective . . . . . . . . . . . . . . . . . . . . . . . . 2. WhyMakeOLEDs? . . . . . . . . . . . . . . . . . . . . . . . . I1. THEORY OF OPERATION OF ORGANIC LIGHT-EMITTING DEVICES . . . . . . . . 1. Conventional OLED Structures . . . . . . . . . . . . . . . . . . . 2. Transport Properties of OLEDs: General Considerations . . . . . . . . . 3. Characteristics of Single-Organic-Layer OLEDs . . . . . . . . . . . . 4. Characteristics of Heterojunction OLEDs . . . . . . . . . . . . . . . 111. IMPROVED OPERATION OF ORGANIC LIGHT-EMITTING DEVICES . . . . . . . . 1. Modifying Emission Color and Eflciency by Doping . . . . . . . . . . . 2. Degradation Processes . . . . . . . . . . . . . . . . . . . . . . . 3. Increasing Eflciency Using Shaped Substrutes . . . . . . . . . . . . . 4. FIexihleOLEDs . . . . . . . . . . . . . . . . . . . . . . . . . . 5 . Transparent OLEDs . . . . . . . . . . . . . . . . . . . . . . . . 6 . Inverted OLEDs . . . . . . . . . . . . . . . . . . . . . . . . . . IV . COLOR-TUNABLE OLEDs IN FULL-COLOR DISPLAYS. . . . . . . . . . . . 1. Side-by-Side Patterning of Red. Green. and Blue OLEDs . . . . . . . . . 2. Color Passband Filtering of White OLEDs . . . . . . . . . . . . . . . 3. Wavelength Down-Conversion . . . . . . . . . . . . . . . . . . . . 4. Microcuvity-Filtered White OLEDs . . . . . . . . . . . . . . . . . . 5 . Stacked Color-Tunable OLEDs . . . . . . . . . . . . . . . . . . . V.SUMMARY. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . REFERENCES. . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Overview

1. HISTORICAL PERSPECTIVE

Efficient electroluminescence (EL) from an organic solid was first demonstrated in the early 1960s in thick (50.p m to 1-mm) crystals of anthracene

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(Helfrich and Schneider, 1965; Pope et al., 1963). Although the EL quantum efficiency of those early devices reached X%, they were impractical due to their high operating voltages (50-1000 V). In the following two decades, a number of studies of carrier injection and transport through both pristine (Basurto and Burshtein, 1975; Kalinowski et al., 1976; Kao and Hwang, 1981; Lohmann and Mehl, 1969; Williams and Schadt, 1970) and doped (Kawabe et al., 1971; Schwob et al., 1970; Wakayama et a!., 1974) organic crystals ensued. Building practical organic light-emitting devices (OLEDs), however, remained a challenge until 1987 when the first vacuum-deposited OLEDs with the active layer thicknesses of only a few hundred angstroms and low operating voltages (< 20 V) were reported (Tang and VanSlyke, 1987).

Subsequent to the demonstration of efficient OLEDs based on vacuumdeposited molecular materials, Burroughes et al. (1990) showed that polymeric thin films could also be employed for use in efficient electroluminescent devices. Since these early demonstrations, widespread interest in bright and efficient OLED-based flat panel displays has resulted in rapid growth of the field (Greenham and Friend, 1995; Rothberg and Lovinger, 1996; Sheats et a/., 1996). The operating voltage, brightness, and the range of colors of present-day OLEDs is adequate for most video display applications (Tang, 1996). Indeed, the first OLED-based displays have become commercially available (Nakada and Tohma, 1996; Wakimoto et al., 1996).

2. WHY MAKEOLEDs? To become viable candidates for use in flat panel displays, OLEDs must demonstrate high brightness, high efficiency (i.e., low operating voltage and current), good color saturation, and long lifetime. Satisfactory performance for indoor and portable display applications requires a brightness of 100 cd/m2 at an operating voltage between 5 and 15 V and a continuous operational lifetime of at least 10,000 h. The AlqJdiamine heterojunction OLEDs discussed in Subsection 4 of Section I1 exceed the requirements, as demonstrated by the development of the first commercially available Alq,-based passive matrix monochrome displays (Nakada and Tohma, 1996; Wakimoto et al., 1996). These greenemitting OLEDs are just the first of many structures that can emit over the entire visible spectrum. Molecular engineering should lead to significant improvements in both luminescent and conducting properties of this class of materials, since the large number of luminescent molecules presently available is small when compared to the enormous potential variety of organic molecular compounds.

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OLED brightnesses of up to 140,000cd/m2 (Kido and Iizumi, 1998), power efficiencies up to 31 lm/W (Baldo et al., 1999a; Schoo et al., 1998; Tang, 1996), low operating voltages from 3 to 10 V (Schoo et al., 1998), and operating lifetimes exceeding 50,000 h (Tang, 1996) have been reported. Saturated-color devices have been demonstrated, spanning almost the entire visible spectrum. In addition, the overall device thickness of less than a micrometer allows for considerable mechanical flexibility, leading to the demonstration of bendable molecular OLEDs grown on ITO-coated polyester substrates (Gu et al., 1997a). Together with the earlier demonstration of flexible polymer OLEDs (Gustaffson et al., 1993; May, 1996), these devices create the potential for realizing roll-up or conformable displays mounted on curved surfaces. Vacuum or vapor phase (Baldo et al., 1997) deposition on plastic sheets enables large-scale production of large area OLEDs via roll-to-roll processing. The successful development of such a process could result in a very low cost manufacture of efficient, flexible, flat panel displays. In this chapter, we discuss the fabrication, physics, and state of the technology of vacuum deposited molecular (i.e., nonpolymeric) OLEDs. We describe the theory of operation of single layer and heterojunction devices, emphasizing carrier transport through the organic layers and injection of carriers at contacts, and relate them to the general properties of commonly employed active organic and electrode materials. With the understanding of the physical processes that govern the operation of OLEDs we can modify their structure to achieve improved performance. Furthermore, the unique optical characteristics of organic materials lead to a number of novel light emitting devices also described in the chapter. Since the main goal of this technology is to develop full-color OLED-based pixelated displays, we conclude the chapter by describing several different approaches to solving this problem.

11. Theory of Operation of Organic Light-Emitting Devices 1. CONVENTIONAL OLED STRUCTURES The simplest OLED configuration is the single organic layer device (shown in Fig. la) consisting of an electron transport or hole transport organic layer (ETL or HTL, respectively) sandwiched between an anode and a cathode. Besides serving to transport carriers, the organic thin film also acts as the electroluminescent layer (LL). Unfortunately, the mobility of electrons and holes in most organic materials are considerably different, leading to exciton formation in the vicinity of one of the two contacts. Since

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I

Substrate

I

Substrate

J

FIG. 1. Cross-sectional diagrams of (a) a single-organic-layer OLED, (b) a double-heterojunction OLED, and (c) and (d) single-heterojunction OLEPs. The hole transport layer (HTL), electron transport layer (ETL), luminescent layer (LL), anode and cathode are indicated.

nonradiative exciton recombination, or quenching, is enhanced at the electrode-organic interface, the structure in Fig. l a typically exhibits a low quantum efficiency. Furthermore, a large percentage of the higher mobility carriers can pass through the organic without forming excitons, also decreasing device efficiency. The luminescence efficiency is significantly improved in heterojunction OLEDs formed by sequential deposition of two or more organic layers (Adachi et al., 1991). Figure l b shows the general structure of the doubleheterojunction OLED where the functions of the ETL, HTL, and LL are performed by separate organic materials. In this device, the transport layers are chosen such that the mobility of one carrier type is much higher than the other. Most carriers are then confined in the LL, where they form excitons, which radiatively recombine. In single-heterostructure OLEDs, the functionality of the ETL and LL (Fig. lc) or the HTL and LL (Fig. Id) is combined into a single organic layer. Indeed, early low operating voltage OLEDs (Tang and VanSlyke, 1987) were based on aluminum tris(8-hydroxyquinoline) (Alq,)/N,N'-diphenyl-N,N'-bis(3-methylphenyl)-1,1'-biphenyl-4,4-diamine(TPD) single heterostructures, where Alq, served as both ETL and LL. In the Alq,-TPD single heterostructure (Fig. 2) holes are injected from TPD (Stolka et al., 1984) into Alq,. The small hole mobility in Alq, (Kepler

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1.o

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0.8

3

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Y

0.6 .-3 u)

c

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0.4

c L 9)

I

3 0.2

I

I

I

IT0

0.0 400

500

600 700

i

800

Wavelength [nm]

EL Ught

FIG. 2. The photoluminescence (PL) spectrum of an Alq, thin film and the electroluminescence (EL) spectrum of an Alq,-TPD OLED sketched in the figure (from Bulovic, 1998).

et at., 1995) results in a short diffusion length of -30-100A (Lin et al., 1997); hence excitons are formed near the ETL-HTL interface. Figure 2 shows the EL spectrum of Alq,-TPD OLEDs and the photoluminescence (PL) spectrum of Alq,. The similarity between the two is evident, indicating that the EL emission originates from the radiative recombination of Alq, excitons. The exciton recombination region near the ETL-HTL interface is known as the “recombination zone.” Its width is determined by the diffusion length of carriers and excitons in the emitting material. It has been found (Garbuzov et al., 1996a) that exciton diffusion lengths in amorphous thin films such as Alq, are approximately 100-300 Hence, in optimal devices, this is on the order of the Alq, layer thickness. A clear demonstration of the extent of the recombination zone was provided by Adachi et at. (1991) using a single-heterostructure OLED, as shown schematically in Fig. 3. For this device, the 500-A-thick hole transport and light-emitting layer consisted of naphthostyrylamine (NSD), and the 500-A-thick ETL consisted of an oxidiazole derivative (structural formulas for these materials are also shown in Fig. 3). To determine the

A.

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m

FIG.3. (Top) Chemical structural formulas for the organic molecules used to determine the width of the recombination region. (Center) Schematic of the doped single-heterostructure OLED. (Bottom) Electroluminescence spectra of six doped single-heterostructure OLEDs, with the positions (A-F) of the dopant layers indicated in the diagram of the OLED (from Adachi et al.. 1991.

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extent of the recombination zone, a series of devices was fabricated, each with a 10-8L-thick slab of the luminescent layer doped with a highly emissive chromophore, a thiadiazole derivative. In each device, the doped section was grown at a different distance from the ETL-HTL interface. The emission spectra of these devices were measured with the results shown in the figure. When the doped section is located within 100 A of the heterointerface, the emission peaks at 590 nm (spectra labeled A and B), the characteristic of the emission from the thiadiazole derivative. However, when the dopant is moved farther from the interface (spectra C-F), the emission peak shifts to 500 nm due to recombination in the undoped NSD. This clearly indicates that nearly all radiative recombination occurs in the HTL within 1OOA of the ETL-HTL interface. Similar conclusions about the width and position of the recombination zone have been confirmed by measurements of exciton dynamics in Alq, and Alq,-TPD interfaces (Kalinowski et al., 1997; Tang et al., 1989), where most excitons radiatively recombine in the ETL within 400 A of the heterointerface. In Alq,-diamine OLEDs, the distribution of excitons within the recombination zone is also weakly dependent on the drive current (Bulovic et al., 1998a). Undoubtedly, these conclusions can be generalized to understand almost all small molecular weight, vacuum-deposited OLEDs.

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2. TRANSPORT PROPERTES OF OLEDs: GENERAL CONSIDERATIONS The thin organic layers of vacuum-deposited molecular or polymer OLEDs are typically amorphous, in contrast to most work on electroluminescence of crystalline organics (Helfrich and Schneider, 1965). The amorphous structure leads to a reduction in quenching from internal conversion processes present in strongly coupled, crystalline organic systems (that have many phonon modes associated with the crystal lattice) and a consequent increase in the radiative recombination efficiency of Frenkel excitons (Pope and Swenberg, 1982). Efficient generation of excitons, however, also depends on the efficiency of charged carrier injection and transport through the organic layers, as well as the likelihood of forming an exciton that will radiatively recombine. Due to the low intermolecular coupling and disorder of amorphous materials, their carrier mobilities are lower than in most crystalline solids. Consequently, the conduction mechanism typically does not follow Ohm’s law but rather is often space charge limited, influenced by the presence of traps, or hopping (see part b of Subsection 3 in Section 11) (Kao and Hwang, 1981). Carrier trapping in diamines such as TPD has a negligible effect on

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its conduction properties, and hole mobilities in these materials are of the order of lo-, cm2/V s (Stolka et al., 1984). In Alq,, however, a high-density distribution of trap states from 0.25 to 0.13eV (Forsythe et al., 1998) significantly affects charge transport, lowering carrier mobilities. The electron mobility in Alq, is 10-4cm2/Vs at an applied electric field of lo6 V/cm (Chen and Liu, 1997), with hole mobilities at least two orders of magnitude lower (Kepler et al., 1995). The lower mobility of holes is responsible for their short diffusion length (Lin et ul., 1997). Since most of the following discussion of carrier transport in organics focuses on Alq,, in Fig. 4 we summarize the room temperature electron and hole mobilities in this material as reported by three different groups (Chen and Liu, 1997; Hosokawa et al., 1992; Kepler et al., 1995). The most detailed study is that of Chen and Liu (1997:), which shows the dependence of mobility on the electric field. The carrier mobility in most organic materials is found to be dependent on both the electric field ( E ) and temperature (T) (Gill, 1972) according to

-

-

where k is the Boltzmann’s constant, L3, = kTo is the activation energy at zero electric field corresponding to temperature To, po, and are the zero 1

.

1

’

1

,U Kepler, et el. (e,h+)

104 0, , 400 600

,

, 800

,

,

. ,

. , . , .

1000 1200 1400 1600 1800

Eln [v/cmI1n FIG.4. Summary of measurements of electric-field-dependent, room-temperature electron (solid symbols) and hole (open symbols) mobilities in Alq, (data from Chen and Liu, 1997, circles; data from Kepler et al., 1995, squares; data from Hosokawa rt a/., 1992, triangles).

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263

field electron mobilities, l& = 1/T - l/To, and and M are constants. This general trend agrees with measurements of Alq, electron mobility from Chen cm2/V s and M = 4.1 x lo-, (cm/V)''2. = 1.4 x and Liu (1997), with We note that thin films of organic materials are designated as either hole or electron transporting. Due to their typically low electron affinities, the majority of organic materials are hole-transporting solids. Unlike the situation for inorganic semiconductors such as silicon, controlled doping of molecular organic thin films to change the polarity of carrier transport has not been convincingly demonstrated. Injecting electrons into low-electron-affinity organic thin films is most efficiently done using low-work-function metal contacts. The stability of such cathodes tends to be poor, although they can be somewhat improved by using coevaporated or layered metals (Tang et al., 1989; Kido et al., 1993) such as Mg:Ag or Li:Al. Also, higher-work-function materials can be used as the cathode layer when defect states (Bulovii: et al., 1997; Parthasarathy et al., 1998) or dipole layers (Campbell et al., 1996b) are formed at the organic-electrode interface. Opaque metal cathodes used in a conventional OLED form the top layer (c.f. Fig. 1). The electroluminescence is observed through the transparent anode, which typically consists of a conducting oxide, such as indium tin oxide (ITO) grown on a transparent glass substrate. Other conducting oxides such as GaInO, and ZnInSnO have also been shown to efficiently inject holes (Cava et al., 1994; Phillips et al., 1995a, 1995b). The low refractive index of GaInO, (1.65) and its high transparency (compared to ITO), however, may be important in some applications (Dodabalapur, 1997). 3. CHARACTERISTICS OF SINGLE-ORGANIC-LAYER OLEDs

The current-voltage (I-V) characteristics of an OLED are determined by both the barrier to the injection of carriers into the organic layer (injection limited I-V, Subsection a in this subsection), and the transport of carriers through the thin film bulk (transport limited I-V, Subsection b in this subsection). In this section, we discuss the relative influence of these two processes on the I-V behavior of single layer OLEDs which are conceptually simple to analyze. In Subsection 4 of Section I1 we then apply our results to the more complex but higher efficiency heterojunction OLED. a.

Carrier Injection at Organic- Electrode Interfaces

The difference between the Fermi level of the cathode (anode) and the lowest unoccupied molecular orbital (LUMO) [highest occupied molecular

v. BULOVICET AL.

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orbital (HOMO)] band of the organic thin film forms a barrier to the injection of electrons (holes) into a single layer OLED. Modeling of the I-V behavior of these devices has shown that for barriers larger than 0.4 eV, the current flow is primarily determined by the efficiency of carrier injection at the contacts (Campbell et al., 1998). This is known as the injection-limited regime. The energy level offset at the organic-electrode interface can be tailored by choosing an electrode with lower work function for more efficient injection of electrons, and vice versa for holes. Furthermore, inserting an oriented dipole layer between the metal electrode and the organic thin film (Campbell et al., 1996b) can change the injection barrier, as shown in Fig. 5. It has been previously shown that the Schottky barrier energy at the Ag-polymer interface can be changed by more than 1 eV when a selfassembled monolayer of an alkyl thiol is attached to the Ag surface prior to polymer deposition (Campbell et al., 1996b). The -2D dipole moment of the oriented thiol molecules was shown to be responsible for this barrier tuning. Similarly, oxygen and hydrogen plasma treatment of the I T 0 surface changed the anode work function (Wu et al., 1997) by modifying the surface charge distribution. It has also been proposed that protonation of I T 0 by nitric acid vapor gives rise to an ionic double layer (adsorbed dipole) that increases the effective electrode work function by -0.8 eV (Niiesch, 1998). From these experiments, it is evident that a dipole layer at the organicelectrode interface can influence the energy level alignment, and therefore the carrier injection efficiency. As interface dipoles are likely to form at most organic-metal and organic-organic interfaces, they must be taken into account when designing N

€4

1

LUMO

4

I

HOMO

4 Metal

Dipole layer

Organlc

FIG. 5. Schematic energy level diagram of metal-dipole layer-organic interface. The dipole layer thickness is on the order of 1 nm. The dipole layer can either decrease (a) or increase (b) the electron Schottky energy barrier.

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efficient OLEDs. Recent spectroscopic studies of organic interfaces have begun to link the interface chemistry and morphology to electronic structure and the electrical characteristics of these contacts. Photoemission and inverse photoemission spectroscopy (Rajagopal et al., 1998; Sugiyama et al., 1998), internal photoemission (Campbell et al., 1996a), electroabsorption (Campbell et al., 1996a), and photocurrent measurements (Niiesch, 1998) are among the recently used spectroscopic techniques. In the early heterojunction OLEDs, hole injection into the HTL was improved by depositing a thin layer of copper phthalocyanine (CuPc) (Tang, 1982), polyaniline (PANI) (Yang and Heeger, 1994; Yang et al., 1995), or 3,4,9,10-peryIenetetracarboxylicdianhydride (PTCDA) (Rajagopal et al., 1998) onto ITO. Although formation of a dipole layer could be responsible for the improved hole injection efficiency, it has also been proposed that organic overlayers introduce additional energy levels at the organic-IT0 interface, which can facilitate hole injection (Tang, 1982). Defect states at the organic-metal interface undoubtedly dominate the injection properties of electrodes. In Mg:Ag-Alq,-Mg:Ag devices, surface states result in the asymmetric I-V response of these otherwise “symmetric” devices (see Fig. 6). A larger current in “forward bias” (bottom electrode positive) is clearly apparent, suggesting a different injection efficiency for the top and bottom Mg:Ag-Alq, interfaces. The log-log plot of these I-V characteristics in the inset of Fig. 6 shows that under forward bias the current is two to three orders of magnitude larger than under reverse bias. Since previous studies show that Alq, is preferentially electron conducting (see Fig. 4), asymmetry in the I-V characteristics implies that the top electrode is a more efficient electron injector than is the botton electrode. This is attributed to differences between contact interface chemistries as well as to their roughnesses. Due to the large latent heat of condensation, vapor-deposited Mg and Ag atoms are presumed to chemically react with the underlying Alq, as they thermalize on the film surface, forming interfacial defects that lower the effective contact energy barrier. In contrast, such reactions are significantly reduced when the relatively low sublimation temperature organic molecules are deposited onto Mg:Ag forming the bottom contact (Bulovii: et al., 1997). Indeed, photoemission spectroscopic investigations of Mg- Alq, and Alq,-Mg interfaces indicate differences in morphology and the density and energies of electron gap states (Rajagopal and Kahn, 1998). Mg deposited onto Alq, diffuses into the organic film, forming interface states that can facilitate electron injection, in contrast to the abrupt interface formed by deposition of Alq, onto Mg. The asymmetry of the organic-metal contacts due to the sequence of layer deposition is a general phenomenon (previously seen in spectroscopic

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2.0

1.5

-

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t

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6

0.5

-

1

1

-10

.

1

.

1

.

1

0 5 Voltage M

-5

.

1

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FIG. 6. Current-voltage characteristics of an Mg:Ag-450-A Alq,-Mg:Ag device with 0.25 mm2 top electrode area. Voltage is applied to the bottom electrode while the top electrode is grounded. Inset: Log-log plot of the forward- and reverse-biased I-V characteristics (from BuloviC et al., 1997).

measurements of PTCDA-In interfaces; Hirose et al., 1996). Physical origins of the asymmetry are pictorially represented in Fig. 7. At the top contact, interface states can facilitate carrier injection by replacing the large contact barrier with a series of smaller energy “steps” more easily surmounted by the injected electron. Similar defect states in CuPc damaged by I T 0 sputter deposition are postulated lo facilitate injection of electrons into metal-free cathode, transparent OLEDs (Parthasarathy et al., 1998). Formation of interface states may also be responsible for the lower operating voltage of Alq,-diamine heterojunction OLEDs with an A1-LiF cathode (Hung et al., 1997). Here, deposition of a “hot” LiF layer can generate surface states in Alq, that facilitate electron injection from Al. Alternatively, it has been proposed that LiF forms an oriented dipole layer that improves injection efficiency (Hung et al., 1997).

5 MOLECULAR ORGANIC LIGHT-EMITTING DEVICES

**..-.....,.....*..*

267

I

FIG. 7. Schematic diagram depicting formation of interface states at the top Mg:Ag-Alq, interface, in contrast to the less defective bottom interface (from Bulovit, 1998).

The forward-biased I-V characteristics in the inset of Fig. 6 follow the power law dependence typical of trap-limited conduction in organic materials (see Subsection b of this subsection). This suggests that the current conduction is transport limited, implying that the top Mg:Ag cathode efficientlyinjects electrons. In these devices, contributions of hole transport to the total current is small due to the low hole mobility in Alq,. Furthermore, in the absence of interface dipoles and surface-state Fermilevel pinning, the barrier to hole injection at the bottom electrode is in excess of 2 eV. Here, we assume that the work function, a, of Mg:Ag is equal to that of Mg (0= 3.7eVj, and the Alq, HOMO energy is at 5.8 k 0.2 eV below the vacuum level (Hamada et al., 1993; Rajagopal et al., 1998; Sugiyama et al., 1998). Similarly, electron transport in Mg:Ag-Alq,-ITO-glass devices is also dominant under forward bias ( I T 0 positive) (Campbell et al., 1998). In reverse bias, the current is four orders of magnitude smaller since untreated I T 0 cannot efficiently inject electrons into Alq,. The influence of the bottom electrode on the I-V characteristics of single layer Alq, devices is shown in Fig. 8. Here, I T 0 (0N 4.7eV), Ag (a = 4.3 eV), and Mg (0= 3.7 eVj were used as bottom electrodes on top of which a 500-a-thick Alq, film was deposited. Circular top electrodes, 1mm2 in area, were defined by shadow masking the substrate while evaporating a 500-A-thick MgAg (mass ratio 9:l) electrode layer and a 1000-A-thick Ag cap. The Alq, layer and electrodes were simultaneously

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FIG. 8. Forward-biased I-V characteristics of three devices with structure Mg:Ag-500-A Alq,-bottom electrode, where the bottom electrodes are ITO, Mg, and Ag. Also shown is the difference in operating voltages between the ITO and Ag devices and the I T 0 and Mg devices as a function of current (dashed lines). The process of electron transport through devices is sketched on the right (from BuloviC, 1998).

deposited on all three substrates to minimize differences between top contacts and bulk transport properties from device to device. Since the current of the three devices is primarily carried by electrons, the differences in the I-V characteristics are not due to the additional hole current, but can be explained by different magnitudes of the builtin dipole at the bottom interface. That is, band bending in Alq, can form an energy barrier W, at the bottom interface as shown in Fig. 8. To first approximation, W, should add a constant voltage offset to the I-V characteristics of the device. Indeed, the I-V curves in Fig. 8 are shifted by 0.8 If: 0.1 V for the device with the Ag electrode, and 0.5 f 0.1 V for the device with the Mg electrode relative to the I T 0 anode. The change in the magnitude of the offset may originate from the difference in morphology of Alq, layers on different substrates. Finally, in Fig. 9 we investigate current spreading by comparing the devices forward-bias current density of the three Mg:Ag-Alq,-Mg:Ag with a 450-A-thick Alq, layer and 0.90-, 0.25, and 0.07-mm2top electrodes. The identical magnitudes of the normalized characteristics indicates that current scales with electrode area, suggesting that enhanced

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1 10'

5 c

104 104

g

10-4

F f

'Od 104

g

10-7

0

101 L

1

Voltage M

10

FIG. 9. Forward-biased I-V characteristics of three Mg:Ag-450-A Alq,-Mg:Ag devices with 0.90,0.25, and 0.07 mm2 top electrode areas, normalized with respect to device area (from Bulovik et a[., 1997).

electric fields at the contact periphery do not significantly affect the current.

b. Carrier Transport through an Amorphous Molecular Organic Material In transport-limited single-organic-layer molecular OLEDs, three different conduction mechanisms are commonly observed: ohmic, space-chargelimited (SCL) conduction, and trapped-carrier-limited (TCL) space-charge conduction. We describe each regime of current conduction in an ETL, noting that similar equations describe conduction in an HTL. Ohmic conduction is observed at low voltages when the density of injected carriers ninjis smaller than the thermally generated background free charge density no. In this regime, the current density is given by Ohm's law: J = 4pnno Vld

(2)

where 4 is the electronic charge, pn is the electron mobility, V is the applied voltage, and d is the ETL thickness. Space-charge-limited conduction is observed when ninj> no under conditions when charge trapping is not significant. The current density is then given by Child's law: J = (9/8)4p,,&V2/d3

(3)

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where E is the dielectric permittivity (Lampert and Mark, 1970). Given the V 2 / d 3dependence of SCL current we expect ohmic conduction to dominate at low voltages and for thicker films. As discussed, the mobility in amorphous organic thin films is also dependent on electric field, increasing the voltage dependence beyond the V z suggested by Eq. (3). Both ohmic and space charge conduction are observed in low-voltage operation of devices in Fig. 8. At very low injection currents (< 1 pA/cm2), Mg:Ag-Alq,-Ag devices follow Ohm’s law. Using Eq. (2) this sets an upper limit of nap, d 107(Vcm s)- Then, using p, lo-* cm2/V s (see Fig. 4), we infer no Q 10l2cm-,, consistent with the insulating nature of Alq, at low voltages. The larger W, in Mg:Ag-Alq,/lTO devices (Fig. 8) results in sufficient accumulation of electrons in Alq, so that ninj= no, and SCL conduction is observed for V < 1 V. At higher voltages, the quasi-Fermi level F , rises toward the LUMO, filling traps that might exist in the energy gap of the organic layer. For a high density of traps, their concentration and energy distribution governs the current, resulting in trapped-charge limited (TCL) space-charge conduction. As the traps below F , fill, they reduce the density of empty traps and increase the electron effective mobility. An analytical expression relating current to voltage in the TCL regime has been derived (Lampert and Mark, 1970) assuming a continuous exponential energy distribution of traps N , ( E ) below the LUMO, as illustrated in Fig. 10. For

’.

-

where EL,,, is the LUMO band energy; N , is the total trap density; k is Boltzmann’s constant; and 7; = E , / k , where E, is the characteristic trap energy. The current density is then given by

Here, NI.UMO is the density of states in the LUMO band, and m = K/T and once again p = p(E). From Eq. (5) we see that trap-limited conduction results in a power-law dependence of current on voltage. This is observed in transport-limited Alq, devices, as shown in Fig. 11 for a series of four OLEDs with structure Mg:Ag-Alq,-ITO. Here, the thicknesses of the Alq, layers are 200, 400, 600, and 8OOA. All four OLEDs have m = 8, implying E, = 8kT 0.2eV,

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Chapter 5

10

1

VDnPeeM FIG.16. Current as the function of applied voltage in an Ag-Mg:Ag-550A A1q3-270ATPD-IT0 OLED at various temperatures (open symbols). The solid lines show a fit to the TCL conduction model using N,=3.1 x loL8~ m and - ~m,N,,,,=4.8 x lOI4(cmV s)-'. Inset: Temperature dependence of the power-law parameter m (from Shen et al., 1996).

FIG.25. (Bottom) Four illuminated elements of an alphanumericTOLED display in front of a printed text background. (Top) Same array with all elements switched off to demonstrate high transparency of the pixels. Faint outlines of patterned IT0 layers are visible (from Universal Display Corporation, Bala Cynwid, PA).

FIG.34. Micrograph of a patterned SOLED. Both organic and electrode layers were vacuum deposited through a laterally translatable triangluar shadow mask. Faint outlines of patterned I T 0 leads V, ,V,, and V, and the bottom IT0 are visible (from Tian, unpublished).

\

Glass substrate FIG.36. Schematic cross-section of a metal free SOLED. To the left are panels showing the device tuned to emit from the top surface. Inset: The biasing scheme for the metal free SOLED containing an insulating layer indicated by the open box (from Parthasarathy, et al., 1999).

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E

Trap Distribution

WE) FIG. 10. Density of trap states D(E) as a function of energy. An exponential trap distribution lies in the carrier energy gap below the LUMO states. The HOMO and LUMO states are shown as narrow energy bands of width -kT

and the injection at contacts or variation of p with voltage do not significantly influence the current conduction process. This value of the characteristic trap depth is consistent with temperature-dependent study of Alq, heterojunction OLEDs (Burrows et al., 1996b; Shen et al., 1996) (discussed in Subsection b of 4 in Section 11).

FIG. 11. Forward-biased I-V characteristics of four Mg:Ag Alq,-IT0 devices with 0.90mm2 top electrode areas and Alq, layer thickness of 200, 400, 600, and 800A (from BuloviC, 1998).

272

v. BULOVI~: ET AL.

FIG. 12. Dependence of the voltage drop over Mg:Ag Alq,-IT0 devices with 0.90-mm' top electrode areas on the square of Alq, thickness, for five constant currents (from Bulovik, 1998).

For such a high value of m, V is approximately proportional to d2 at a constant current. In Fig. 12, we plot I/ vs dZ for devices in Fig. 11, over 5 decades of current. Straight line fits to the data intercept the voltage axis above 0 V. This "zero thickness voltage" can be interpreted as the voltage dropped over the electrode-organic interfaces. Therefore, for a 600-A-thick A, approxiAlq, device with 0.90-mm2top electrode area operating at mately 30% of the applied voltage is dropped over the contacts. Our approximation of the voltage distribution in single-organic-layer OLEDs does not take into account the local distribution of electron density and electric field within the device. Furthermore, the change of electron mobility with electric field is neglected since the voltage drop over the device is proportional to p!,l'("'+')I. Our analysis suggests that the I-V characteristics of transport-limited Alq, devices are consistent with TCL conduction through the organic film for V > 2 V. Since m is sensitive to changes in temperature, measuring the temperature dependence of current transport gives further insights into the transport properties of Alq,, as described in Subsection 4 of Section 11. c. Luminescence

At sufficiently low current densities, the number of excitons generated in an OLED is proportional to the number of injected carriers. The lumines-

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0

Current [mA]

200 400 800 800 1000

273

0.00

Alq, Thickness [A]

FIG. 13. (a) Linear dependence of EL intensity on the injected current for a Mg:Ag-900-A Alq,-IT0 OLED. (b) Dependence of EL equantum efficiency on the Alq, layer thickness of Mg:Ag Alq,-IT0 OLEDs (from BuloviE, 1998).

cence output then also scales linearly with current, as shown in Fig. 13a for an ITO-900-A Alq,-Mg:Ag OLED. The luminescence quantum efficiency q in these electron-current-dominated devices is low, since most electrons pass through the OLED without forming excitons. In thicker devices, exciton formation occurs farther from the metal electrode, which can act as a quenching site leading to nonradiative energy transfer (Bulovik et al., 1998a). Reduction in nonradiative processes is reflected in the corresponding rise in q with organic layer thickness (Fig. 12b). Numerical analysis of current injection in single layer OLEDs has shown that the highest q is achieved for devices with a balanced injection of both electrons and holes (Crone et al., 1998; Malliaras and Scott, 1998). Since most organic solids preferentially transport a single carrier type, balanced injection can be achieved by increasing the injection barrier for the higher mobility carrier. Although this leads to a concomitant increase in the operating voltage, it is predicted that for some OLEDs the power efficiency may be improved by this method (Crone et al., 1998).

4. CHARACTERISTICS OF HETEROJUNCTION OLEDs a. Luminescence The low efficiency of single-layer OLEDs is in part due to the low probability for exciton formation in the organic film. In multilayer hetero-

v. BULOVI~: ET AL.

274

junction OLEDs, the probability of charge carrier encounter is increased by designing structures in which at least one of the injected carriers is confined at an organic heterointerface. As an example, in Fig. 14 we illustrate the process leading to the formation of excitons in Alq,-based heterojunction devices using 4,4‘-bis[N-( 1napthy1)-N-phenyl-amino] biphenyl (a-NPD) for the HTL. Here, electrons are injected from the cathode into Alq, while holes are injected into a-NPD from the anode. Under the influence of an applied electric field, the carriers hop from molecule to molecule to eventually arrive at the HTL-ETL interface. Energy level lineup at this interface forms a barrier to the injection of electrons from Alq, into a-NPD, and a smaller barrier to hole injection from a-NPD into Alq, (Rajagopal et al., 1998). Penetration of holes into Alq, is limited by their very low mobility in this material (Lin et al., 1997). In the vicinity of the ETL-HTL interface, a hole can encounter an electron in a “self-trapped state” or polaron of the Alq, molecule thereby forming a Frenkel exciton. After a time approximately equal to the exciton lifetime (1- 10 ns), the electron-hole pair either radiatively recombines generating light, or nonradiatively recombines to generate heat. The Alq, molecule is then free to accept the next injected carrier. We note that in radiative processes the energy of the emitted photon is smaller than the HOMOLUMO gap due to both polaron relaxation (“self-trapping”) and Coulomb binding of the electron-hole pair. In the EL process, the injected electron and hole are presumed to be

ELUMO

FIG. 14. Schematic energy level diagram showing the process of charge transport and subsequent relaxation of trapped electrons in Alq3 with a minority hole, to form a distribution of Frenkel excitons in the Alq, bandgap. These self-trapped excitons subsequently recombine to generate electroluminescence. Here F , and F , represent the quasi-Fermi levels for electrons and holes, respectively (from Burrows et ul., 1996b).

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uncorrelated and can form both singlet and triplet excitons. It has thus been assumed that the triplet-to-singlet ratio in OLEDs is 3 to 1 (Brown et al., 1992), corresponding to their multiplicity. Studies on Alq,-based OLEDs indicate that only -40% of the injected charges combine to form excitons that can contribute to a luminescent process, and of these only 22 & 3% are singlet states (Baldo et al., 1999b; Bulovic et al., 1998a). The remaining 60% of the injected carriers either do not form excitons or form excitons that are rapidly quenched at defects within the device, transfer their energy to surface plasmon modes in the metal electrode, or to lossy waveguide modes (Bulovic et ul., 1998a). The EL quantum efficiency of heterojunction OLEDs can be more than an order of magnitude larger than that of single layer devices. This is evident in Fig. 15, where we plot the external quantum efficiency q for a series of Ag-Mg:Ag-Alq,-tx-NPD-CuPc-ITO-glassOLEDs. Here, q 0.6% compared with 0.02% for single layer OLEDs (see Fig. 13). The change in q with Alq, layer thickness can be explained in term of microcavity effects (Bulovik et al., 1998a). Exciton quenching at interface defects may also be responsible for the lower efficiencies of the thinnest devices. The EL efficiency also increases at lower temperatures as the number of phonon modes (and hence routes to nonradiative internal conversion) in the luminescent film decreases. For example, for Alq, -TPD OLEDs with 550-A-thick Alq, layer, q increases from 0.6% at 375 K to -3% at 100 K. Similarly, the PL efficiency of 1000-A-thick Alq, film changed from 32% at 300 K to 47% at 120 K (Burrows et al., 1996b; Shen et al., 1996).

-

0

500

1000

Alq, Thickness [A] FIG.15. Dependence of EL quantum efficiency (measured in the forward-scattered direction) on the Alq, layer thickness of 500-8, Ag-1000-8, Mg:Ag Alq,-350-8, a-NPD-50-A CuPc-IT0 OLEDs (from BuloviC et al., 1998a).

276 b.

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Trapped-Charge-Limited Space-Charge Conduction in Alq,

In part b of Subsection 3 in Section I1 we showed that the TCL model fits the I-V characteristics of single organic layer Alq, OLEDs. We now apply the same model to the Alq,-diamine heterojunction OLEDs and measure their I-V behavior a function of temperature. The I-V characteristics of a Ag-Mg:Ag-550-A Alq3-270-A TPDITO-glass OLED measured over the range 120 K < T < 300 K are shown as symbols in Fig. 16 (see color insert), together with the fit to the TCL model (solid lines). At all temperatures, Eq. ( 5 ) fits the data over 6 decades of current with m varying from 6 2 1 at 300 K to 15 f 1 at 120 K. The inset to Fig. 16 is a plot of m vs 1/T, which is fit by a straight line corresponding to ?; = 1780 f 200 K or E, = 0.15 -t 0.02 eV. Since E, = 6kT, most of the traps are buried deep within the HOMO-LUMO gap as required by the TCL model. Using the measured value of IT; to determine the energy dependent density of traps in Eq. (9,we then find N,=3.1 x lo1*cm-, and p,,NLUMO =4.8 x lot4 (cm V s)-'. Then, using p,, = 5 x 10- cm2/v s we obtain NLUMo = l O I 9 ern-,. The fact that NLUMo and N, are roughly equal (and within two orders of magnitude of the density of molecular sites) suggests that each conduction electron is self-trapped on an Alq, molecule due to the resulting formation of a polaron. Therefore, each molecular site serves as both the available electron state and a trapping site. This is consistent with molecular orbital calculations that indicate that the Alq, anion LUMO is -0.2eV lower in energy than the neutral state (Burrows et al., 1996b), and with the thermally stimulated luminescence (Forsythe et al., 1998) and photoemission (Probst and Haight, 1997) measurements that identify a distribution of trapped states from 0.25 to 0.13 eV. The exponential distribution of trap energies, as suggested by the preceding analysis, is possibly due to the presence of a large number of molecular conformations resulting from local inhomogeneities in molecular density in the amorphous thin film. Alternatively, it has been proposed that the exponential trap distribution originates from the tail of broad Gaussian-shaped LUMO and HOMO levels (Campbell et al., 1997). Here, the states in the centers of Gaussian distributions act as transport sites between which the carrier hops, while the tails act as traps. Finally we note that extrinsic trap states may also be responsible for N,, although there is no independent evidence for such a high density of impurities in the Alq,. To investigate the fraction of the applied voltage that is dropped over the Alq, layer, a series of Ag-Mg:Ag-Alq,- a-NPD-CuPc-IT0 OLEDs with different Alq, layer thicknesses were studied (Fig. 17) (Bulovik, 1998). The power-law dependence changes from m = 7 for the OLED with a 950-A-

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FIG. 17. Dependence of the I-V characteristics on the Alq, layer thickness of 500-A Ag-1000-A Mg:Ag A1q3-350-A a-NPD-50-a CuPc-IT0 OLEDs with 0.90-mm2 top electrode areas (from BuloviC, 1998).

thick Alq, layer, to m = 11 for the OLED with a 200-A-thick Alq, layer. This is in contrast to single-layer Alq, OLEDs, for which m = 8 for all thicknesses. The increase in m for OLEDs with a thinner Alq, layer is understood when considering that fraction of voltage dropped over the cr-NPD-CuPc-IT0 layers and contacts is larger for OLEDs with thinner Alq, layers. Change in this voltage offset results in a larger value of m. In Fig. 18 the voltage drops over both single layer Alq, devices and Alq,-cr-NPD OLEDs at I = A (device area 0.90 mm’) are shown. The operating voltage of the 650-A-thick Alq, single-layer OLED is -0.8 times that of the heterojunction device. In part b of Subsection 3 in Section I1 we estimated that -70% of the voltage is dropped over the Alq, layer in the single layer device, with the remaining voltage dropped across the contactorganic interfaces. From this, we determine that most of the voltage applied to the heterojunction OLED is dropped over the 650-A-thick Alq, layer. As long as the fraction of voltage dropped over the Alq, layer is independent

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12.

0 0

I

I

I

I

I

I

I

I

200

400

600

800

I

1

1000

Alq, Thickness [A] FIG. 18. Dependence of the voltage drop over devices sketched in the figure and with 0.90-mm2 top electrode areas, at a constant current of I = 10-4A (from BuloviC, 1998).

of current, the TCL conduction model, as described, provides a reasonable approximation to OLED performance.

c.

Other Conduction Models

In this chapter we have discussed current conduction and luminescence processes in OLEDs, focusing on trapped-charge-limited space-charge conduction. However, other conduction models have also been proposed including both injection and bulk transport dominated mechanisms. For example, trap-free space-charge-limited conduction has been used to explain the field and temperature dependence of the I-V characteristics (Blom et al., 1996, 1997). Injection dominated mechanisms such as tunneling (Marks et al., 1993; Parker, 1994) and thermionic emission (Gmeiner et al., 1993) also agree with the measurements, although in all cases these other models can fit the data only over a limited range of operating currents and temperatures. Many of these models may be valid in some regime of operation, for some device structures or materials systems. The TCL model, however, exhibits

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excellent agreement with the I-V-T and thickness-dependent characteristics of the widest range of materials, device structures, and operating conditions commonly found in practical molecular organic LED applications, and has also been successfully applied to polymer LEDs (Campbell et al., 1997). 111. Improved Operation of Organic Light-Emitting Devices

The “conventional” Alq,-diamine OLED structure described in Subsection 4 of Section I1 has served as the standard for research on electroluminescence of organic materials. Soon after the first demonstration, a number of improvements to the conventional OLED structure were suggested, including doping to increase the luminescence output (Tang et al., 1989), packaging in an inert atmosphere (Burrows et al., 1994), and growth on shaped (Gu et al., 1997b) or flexible (Gu et al., 1997a; Gustaffson et al., 1993; May, 1996) substrates. Some of these improvements are reviewed in the following. BY DOPING 1. MODIFYING EMISSION COLORAND EFFICIENCY

One method for improving the luminescence efficiency of OLEDs is by doping the active layers with high-quantum-efficiency luminescent dyes (Tang et al., 1989). Due to Forster or Dexter energy transfer (Pope and Swenberg, 1982) the exciton in a host material can transfer its energy to a luminescent guest molecule. Similarly, the energy level structure of the guest molecule might render it an electron or a hole trap for carriers traveling through the host, which can again lead to exciton formation directly on the guest molecule. In general, during energy transfer an excited donor molecule D* transfers its energy to a luminescent acceptor in the ground state A, which in turn is promoted to an excited state:

Here hv, and hvT are, respectively, the optical and thermal energies emitted in the process. That is, in some doped organic thin films, the exciton energy can be transferred from the host to the guest molecules, quenching the luminescence of the host while increasing that of the guest. The interaction between D* and A is described by the perturbation Hamiltonian, H . In this case, the D* A is not a stationary state of the

+

280

V. B U L O VET I ~ AL.

total Hamiltonian, but can evolve to other isoenergetic states, such as D + A*. Time-dependent perturbation theory then assigns the probability of evolution from D* A , described by Y i , to D + A*, described by Y f , as

+

where p is the density of coupled isoenergetic donor-acceptor transitions. For organic molecules with broad spectral features, p can be estimated by calculating the overlap integral of donor luminescence and acceptor absorption. In general, after the exciton transfers from D* to A it promptly relaxes to the lowest vibrational level of A*. The pertubation H contains several terms, of which the most important are the electrostatic (coulombic or Forster) and electron exchange (or Dexter) interaction, both independently capable of inducing energy transfer. The electrostatic interaction can be expressed as a series of multipolemultipole terms. The dominant, dipole-dipole interaction has been treated by Forster who found that the rate of energy transfer depends on the distance R between the donor and the acceptor molecules (Pope and Swenberg, 1982; Forster, 1965):

Here, R , is the Forster radius, and z is the average donor exciton lifetime for recombination in the absence of energy transfer corresponding to rate K , = l/z. When R = R, then K,, = K,, and the probability that an exciton will recombine at the donor is equal to its transfer probability. The critical distance R , is given by the integral (Pope and Swenberg, 1982; Forster, 1965) over all wavelengths A

"s

RX = 1.25 x l O " 2 n4c3

FD(A)gA(A)Az dA

(9)

where 4E is the quantum efficiency of donor emission, n is the refractive index of the host, F , is the normalized emission spectrum of the donor, and aA is the molar extinction coefficient of the acceptor. Using Eq. (9) it has been shown (Pope and Swenberg, 1982) that dipole-dipole interaction in an organic guest:host thin film can be significant even at distances as large as 100A.

5 MOLECULAR ORGANIC LIGHT-EMITTING DEVICES

28 1

In Forster energy transfer, the spin of both D and A is conserved. Therefore. the allowed transitions are

where superscripts 1 and 3 indicate singlet and triplet states, respectively. The triplet-singlet transition

is forbidden, but is sometimes observed since ,0* has a long lifetime and, although slow, K,, can be faster than ,D* -+ 'D (Pope and Swenberg, 1982). Energy transfer due to electron exchange has been analyzed by Dexter to give (Pope and Swenberg, 1982; Dexter, 1953)

where R is distance between D* and A, and L is a constant. Because this process involves exchange of electrons, it occurs only over a short distance of N 10A. Under the Dexter transfer process, the total system spin is conserved (Pope and Swenberg, 1982). Triplet-triplet energy transfer, which was doubly forbidden in the Forster process, is now allowed: ,D*

+ 'A

-+

'D

+ ,A*

(12)

A pictorial representation of both Forster and Dexter processes is shown in Fig. 19 for singlet-singlet energy transfer. Even a small concentration of guest molecules (< YO) in the electrically pumped host layer is sufficient for complete Forster energy transfer to the guest, quenching the host luminescence. A low dopant concentration results in a large average distance between guest molecules, which reduces excitonexciton quenching and yields highly efficient luminescence. The lifetime and efficiency of these "guest-host" OLEDs have both been found to depend on the nature and concentration of dopant molecules in the light emitting layer (Tang, 1996). Some examples of doped Alq, OLEDs are shown in Fig. 20. For comparison, the luminescence spectra of undoped Alq, is also shown. The Forster energy transfer in Alq,:(2%) DCM2 OLEDs leads to orange EL with an external quantum efficiency of vex 0.5% (Tang et al., 1989; Bulovik N

282

V. BULOVICET AL. FBrster, Coulombic (long range -30-1 00 A)

Dexter,e- exchange (short range -6-20 A)

FIG. 19. The Forster and Dexter processes in singlet-singlet energy transfer

et al., 1998b). Dexter energy transfer in Alq,:(6%) PtOEP OLEDs leads to saturated red phosphorescence with vex 4% at low intensities (Baldo et al., N

-

1998). Also, coumarin 6 doped in Alq, improves the color saturation and efficiency (vex 2%) of green OLEDs (Tang et al., 1989), presumably due to a combination of carrier trapping and Forster transfer.

1.0

2

'm c

0.8

'

.

s!

C 1

0.6

w

F.Cr

0.4

!i p

0.2

0.0 I ( 400

500

,

1

600

700

I 800

Wavelength [nm] FIG. 20. The normalized electroluminescence spectra of heterojunction OLEDs with an undoped Alq, active layer (Alq,), DCMZ-doped Alq, (DCMZ:Alq,), PtOEP-doped Alq, (PtOEP:Alq,), and coumarin 6-doped Alq, (C6Alq3) (from Bulivit, 1998).

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In doped OLEDs, host molecules transport electrons or holes and efficiently transfer energy to the luminescent center (the guest). Separation in functionality between host and guest molecules allows for greater flexibility in OLED design.

PROCESSES 2. DEGRADATION Atmospheric oxygen and water vapor often form chemical complexes with organic compounds. For example, exposure of Alq, to moisture leads to formation of nonemissive species that can act as luminescence quenching sites (Papadimitrakopolus et al., 1996; Papadimitrakopolus and Zhang, 1997). Indeed, it was observed that degradation of aluminum tris-(g-hydroxyquinoline) (Alq,) based OLEDs occurs in air on a time scale of a few hours (see inset of Fig. 21a), while even simple encapsulation in an inert atmosphere, using an epoxy sealed glass package (Burrows et al., 1994) can extend the device half-life to > 1000 h at video brightness (Fig. 21a). Here the half-life is defined as the time for the brightness to decline to 50% of its initial value when the OLED is operated at a constant current. The primary mode of luminescence degradation involves formation of nonemissive “dark spot defects” that also increase the device resistance (hence the increase in operating voltage in Fig. 21b). The dark spot growth is accelerated at higher electric fields. Even when packaged, however, OLEDs are susceptible to degradation at elevated temperatures. Thermally induced failure of these structures is found to be due to morphological changes in organic layers with a low glass transition temperature (T,) (Fenter et al., 1997). Particularly, it was shown that by increasing the T, of hole transporting layers the lifetime of OLEDs under pulsed drive current operation can be extended to 4000h at a brightness of 500 cd/m2 (VanSlyke et al., 1996). These results also show that the rate of degradation is proportional to the drive current (and hence the luminescence intensity) of the OLED, which implies that > 20,000 h of operation of these OLEDs at video brightness (of 100 cd/m2) can be realized. Doping the luminescent layer of OLEDs with emissive centers can further extend operating lifetime as well as enhance the luminescence efficiency and color saturation (see Subsection 1 of Section 111). Figure 22 compares the stability of three green Alq,-based devices (Tang, 1996) with both doped and undoped luminescent layers, all driven at 20 mA/cm2. The half-life of the undoped Alq, device is 4000 h, with an initial luminance of 510 cdjm2. The initial luminescence of the quinacridone (QA)-doped device is 1600 cd/ m2, but with a half-life of only 300-400 h. Doping methyl quinacridone

-

284

V. BULOVICET AL.

2.4

-

2.0

-

0

0

0.0

'

0

.

'

200

-

'

*

400

'

-

600

,

'

800 1000

Time [hours] 24

20

16 Q)

-3 1 2 U

3 8 4

0

Time [hours] FIG.21. (a) Time evolution of the average optical power output of a population of approximately 10 encapsulated OLEDs. The inset shows the corresponding evolution of identical, unencapsulated devices. (b) Time evolution of the average voltage required to maintain the current density of 10mA/cm2 (from Burrows et al., 1994).

(MQA) molecules into Alq,, however, results in OLEDs with both high initial luminescence efficiency (1400 cd/m2) and a long half-life (6000 h), probably due in part to removal of dissipative triplet excitation and reduction of dopant-dopant interactions. Again assuming that the rate of degradation is proportional to the drive current, this translates to a half-life in excess of 50,000 h for an initial luminance of 100cd/m2.

MOLECULAR ORGANIC LIGHT-EMITTING DEVICES

5

285

I.o

3 0.9 2.

.P c 0.6

CL 0.5

I

5 10

50 100

1000

loo00

Hours of operation FIG. 22. Stability of Alq, devices driven at 20mA/cm2. The initial luminance for Alq, devices is 510 cd/m2, for QA-doped Alq, devices is 1600 cd/m2, and for MQA-doped Alq, devices is 1400cd/m2 (from Tang, 1996, courtesy of Society for Information Display).

Recrystallization of luminescent species can also contribute to the decrease in luminescence efficiency (Papadimitrakopolus et al., 1998). Crystal formation increases the number of phonon modes that quench luminescence and concomitantly red shifts the emission spectrum. One approach to solving this problem is to use molecules that are sterically prevented from crystallization (Adachi et al., 1995; Shirota et al., 1994). Alternatively, it has been shown that crystallization can be suppressed by doping (Sano, 1996).

EFFICIENCY USINGSHAPED SUBSTRATES 3. INCREASING Various dielectric and conducting layers that comprise an OLED typically have different refractive indices. As a result, the emitted photon undergoes numerous reflections and refractions before escaping from the substrate. A fraction of this light is emitted in the forward direction, while the remainder that exceeds the glass-air critical angle (0, 19O) (Gu et al., 1997b; Garbuzov et al., 1996b) can be waveguided in the organic layers or can radiate through the substrate edge (Fig. 23a). By redirecting the edge-emitted light, forward emission can be enhanced by 70 to 90% (Bulovic et al., 1998a), as has been demonstrated (Gu et al., 1997b) using shaped substrates shown in Fig. 23b. With these mesa structures, light emitted at large angles from the normal can be reflected into angles < 0,and therefore be efficiently coupled out from the substrate. This can lead to a factor of 2

-

286

V. BULOVICET AL. Glass Substrate

,OLED

(a)

,OLED

(b)

fomardacattered light

FIG. 23. (a) A fraction of light generated by a conventional OLED grown on a flat substrate is emitted through the substrate edges. (b) By growing OLEDs on shaped substrates, much of the edge-emitted light can be redirected into the viewing direction (from Gu et al., 1997b, courtesy of IEEE).

enhancement in external quantum efficiency as compared with similar OLEDs grown on flat glass substrates.

4. FLEXIBLE OLEDs The total layer thickness in an OLED is typically <0.5pm. Such thin layers are expected to exhibit considerable flexibility when bent over macroscopic distances. In contrast to inorganic materials, which are often constrained in the choice of the substrate due to the requirement for epitaxial growth, amorphous organic materials form continuous, smooth layers on any sufficiently flat substrate. Also, during film deposition, substrates are typically maintained at room temperature, enabling the use of low melting point substrates. And finally, many low-molecular-weight, luminescent organic materials are bonded by flexible van der Waals forces (Forrest, 1997). Taken together, these three properties identify the lowmolecular-weight organics as a unique class of materials useful for forming flexible, conformable, light-emissive elements. Flexible polymeric OLEDs have been demonstrated on polyaniline substrates (Gustaffson et ul., 1993) and ITO-coated plastic substrates (May, 1996). Similarly, large area OLEDs (- 1 cm2) employing low-molecularweight organic materials have also been demonstrated on 175-pm-thick polyester sheets precoated with I T 0 (Gu et al., 1997a). These latter devices were sufficiently flexible to permit repeated bending over a radius of > 0.7 cm without mechanical damage or changes in the current-voltage response and luminescent efficiency.

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5. TRANSPARENT OLEDs

Many organic materials exhibit a large Franck-Condon effect where the luminescence band is significantly red-shifted from the absorption band (Pope and Swenberg, 1982). This renders them transparent to their own luminescence and often throughout most of the visible spectrum. Using the Franck-Condon effect, along with a compound transparent cathode, a novel class of OLEDs that are greater than 70% transparent when turned off and that emit light from both top and bottom surfaces with up to 0.4% external quantum efficiency when turned on, have been demonstrated (Bulovii: et al., 1996; G u et al., 1996). Such transparency offers the potential for very high-definition full-color displays in which the red (R), green (G), and blue (B) emission layers are placed in a vertically stacked array providing for a simplified fabrication process, as well as a minimum R-G-B pixel size and maximum fill factor. The transparent device can also be used for helmet-mounted windscreen-mounted or other “head-up’’ display applications. The transparent OLED (TOLED) shown in Fig. 24a is grown on a glass substrate precoated with an I T 0 anode. The device consists of a 200-&thick layer of the hole conducting compound N,N’-dipenyl-N,N’-bis(3-methylphenyl)-l,I’-biphenyl-4,4‘-diamine(TPD), followed by a 400-A-thick layer of Alq,. Typically, the top, electron injecting contact of a conventional OLED consists of a thick opaque electrode composed of a low-workfunction alloy such as Mg:Ag. In contrast, the top “compound electrode of the transparent devices consists of Mg:Ag that is thinner than the optical skin depth (-200A), followed by an overlayer of rf sputtered ITO. This results in a transparent top electrode with injection characteristics similar to that of the conventional OLED top electrode.

FIG.24. Schematic illustration of the transparent OLED structure (from BuloviC et al., 0 1996, with permission of Nature).

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The feasibility of manufacturing more complex display devices is demonstrated in Fig. 25 (see color insert) where TOLED-based alphanumeric display in on (Fig. 25, bottom) and off (Fig. 25, top) state is shown (Universal Display Corp.). In this structure I T 0 has been lithographically patterned in the shape of the emissive elements prior to the blanket deposition of the organic layers and the cathode. The individual elements appear to be very transparent even though these images are taken under the condition that light illuminating the background must pass twice through the TOLED. To study the contribution of different layers to the absorption of the completed device, the transmission spectra of the ITO-glass substrate (filled circles), the Alq,-TPD-ITO-glass structure (open circles), and the full TOLED (open squares) are compared in Fig. 26. The device becomes absorptive at short wavelengths (< 400 nm) owing to a combination of Mg:Ag absorption and the strong molecular transitions of Alq, (Garbuzov et al., 1996a), and at long wavelengths due to absorption by the Mg:Ag. However, the device is 63% transparent at the peak (A = 530nm) emission wavelength of Alq,, and this transparency extends across the visible spectrum.

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400

500

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Wavelength [nm] FIG.26. Transmission spectra of various TOLED layers. The thicknesses of top ITO, Mg:Ag, Alq,, TPD, and bottom I T 0 are 400, 100, 400, 200, nd 1600A, respectively. Inset: Transmission of the ITO-MgAg cathode contact at a wavelength of 1, = 530nm as a function of Mg:Ag layer thickness (from Rulovic et a)., $3 1996, with permission of Nature).

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The inset of Fig. 26 shows the transmission of the Mg:Ag contact at a wavelength of /z = 530 nm, with Mg:Ag film thickness ranging from 50 f 10 to 400 f 10A. For a 75-A-thick Mg:Ag layer, a contact transparency of 81% and a total device transparency of 71% is measured. An 81% transmissive film of Mg should be 15 A thick (the real and imaginary coefficients of the refractive index of Mg at II = 530nm are n = 0.57 and k = 3.47, respectively; O’Bryan, 1936). This suggests that at least 60 I lOA of the Mg:Ag cathode is compositionally changed in the deposition process, presumably due to a redox reaction with I T 0 during sputtering, or diffusion of Mg into Alq, (Rajagopal and Kahn, 1998). Figure 27 shows the EL spectra emitted from the top and substrate surfaces of a TOLED. Typical operating conditions for a 1-mm-diameter device are 100pA drive current at lOV, leading to a brightness of -500cd/m2 from one surface. The emission spectra are similar to those obtained for conventional Alq,-based OLEDs, although there is a 10-nm red shift of the peak of the emission spectrum from the top (Mg:Ag-ITO) contact relative to that from the substrate due to filtering by the Mg:Ag film and microcavity effects (Bulovic et al., 1998a). Approximately 10% less power is emitted from the top contact than from the substrate of the TOLED. The I-V characteristics of TOLEDs are comparable to those of a conventional OLEDs, as shown in Fig. 28. Here, the transparent cathode is

1.o

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700

800

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400

500

600

Wavelength [nm] FIG. 27. The electroluminescence spectra from the top and bottom surfaces of a TOLED with 100-A-thick Mg:Ag layer (from Gu et al., 1996).

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FIG. 28. Current-voltage characteristics of conventional and transparent OLEDs made under similar conditions (from Gu et ul., 1996). Inset: Optical output power from both substrate surfaces vs. drive current of a metal-free TOLED with a 1 % coumarin 6 in Alq, luminescent layer (from Parthasarathy et ul., 1998).

formed by a 100-A-thick Mg:Ag layer with a 400-a-thick I T 0 overlayer, while the conventional cathode consists of 1500-A-thick Mg:Ag layer and a 500-A-thick Ag overlayer. The TOLED has a larger leakage current due to the conducting paths resulting from the damage in the organic layers incurred during I T 0 sputtering. Also the operating voltage is somewhat higher for TOLEDs due to the additional resistance in the Mg:Ag-IT0 interface, resulting from the barrier formed by the complexation of Mg:Ag with ITO, and surface defects formed during the sputtering process. The transparency of the TOLED was further improved using a metal-free cathode composed of a CuPc ( - 50 A)-ITO bilayer, resulting in 85% transmission in the visible (Parthasarathy et al., 1998). Here, it is presumed that defect states in the CuPc layer generated during I T 0 sputter deposition contribute to efficient electron injection. The total external quantum efficiency of coumarin 6-doped metal-free TOLEDs (MF-TOLEDs) is similar to that of doped TOLEDs with vext;= 0.38 f 0.05% (inset in Fig. 28)

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(Parthasarathy et al., 1998). The luminance of both devices is 2000 cd/m2 at a 100mA/cm2 drive current. The metal-free contacts might find uses in devices such as electrically pumped organic lasers, whereby reducing light absorption at metal contacts can result in a reduction in the lasing threshold current. The use of TOLEDs in display applications is further discussed in Subsection 5 of Section IV, where we describe construction of independently addressable stacks of multicolor EL elements. The development of a transparent, thin film organic light-emitting device demonstrates a significant advantage of organic materials in realizing highdefinition, full-color, and head-up displays. The TOLED structure works equally well for OLEDs emitting across the visible spectrum, since the large Franck-Condon shift between emission and absorption is characteristic of many organic materials.

6 . INVERTEDOLEDs

To be useful in low-cost active matrix displays, OLEDs must be integratable with the back plane electronics. A conventional OLED is grown on a transparent anode such as ITO, and its emission is viewed through the substrate, complicating integration with electronic components such as silicon based display drivers. The TOLED structure, however, can be integrated as it emits through the transparent top contact, and it can be grown directly on a silicon substrate. One limitation of TOLEDs is that their anode forms the bottom contact, whereas for display drivers employing n-channel field-effect transistors (FETs) it is desirable that the bottom contact be the cathode. That is, for stable FET operation the source ( S ) terminal should be connected to a fixed potential, with the OLED connected to the drain (D)terminal, as shown in Fig. 29. OLEDs with cathodes as the bottom contact have an inverted biasing arrangement as compared to conventional devices and their electroluminescence is emitted through the semitransparent anode. This class of structures is called “organic inverted LEDs” (OILEDs) (Bulovii: et al., 1997). In the OILED configuration, the I T 0 anode must be sputter-deposited onto the relatively fragile hole-conducting organic thin film, which can result in unacceptable degradation of the device operating characteristics. Damage, however, can be minimized by using a crystalline organic layer (Bulovik et al., 1997), to protect the underlying HTL from the sputter deposition of the I T 0 anode. Hence, this crystalline organic layer serves as a “protective cap layer” (PCL). Prior to the demonstration of OILEDs a similar integration of molecular OLEDs with silicon was achieved using a tunneling SiO, interface (Kim et al., 1994). The tunneling interface, however, increases the operating voltage

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AL.

EL Emisslon

OILED

Organic Layers Comprising an OILED

FIG. 29. An OLED with a cathode as the bottom contact (OILED) can be directly integrated with n-FET-based drivers on display backplanes (from Bulovii: ef a/., 1997).

and can be avoided in the surface-emitting OILED structure. Also, polymerbased surface-emitting OLEDs were previously demonstrated on silicon with a transparent I T 0 and a semitransparent Au or A1 top anode (Baigent et al., 1994; Kim et al., 1994). In those inverted polymer OLEDs, the PCL was not necessary since the glass transition temperature (T,) of polymer materials is much higher than the temperatures known to damage molecular organics. Furthermore, polymers are highly conjugated, which efficiently distributes the deposition energy over many bonds. The OILED structure, shown in Fig. 30 was grown starting with the thermal evaporation in vacuum of a 1000-A-thick cathode consisting of 2 5 1 Mg:Ag alloy, followed by a 500-A-thick Alq, ETL and LL, and a 250-A-thick HTL of TPD or a-NPD. To protect the fragile HTL from the sputter deposition of the top, I T 0 anode contact, either a 3,4,9,10-perylenetetracarboxylic dianhydride (PTCDA) or a copper phthalocyanine (CuPc) film was employed. Both CuPc (VanSlyke et al., 1996) and PTCDA (Bulovii: and Forrest, 1996; Burrows and Forrest, 1994) can serve as efficient hole injection layers in conventional OLEDs. Furthermore, the use of PTCDA in a photodetector structure with an I T 0 electrode deposited on the film surface (So and Forrest, 1989) has previously demonstrated that this material can withstand sputter-deposition of I T 0 with minimal degradation to its conducting properties. The top, I T 0 layer was deposited by a low-power rf magnetron sputtering process (Bulovii: et al., 1996). The forward-bias current-voltage characteristics of OILEDs with PTCDA and CuPc PCLs as well as of a device with no PCL are shown in Fig. 31 to be similar to those of conventional OLEDs (Burrows et al., 1996b;

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EL Light

IT0 PTCDA or CUP (40-1OOA)

PTCDA

/

'

TPD or a-NPD

A'% Mg:Ag Silicon FIG.30. Schematic illustration of the OILED structure (from Bulovic et al., 1997).

1

Voltage M

10

FIG. 31. Forward-biased current-voltage characteristics of 0.05-mm2 OILEDs with different protection layer compositions and thicknesses. All devices were driven to their maximum current (Imax)prior to failure (from Bulovic et al., 1997).

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Shen et al., 1996; Burrows and Forrest, 1994). The luminance at a current density of 10 mA/cm2 is between 40 and 100 cd/m2 for all devices, independent of details of the HTL and PCL anode structure. The operating voltage of OILEDs employing CuPc as the PCL was independent of CuPc thickness between 40 and 170k In contrast, the operating voltage of PTCDA-protected OILEDs abruptly decreased by 1.5 V as the PTCDA thickness increased from 40 to 60 A.The voltage drop across the PCL is small compared to that across the rest of the device since PTCDA and CuPc are both thinner and more conductive (Burrows and Forrest, 1994) than Alq,. Therefore, the abrupt change in PTCDA-based devices reflects a change in the carrier injection efficiency from the I T 0 contact, which upon sputter deposition inflicts film damage to the topmost organic layer. This damage resulted in a significantly reduced device yield for devices lacking a PCL. The abrupt increase in the operating voltage for OILEDs with < 40-A-thick PTCDA layers occurs when the thickness of the damaged region is comparable to the E’CL thickness. Figure 32 shows the light intensity vs current response of the OILEDs in Fig. 31. The external EL quantum efficiency of the protected OILEDs is = 0.15 f 0.01% vs q = 0.30 0.02% for the unprotected devices. This difference is due in part to absorption by the PCL. A different PCL material which is transparent to Alq, emission is thus expected to increase the OILED efficiency somewhat. The injection efficiency of the “inverted” Mg:Ag-Alq, contacts can be inferred from the I-V characteristics of Mg:Ag-450-A Alq,-Mg:Ag devices, described in Subsection 3 of Section 11. This seemingly symmetric device exhibited a strongly asymmetric response with two orders of magnitude larger current in “forward bias” (bottom electrode positive). This suggested a different injection efficiency for the top and bottom Mg:Ag-Alq, interfaces, which was attributed to differences between contact interface chemistries and their roughness. The asymmetry implies that the top electrode is a more efficient electron injector into Alq, than is the bottom electrode. Its improved injection efficiency was ascribed to the presence of interface defects formed during cathode deposition (Fig. 7). Since the OILED cathode corresponds to the bottom contact of the Mg:Ag-Alq,-Mg:Ag sandwich, the electron injection efficiency is lower in OILEDs as compared to conventional OLEDs, leading to a concomitant increase in the OILED operating voltage. The lower injection efficiency in OILEDs can also be responsible for decreased q, as previously observed in a series of Alq, devices using cathodes of various composition (Hung et al., 1997). It is clear that the present OILED structures can be improved by decreasing the potential drop over the cathode interface and increasing the electron injection efficiency. We also need to better understand the electronic properties of the cap layer, which mechanically protects the underlying

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Wavelength [nm]

0.0

0.4

0.8

1.2

1.6

Current [mA] FIG.32. Light output power-current characteristics of OILEDs with different protection layer composition and thickness. All devices driven to their maximum current (Imax) prior to failure. Inset: The EL spectrum of an OILED with a 60-wPTCDA PCL and a conventional OLED (from Bulovii: et al., 1997).

organics. Nevertheless, the demonstration of the surface-emitting organic “inverted LED teaches us about the nature of organic-electrode contacts and points to the versatility of organic materials and the ease with which they can be integrated with conventional electronic components.

IV. Color-Tunable OLEDs in Full-Color Displays Small monochrome, passive matrix OLED displays have become commercially available (Nakada and Tohma, 1996; Wakimoto er al., 1996). The goal of manufacturing a practical full-color OLED display is being pursued

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BLUE

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(b)

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ED

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,

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GREEN

BLUE

FIG. 33. Five approaches for implementing a full-color pixel with OLEDs: (a) side-by-side patterning of R, G , and B subpixels, (b) filtering of white-light-emitting OLED by color passband filters, (c) downconversion of blue light lo generate green and red light, (d) filtering of a broadband OLED by microcavity-based filters, and (e) three color tunable pixels (from Burrows et a[., 1997).

(Hosokawa et al., 1998; Miyaguchi et al., 1998). We review five commonly proposed architectures (Burrows et al., 1997) (see Fig. 33) for achieving full color. 1. SIDE-BY-SIDE PATTERNING OF RED, GREEN,AND BLUEOLEDs

Conceptually, the simplest full-color pixel consists of an R, G, and B subpixel positioned in a side-by-side (S x S) arrangement (Fig. 33a). This method was initially demonstrated in both polymers (Lidzey et al., 1998; Wu

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et al., 1996) and low-molecular-weight organics (Kijima, 1997) at very low resolution. Miyaguchi et al. (1998), however, developed a high-resolution 5.2411. molecular organic &VGA display. In the initial demonstrations, the large feature size was due to limitations of shadow mask patterning, suggesting the need for a higher-resolution photolithographic process. It has been demonstrated that postdeposition lithographic patterning of vacuum deposited crystalline organic semiconductors can achieve feature sizes of from 1 to 10pm (Cheng et a/., 1985; Yap et al., 1995). However, lithographic patterning of amorphous organic thin films used in OLEDs is more difficult as they are often soluble in solvents used in standard photolithography. Postdeposition patterning can be avoided by prepatterning walls (Tang, 1994) or using other photoresist structures (Tian et al., 1997a, 1997b) on the substrate (Fig. 34, see color insert). These structures serve as in situ microshadow masks. OLED layers are then deposited at an angle to the substrate so that areas reserved for subsequent layers and contacts are shadowed. In this way, monochromatic OLED arrays have been demonstrated (Tang, 1994; Tian et al., 1997a; Wakimoto et al., 1996), while further work on full-color arrays is in progress (Miyaguchi et al., 1998; Tang and Littman, 1994). Photolithographic patterning is feasible on small substrates (up to a few inches in diameter), but becomes more difficult when applied to large-area displays. Research on a number of simpler techniques that may eventually be adaptable to large substrates includes stamping (Kim et al., 1995; Xia et al., 1995), printing (Garner et al., 1994), screen printing (Bao et al., 1997), and ink-jet printing (Bharathan and Yang; 1998; Henber et al., 1998).

2. COLORPASSBAND FILTERING OF WHITEOLEDs White OLEDs have been demonstrated by replacing the single luminescent layer in heterojunction OLEDs with two or more sequentially deposited EL layers (Kido et al., 1995; Littman and VanSlyke, 1995) whose mixed color emission spectrum appears white. To form a full-color display the white OLEDs are grown on substrates with prepatterned color bandpass filters that each transmits only R, G, or B, as shown in Fig. 33b. No further postdeposition patterning is necessary. Unfortunately, in this approach most of the white light is removed by the color filters to achieve the required primary color. As a result, the display output power efficiency is low, requiring higher driving currents to achieve a desired brightness. Also, the rate of OLED degradation increases with device current, reducing display lifetime (Sato et al., 1998). Light absorbed by color filters is converted to heat increasing the temperature of the display, further contributing to

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accelerated OLED degradation (Sato et al., 1998). Overall, the color subtraction approach is not suitable for full-color applications requiring low power consumption.

3. WAVELENGTH DOWN-CONVERSION Many luminescent organic materials can efficiently absorb short-wavelength photons and reemit a green or a red photon with up to 100% quantum efficiency in a process called “down-conversion.” To make a full-color display (Fig. 33c), a single blue OLED can be used to pump the green and red prepatterned films of ef€icient fluorescent or phosphorescent organic down-converters (Tasch et al., 1997). A low resolution full-color display (Hosokawa et al., 1998; Nakamura et al., 1996) based on this concept was recently demonstrated. The greatest advantage of this method is the ease in patterning the single blue OLED structure over prepatterned films of spectral down-converters. The down-conversion films should be thick to absorb all of the blue light, which might raise difficulties associated with patterning and growth on nonplanar substrates. Light conversion efficiencies lower than unity will decrease the display power efficiency, which is further reduced by the inherent power loss in converting higher into lower energy photons. Additionally, light propagating through the substrate can be unintentionally absorbed and reemitted by neighboring subpixels, resulting in “color bleeding.’’ Use of shaped substrates (Gu et al., 1997b) should minimize this effect.

4. MICROCAVITY-FILTERED WHITEOL,EDs The broad EL spectrum of OLEDs can be narrowed (Dodabalapur et al., 1994) and directed (Takada et al., 1993) by placing the luminescent layer inside a microcavity. The thickness of the microcavity determines the emission color, thereby forming the R , G, and B subpixels of a full-color display (Fig. 33d). The microcavity is formed between the top, reflective metal electrode and a predeposited dielectric distributed Bragg reflector coated with a transparent I T 0 electrode. Both individual R, G, and B subpixels (Dodabalapur et al., 1996) and a red, yellow, and green pixel (Lidzey rt al., 1998) were realized using microcavity filtering. One characteristic of microcavities, however, is that in all cases the emission color is dependent on the angle of observation (Takada et a!., 1993). This “color directionality” can be reduced by incorporating a

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scattering layer outside the microcavity (Dodabalapur et al., 1995; Lidzey et al., 1998; Rothberg et al., 1994), at the expense of OLED efficiency. As with color filtering, each subpixel selects only a section of the emitted spectrum, further reducing the display efficiency. Additionally, different thicknesses of subpixel Bragg reflectors might give rise to difficulties associated with growth and patterning of the overlying OLEDs on nonplanar substrates. 5. STACKED COLOR-TUNABLE OLEDs

Displays employing coaxially emitting stacked color-tunable OLEDs (Fig. 33e), have at least a threefold increase in resolution and fill factor over S x S architectures. Each subpixel of the stacked OLED (SOLED) has to be transparent to its own radiation, and the radiation of other subpixels, as shown in Fig. 35. Therefore, the high optical transparency of TOLEDs is a critical component of stacked architecture. For example, the top I T 0 surface of a TOLED can serve as a hole injecting electrode for the second TOLED built on top of the first device. Each device in the stack is then independently addressed and can be tailored to emit its own color through the adjacent transparent organic layers, contacts and substrate. Stacked OLEDs with both two (Burrows et al., 1996a) and three (Shen et al., 1997; Parthasarathy et al., 1999) subpixels have been demonstrated. The three-color SOLED structure is illustrated in Fig. 36 (see color insert), with the color emissions of individual subpixels shown in the inset. The device in Fig. 36 is completely metal free in that the contacts consist

Trmepunt

Electrodes 2

4

RGB FIG. 35. Schematic cross section of a three-color stacked OLED.

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only of I T 0 for hole injection, and ITO/bathocuproine for electron injection (Parthasarathy et al., 1999). By eliminating all metals in the contacts, optical reflections within the stack are minimized, thereby reducing microcavity effects to an acceptable level. In addition, to simplify biasing, an insulating organic film is inserted between contacts E3 and E4, with contacts E4 and E2 shorted to ground. In this manner, each subpixel in the stack is separately referenced to ground. The green emission is broadened, after being filtered by microcavity effects, and appears pink (middle inset of Fig. 36). To minimize the influence of microcavity effects on the spectral emission, the thicknesses of device layers have to be carefully tailored. In the stacked configuration, only one layer structure is grown over the entire area of the substrate, obviating the need for postdeposition patterning or interruption of the deposition sequence for midprocess patterning. The stacked OLEDs can be grown on lithographically patterned substrates with photoresist structures that act as in situ shadow masks (Tian et al., 1997). Layers deposited at an angle to the stationary (Fig. 37a) or rotating (Fig. 37b) substrate leave different footprints that in sequence form active organic layers with separately addressable electrodes. A complete two-color SOLED is shown in Fig. 37c. To operate individual subpixels, voltage is applied between pairs of prepatterned I T 0 pads. Alternatively, patterning of both active and contact layers can be achieved by precise translation of an in situ shadow mask (Tian et al., 1999).

V. Summary After more than 10 years of research, organic light-emitting devices have evolved from a scientific curiosity to a commercially viable building block of flat panel displays. The low cost of organic materials, the simplicity of organic thin film growth, and the ease of integration with conventional semiconductor backplane technologies together with the demonstrated potential to grow OLEDs on flexible substrates are some of the benefits of this promising technology. Passive matrix, monochrome displays are commercially available, and it is only a matter of time before pixelated full-color displays reach consumer markets. This chapter reviewed the present state of OLED science and technology, emphasizing the structural properties of various devices and their implications on device performance. The new device structures demonstrated during the past several years together with conventional OLEDs are building blocks for many different organic light emitters and full-color displays. It is important to realize that although in their infancy, these

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devices offer adequate performance for a number of display applications, with the added advantage of extremely small thickness, flexibility, light weight, integrability, and relative simplicity of fabrication. It is therefore conceivable that in the not too distant future, OLEDs will become commonplace light-emitting elements in flat panel displays.

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ACKNOWLEDGMENTS We would like to acknowledge our collaborators and friends M. E. Thompson, Z. Shen, G. Gu, D. Garbuzov, G. Parthasarathy, M. Baldo, P. Tian, D. OBrien, J. Mahon, T. Zhou, and S . Mao for many hours of joint research, which led to the advances described in this chapter. We are grateful to National Science Foundation Materials Research Science and Engineering Center, the Defense Army Research Project Agency, and Universal Display Corporation for their financial support of this work.

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Index

A

degradation, 283-285 doping, 279-283 photoluminescence spectrum, 259 Alq-based devices degradation, 2 17-220 lifetime, 242-243 spectral range, 242 temperature dependence, 228-23 1 Applications LEDs, 32-33, 35-39,42-43, 132-133 organic light-emitting diodes (OLEDs), 249-250, 256-257 Arsine gas purification, 77 Atomic ordering, 59, 85 Automotive exterior signals, 37-39

Aging treatment for OLEDs, 220 AlGaAs LEDs evolution, 5 performance, 3 reliability, 119- 122 structures, 17 AlGaAs superstrates, 65 AlGaInP LEDs, 50-123 band structure, 51-53 current-voltage characteristics, 104- 107 design features, 57-72 device fabrication, 99-101 electroluminescent spectra, 111 electron mobility, 57 epitaxial growth methods, 26, 72-73, 76-92 high-power devices, 116-1 18 historical overview, 50-51 luminous efficiency, 109-1 12, 132-133 manufacturing issues, 87-88, 92-98 market growth, 45 packaging, 71-72, 101-102 performance, 3, 33, 40-41, 50-51, 102-112 photoluminescence intensity, 54-56 reliability, 17, 118-123 source materials, 76-78, 90 structures, 17, 88-92 thermal management, 112-1 16 AlGaN doping, 175-177 layer growth, 148-149 AIN buffer layers, 140 Alq, (aluminum tris(8-hydroxyquinoline)) carrier transport properties, 262-263, 269-272, 275-279

B Bandgap energy (E,) AlGaInP alloys, 51-53 GaInN, 174-175 GaN, 158-159 and materials selection, 5, 6 nitride materials, 130 and quantum efficiency, 103-104 Buffer layers in nitride-based LEDs on sapphire, 140-143 in OLEDs, 222

C

Carrier confinement, 61-63, 113-114, 179, 185 Carrier injection efficiency in AlGaInP devices, 60-61

307

308

INDEX

Carrier injection efficiency (Continued) in nitride-based LEDs, 184-185 in OLEDs, 222-223, 263-269 Carrier transport properties AlGaInP devices, 57 Alq,, 262-263, 269-272, 275 - 279 defined, 8-9 majority us. minority mobility, 9 nitride-based devices, 164-177 in OLEDs, 261-273 Chemical hazards of manufacturing, 93-94 Child’s Law, 269-270 Chloride VPE, 24, 25 Codoping, 186-187 Contact design, 31, 99-100 in AlGaInP devices, 66 in OLEDs, 263-269 in transparent OLEDs, 288-291 Copper phthalocyanine, 222, 292-294 Crystal growth, 20-23 GaAs, 20-22 GaP, 20, 22-23 nitride-based LEDs, 133-152 Crystalline structure wurtzite, 7 zinc blend, 5 Current spreading in AlGalnP devices, 64-66 nitride-based LEDs, 182-183 in nitride-based LEDs, 179-180

D Defects deep levels, 9 and HSR recombination, 11 in nitride-based LEDs, 154- 157 residual impurities, 8, 163 Device fabrication, 30-32 AlGaInP devices, 99-101 OLED panels, 244-245 Dexter energy transfer, 280-282 Dicing, 100--101 9,lO-Diphenylanthracene doping, 225 Diphenyltetracene (DPT) doping, 225 Direct us. indirect recombination, 9- 12 Distributed Bragg reflectors (DBRs) in AlGaInP devices, 68-70 growth process, 91-92

Dopant materials, 7-8, 22 Ge, 148 Mg, 79,80, 146, 168-174, 176-177 0, 148, 167 rubrene, 224-226 Si, 78-79, 147--148, 163-164, 176, 186-187 Te, 79, 88 Zn, 78, 79-80, 146, 175, 185-187 Doping. See also Codoping AIGaInP, 60, 76, 78-81 AIGaN, 175-177 GaInN, 175, 185-187 GaN, 145-148, 162- 174 in G a N devices, 145-148 homojunction LEDs, 14 organic LEDs, 223-228,279-283 single-heterostructure LEDs, 15 Double-heterostructure (DH) LEDs, 15-16 AlGaInP devices, 58-63, 88-89 GaInN devices, 177-179, 185-190

E Edge-emitting laser diodes, 33, 34 Electronic signs, 37 Electrostatic discharge (ESD) sensitivity, 190-191 Environmental and safety issues arsine gas alternative, 76 MOVPE manufacturing of AlGaInP devices, 92-94 phosphine gas alternative, 76 Epitaxial film growth, 23-30 AlGaInP devices, 26, 72-73, 76-92 nitride films, 133-152 selective area epitaxy, 150-152 Epitaxial laterally overgrown GaN (ELOG), 152 Etching, 31-32

F Flat panel displays OILED, 291 -295 OLED, 239-248 Forster energy transfer, 280-281

309

INDEX I

Future trends LEDs, 40-44 OLEDS, 210-213

G GaAsP LEDs epitaxial growth methods, 26 evolution, 5 limitations and performance, 1-2 market growth, 44-45 structures, 17 wafer fabrication, 31 GaAs substrate, 3, 5, 17 crystal growth, 20-22 GaInN bandgap us. composition, 175 emission mechanism, 161-162 layer growth, 149-150 GaN buffer layers, 140-143 GaN substrate, 7, 26 GaP LEDs limitations and performance, 1-2 market growth, 44-45 structures, 17 GaP substrates, 3, 5, 17 crystal growth, 20, 22-23 G a P superstrates, 65 Gradient freeze (GF) method, 21 Grinding and cropping, 22

H Halide vapor phase epitaxy (HVPE), 133-134 Hall-effect measurements, 164-166 Hall-Shockley-Read (HSR) recombination, 11 Hazardous waste in manufacturing, 94 Hemispherical LEDs, 71-72 High-power LEDs, 116-118 Homojunction LEDs, 12-14 Horizontal Bridgman (HB) method, 21 Horizontal tube MOVPE reactors, 29 Hydride VPE, 25-26, 183 Hydrogen passivation in AlGaInP devices, 80-81 in GaN:Mg films, 168, 171-174

111-V roadmap, 5 Indirect recombination statistics, 11 Indium tin oxide (ITO) films as anode in OLEDs, 263 and current spreading, 66 InGaN LEDs eptiaxial growth methods, 26 luminous efficiency, 132-133 performance, 3-4, 33, 41 structures, 20 substrate materials, 7, 20

L Lattice constant (aO) AlGaInP alloys, 51-53 and device design, 5-7 Lattice match AlGaInP devices, 51-52, 86-87 defined, 5-7 nitride-based LEDs, 154- 156 LEC. See Liquid-encapsulated Czochralski (LEC) method Light-emitting diodes (LEDs). See also Organic light-emitting diodes (OLEDs); Polymeric light-emitting diodes (PLEDs) AlGaInP devices, 50-123 applications, 32-33, 35-39, 132-133 us. edge-emitting diodes, 33, 34 future trends, 40-44 historical overview, 1-4 market growth, 44-45 nitride devices, 130- 196 reliability, 16- 17 structures, 12-20 us. VCSELs, 33, 34 Light extraction, 16, 31, 180 in AlGaInP devices, 65-72, 100 in high-power LEDs, 117 Lighting and illumination applications, 32-33, 39,42-43 Liquid-encapsulated Czochralski (LEC) method, 21-22 Liquid-phase epitaxy (LPE), 23-24 Low-energy electron-beam irradiation (LEEBI), 171

310

INDEX

Luminous efficiency AlGaAs LEDs, 132-133 AIGaInP LEDs, 109-1 12, 132-133 and doping of OLEDs, 279-283 GalnN LEDs, 132-133, 194-195

M Manufacturing issues environmental and safety issues, 92-94 nitride-based devices, 135- 145, 196 reproducibility, 96-98 switching between AsH3 and PH3,87-88 uniformity, 95-96 Material properties, defined, 5- 12 Metal-organic vapor phase epitaxy (MOVPE), 27-30, 74-75, 134-135 AIGaInP devices, 73, 76-92 environmental and safety considerations, 92-94 growth temperature, 86-87 nitride devices, 135-145, 196 reactor designs, 28, 29-30, 135-138 reproducibility considerations, 96-98 switching between ASH, and PH,, 87-88 uniformity considerations, 95-96 Microcavity-filtered OLEDs, 298-299 Minority carriers in double-heterostructure LEDs, 15, 60-61 properties, 8-9 Multiple quantum wells (MQW) in AlGaTnP devices, 63-64 in GalnN devices, 190

N

Nitride-based LEDs, 130-196 AlGaN layer growth, 148-149 applications, 132-133 carrier transport properties, 164-177 device structures, 177-179, 185-190 doping, 145-148, 162-177, 185-187 electrostatic discharge sensitivity, 190- 1Y 1 epitaxial film growth, 133-152 epitaxial laterally overgrown GaN (ELOG), 152

GaInN emission mechanism, 161-162 GaInN layer growth, 149-150 growth parameters, 145 historical background, 131 low production yield, 196 luminescence properties, 157-162 n-type layer conductivity, 182-183 nucleation layers on sapphire substrate, 139- 143 p-type layer conductivity, 184-185 reactor designs, 135-138 reliability, 191-193 selective area epitaxy, 150-152 source materials, 133-135 structural defects, 154-157 substrate materials, 131, 138-139, 179-182 yellow luminescence, 161 Nucleation layers, 139- 143

0 OLED panels, 239-248 fabrication process, 244-245 flexibility, 286 material systems, 239-244 monochrome panels, 246 multi-color and full-color panels, 247-248 performance requirements, 256-257 stacked-RGB cells, 243, 247-248 Organic inverted light-emitting diodes (OILEDs), 291-295 Organic light-emitting diodes (OLEDs), 209-251, 255-301. See ulso Polymeric light-emitting diodes (PLEDs) AC us. DC operation, 235-236 aging treatment, 220-221 applications, 249-250, 256-257 carrier injection efficiency, 222-223, 263269 carrier transport properties, 261-273 chemical instability, 216 color-tunable, 295 -300 constant-current us. constant-voltage operation, 231-236 crosstalk reduction, 233-235 crystallization of amorphous films, 216 current-voltage characteristics, 267-273

31I

INDEX decay of luminance, 217-221 degradation mechanisms, 214-217, 283-285 dipole layers and carrier injection, 264-269 doping, 223-228, 279-283 double-heterojunction type, 221-224, 258 emission mechanism, 225 historical overview, 209-213, 255-256 IT0 buffer layer, 222 OLED panels, 239-248 organic inverted LEDs, 291-295 passive us. active matrix, 236-239 reliability, 214-231 shaped substrates, 285-286 single-heterojunction type, 215-217, 258-259 single-organic-layer type, 257-258, 263-273 structures, 217, 220-222, 251-261 temperature dependence, 228-23 1 transparent OLEDs, 287-291 white light emission, 242-243 Organometallic vapor phase epitaxy (OMVPE). See Metal-organic vapor phase epitaxy (MOVPE) Oxygen incorporation, 76-78, 81-85

P Packaging AlGaInP LEDs, 71-72, 101-102 LEDs, 32 nitride-based LEDs, 179-182 OLED panels, 244-245 Passive us. active matrix displays, 236-239 Perylene doping, 226, 227 Phosphine gas alternative, 76 Photoluminescence intensity AlGaInP films, 54-55 Alq, films, 259 GaN films, 158-162 Planetary MOVPE reactors, 29 Polishing, 22 Polymeric light-emitting diodes (PLEDs), 248-249 Power conversion efficiency edge-emitting laser diodes, 34

equation, 107 vertical-cavity surface-emitting laser diodes (VCSELs), 33, 34 Power signaling, 36-39

Q Quantum efficiency AlGaInP devices, 102-109 GaInN devices, 195 heterojunction OLEDs, 273-275 single-organic-layer OLEDs, 273 us. wavelength, 108-109 Quantum well devices multiple quantum well (MQW), 63-64, 190 single quantum well (SQW), 188-190 Quinacridone doping, 226, 227

R Radiative efficiency equation, 12 Radiative recombinations in direct bandgap devices, 9- 11 in indirect bandgap devices, 11 Reliability AlGaAs devices, 119-122 AlGaInP devices, 118-123 nitride-based LEDs, 191-193 organic LEDs, 214-231 Residual impurities and doping concentration, 8, 163 Resonant cavity LEDs (RCLEDs), 72 Rotating disk MOVPE reactors, 29 Rubrene doping, 224-226

S

Sapphire (A1203) substrate, 131, 139-143, 179-180, 181 Secondary ion mass spectrometry (SIMS) G a N films, 166-167 Selective area epitaxy, 150-152 Shaped substrates, 285-286 Silicon carbide (Sic) substrate, 143-144, 181-182 Single-heterostructure (SH) LEDs, 14-15

INDEX Single quantum well (SQW) devices, 188-190 Slicing and edge profiling, 22 Snell's law, 67 Source materials AlGaInP devices, 76-78,90 nitride-based devices, 133-135 switching between AsH3 and PH3,87-88 Space-charge-limited conduction, 269-270 Stacked color-tunable OLEDs, 299-300 Stacked-RGB cells, 243,247-248 Substrate materials GaAs, 3, 5, 51 -52 Gap, 3, 5, 70-71 nitride-based LEDs, 138-139, 179-182 sapphire (A1203), 131, 139- 143, 179-180, 181 silicon carbide (Sic), 143-144, 181-182 transparent substrate (TS) devices, 70-71, 100, 181-182 Surface texturing, 71

Trap-free space-charge-limited conduction, 278 Trapped-charge limited space-charge conduction, 270-272,276-278

U Uniformity in manufacturing, 95-96

V Vapor phase epitaxy (VPE), 24-27. See also Metal-organic vapor phase epitaxy (MOVPE) Vertical-cavity surface-emitting laser diodes (VCSELs), 33, 34 Visible LEDs. See Light-emitting diodes (LEDs)

W T Thermal management AlGaTnP devices, 112-1 16 organic light-emitting diodes (OLEDs), 217-220, 228-231 Traffic signals, 36-37, 116-117 Transparent substrate (TS) devices AIGaInP, 70-71 nitride-based LEDs, 180-182 OLEDS, 263,287-291 performance, 3 structures, 17 wafer bonding, 100

Wafer bonding, 100 Wafer fabrication, 20-23, 30-32, 99-101 Wavelength down-conversion, 298 White light generation, 42-43, 132, 193, 242-243 Window layers, 64-65, 68 growth process, 89-91

Z

Zone melting, 21

Contents of Volumes in This Series

Volume 1 Physics of 111-V Compounds C. Hilsum, Some Key Features of 111-V Compounds F. Bassani, Methods of Band Calculations Applicable to 111-V Compounds E. 0.Kane, The k-p Method K L. Bonch-Brueuich, Effect of Heavy Doping on the Semiconductor Band Structure D.Long, Energy Band Structures of Mixed Crystals of 111-V Compounds L. M. Roth and P. N. Argyres, Magnetic Quantum Effects S. M. Puri and T. H. Geballe, Thermomagnetic Effects in the Quantum Region W. M. Becker, Band Characteristics near Principal Minima from Magnetoresistance E. H. Putley, Freeze-Out Effects, Hot Electron Effects, and Submillimeter Photoconductivity in InSb H. Weiss, Magnetoresistance B. Ancker-Johnson, Plasma in Semiconductors and Semimetals

Volume 2 Physics of 111-V Compounds M. G. Holland, Thermal Conductivity S. I. Novkovu, Thermal Expansion U. Piesbergen, Heat Capacity and Debye Temperatures G. Giesecke, Lattice Constants J. R. Drubhle, Elastic Properties A. U.Mac Rae and G. U? Gobeli, Low Energy Electron Diffraction Studies R. Lee Mieher, Nuclear Magnetic Resonance B. Goldstein, Electron Paramagnetic Resonance T. S. Moss, Photoconduction in 111-V Compounds E. Antoncik and J. Tuuc,Quantum Efficiency of the Internal Photoelectric Effect in InSb G. W. Gobeli arid I. G. Allen, Photoelectric Threshold and Work Function P. S. Pershun, Nonlinear Optics in 111-V Compounds M. Gershenzon, Radiative Recombination in the 111-V Compounds F. Stern, Stimulated Emission in Semiconductors

313

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Volume 3 Optical of Properties 111-V Compounds M . Hass,Lattice Reflection W. G. Spitzer, Multiphonon Lattice Absorption D. L. Stierwalt and R. F Potter, Emittance Studies H. R. Philipp and H . Ehrenveich, Ultraviolet Optical Properties M . Curdona, Optical Absorption above the Fundamental Edge E. J. Johnson, Absorption near the Fundamental Edge J. 0.Dimmock, Introduction to the Theory of Exciton States in Semiconductors B. Lax and J. G. Muvroides, lnterband Magnetooptical Effects H . Y. Fun, Effects of Free Carries on Optical Properties E. D. Palik and G. 5. Wright, Free-Carrier Magnetooptical Effects R. H. Bube, Photoelectronic Analysis 5.0.Seraphin and H. E. Bennett, Optical Constants

Volume 4 Physics of 111-V Compounds N. A . Goryunova. A. S. Borschevskii, and D. N. Tretiakov, Hardness N. N. Sirota, Heats of Formation and Temperatures and Heats of Fusion of Compounds A"'BV D. L. Kendull, Diffusion A. G. Chynoweth, Charge Multiplication Phenomena R. W. Keyes, The Effects of Hydrostatic Pressure on the Properties of 111-V Semiconductors L. W. Aukerman, Radiation Effects N. A. Goryunova, F. P. Kesamanly, and D. N. Nasledov, Phenomena in Solid Solutions R. T. Bate, Electrical Properties of Nonuniform Crystals

Volume 5 Infrared Detectors H . Levinstrin, Characterization of Infrared Detectors P. W. Kruse, Indium Antimonide Photoconductive and Photoelectromagnetic Detectors M . 5.Prince, Narrowband Self-Filtering Detectors I. Melngalis and T. C. Harman, Single-Crystal Lead-Tin Chalcogenides D. Long and J. L. Schmidt, Mercury-Cadmium Telluride and Closely Related Alloys E. H. Putley, The Pyroelectric Detector N. 5.Stevens, Radiation Thermopiles R. J. Keyes and T. M. Quist, Low Level Coherent and Incoherent Detection in the Infrared M. C. Teich, Coherent Detection in the Infrared F. R. Arams, E. W. Surd, B. J. Peyton. and F. P. Puce, Infrared Heterodyne Detection with Gigahertz IF Response H. S. Sommers, Jr., Macrowave-Based Photoconductive Detector R. Sehr and R. Zuleeg, Imaging and Display

Volume 6 Injection Phenomena M. A . Lumpert and R. B. Schilling, Current Injection in Solids: The Regional Approximation Method R. Williams, Injection by Internal Photoemission A . M. Barnett. Current Filament Formation

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315

R. Baron and J. W Mayer, Double Injection in Semiconductors W. Ruppel, The Photoconductor-Metal Contact

Volume 7 Application and Devices Part A J. A. Copeland and S. Knight, Applications Utilizing Bulk Negative Resistance F A. Padovani, The Voltage-Current Characteristics of Metal-Semiconductor Contacts P. L. Hower, W. W. Hooper, B. R. Cairns, R. D. Fairman, and D. A. Tremere, The GaAs Field-Effect Transistor M. H. White, MOS Transistors G.R. Antell, Gallium Arsenide Transistors T. L. Tansley, Heterojunction Properties

Part B T. Misawa, IMPATT Diodes H. C. Okean, Tunnel Diodes R. B. Campbell and Hung-Chi Chang, Silicon Junction Carbide Devices R. E. Enstrom, H. Kressel, and L. Krassner, High-Temperature Power Rectifiers of GaAs, -xPx

Volume 8 Transport and Optical Phenomena R. J. Stirn, Band Structure and Galvanomagnetic Effects in 111-V Compounds with Indirect Band Gaps R. W. Ure, Jr., Thermoelectric Effects in 111-V Compounds H. Piller, Faraday Rotation H. Barry Bebb and E. W. Williams, Photoluminescence I: Theory E. W. Williams and H. Barry Bebb, Photoluminescence 11: Gallium Arsenide

Volume 9 Modulation Techniques B. 0. Seraphin, Electroreflectance R. L. Aggarwal, Modulated Interband Magnetooptics D. I? Blossey and Paul Handler, Electroabsorption B. Batz, Thermal and Wavelength Modulation Spectroscopy I. Balslev, Piezopptical Effects D. E. Aspnes and N . Bottka, Electric-Field Effects on the Dielectric Function of Semiconductors and Insulators

Volume 10 Transport Phenomena R. L. Rhode, Low-Field Electron Transport J. D. Wiley, Mobility of Holes in 111-V Compounds C. M.Wove and G. E. Stillman, Apparent Mobility Enhancement in Inhomogeneous Crystals R. L. Petersen, The Magnetophonon Effect

316

CONTENTS OF VOLUMES IN THISSERIES

Volume 11 Solar Cells H. J. Hovel, Introduction; Carrier Collection, Spectral Response, and Photocurrent; Solar Cell Electrical Characteristics; Efficiency; Thickness; Other Solar Cell Devices; Radiation Effects; Temperature and Intensity; Solar Cell Technology

Volume 12 Infrared Detectors (11) W. L. Eisemun, J. D. Merriam, and R. F. Potter, Operational Characteristics of Infrared Photodetectors P. R. Bratt, Impurity Germanium and Silicon Infrared Detectors E. H. Putley, InSb Submillimeter Photoconductive Detectors G. E. Stillman, C. M . Wove, and J. 0. Dimmock, Far-Infrared Photoconductivity in High Purity GaAs G. E. Stillman and C. M. Wolfe, Avalanche Photodiodes P. L. Richards, The Josephson Junction as a Detector of Microwave and Far-Infrared Radiation E. H. Putley, The Pyroelectric Detector- An Update

Volume 13 Cadmium Telluride K. Zanio, Materials Preparations; Physics; Defects; Applications

Volume 14 Lasers, Junctions, Transport N. Holonyak, Jr. and M. H. Lee, Photopumped 111-V Semiconductor Lasers H. Kressel and J. K. Buller, Heterojunction Laser Diodes A Van der Ziel, Space-Charge-Limited Solid-state Diodes P. J. Price, Monte Carlo Calculation of Electron Transport in Solids

Volume 15 Contacts, Junctions, Emitters B. L. Sharma, Ohmic Contacts to 111-V Compounds Semiconductors A . Nussbaum, The Theory of Semiconducting Junctions J. S. Escher. NEA Semiconductor Photoemitters

Volume 16 Defects, (HgCd)Se, (HgCd)Te H. Kressel, The Effect of Crystal Defects on Optoelectronic Devices C. R. Whitsett, J. G. Bruerman, and C. J. Summers, Crystal Growth and Properties of Hg, -.Cd,Se alloys M. H. Weiler, Magnetooptical Properties of Hg, -xCd,Te Alloys P. W. Kruse and J. G. Ready, Nonlinear Optical Effects in Hg, -xCd,Te

Volume 17 CW Processing of Silicon and Other Semiconductors J. F: Gibbons, Beam Processing of Silicon A. Lietoila, R. B. Gold, J. F. Gibbons, and L. A. Christel, Temperature Distributions and Solid Phase Reaction Rates Produced by Scanning CW Beams

CONTENTS OF VOLUMES RV THISSERIES

317

A. Leitoila and J. F. Gibbons, Applications of CW Beam Processing to Ion Implanted Crystalline Silicon N. M. Johnson, Electronic Defects in CW Transient Thermal Processed Silicon K. F. Lee, T. J. Stuhz, and J. F. Gibbons, Beam Recrystallized Polycrystalline Silicon: Properties, Applications, and Techniques T. Shibata, A. Wakita, T. W. Sigmon, and J. I;: Gibbons, Metal-Silicon Reactions and Silicide Y. i. Nissim and J. F. Gibbons, CW Beam Processing of Gallium Arsenide

Volume 18 Mercury Cadmium Telluride P. W Kruse, The Emergence of (Hg, _=Cd,)Te as a Modern Infrared Sensitive Material H. E. Hirsch, S. C. Liang, and A. G. White, Preparation of High-Purity Cadmium, Mercury, and Tellurium W. F. H. Micklethwaite, The Crystal Growth of Cadmium Mercury Telluride P. E. Petersen, Auger Recombination in Mercury Cadmium Telluride R. M. Broudy and V. J. Mazurczyck, (HgCd)Te Photoconductive Detectors M. 8. Reine, A. K. Soad, and T. J. Tredwell, Photovoltaic Infrared Detectors M. A . Kinch, Metal-Insulator-Semiconductor Infrared Detectors

Volume 19 Deep Levels, GaAs, Alloys, Photochemistry G. F. Neumark and K Kosai, Deep Levels in Wide Band-Gap 111-V Semiconductors D. C. Look, The Electrical and Photoelectronic Properties of Semi-Insulating GaAs R F. Brebrick, Ching-Hua Su, and Pok-Kai Liao, Associated Solution Model for Ga-In-Sb and Hg-Cd-Te Y. Ya. Gurevich and Y. V. Pleskon, Photoelectrochemistry of Semiconductors

Volume 20 Semi-Insulating GaAs R. N. Thomas, H. M. Hobgood. G. W. Eldridge, D. L. Barrett, T. T. Braggins, L. B. Ta, and S. K. Wang, High-Purity LEC Growth and Direct Implantation of GaAs for Monolithic Microwave Circuits C. A. Stolte, Ion Implantation and Materials for GaAs Integrated Circuits C. G. Kirkpatrick, R. T. Chen, D. E. Holmes, P. M. Asbeck, K. R. Elliott, R. D. Fairman, and J. R. Oliver, LEC GaAs for Integrated Circuit Applications J. S. Blakemore and S. Rahimi, Models for Mid-Gap Centers in Gallium Arsenide

Volume 21 Hydrogenated Amorphous Silicon Part A J. I. Pankove, Introduction M. Hirose, Glow Discharge; Chemical Vapor Deposition Y. Uchida, di Glow Discharge T. D. Moustakas, Sputtering I. Yamada, Ionized-Cluster Beam Deposition B. A. Scott, Homogeneous Chemical Vapor Deposition

318

CONTENTS OF VOLUMESIN THISSERIES

F. J. Kumpas, Chemical Reactions in Plasma Deposition

P. A. Longeway, Plasma Kinetics H. A. Weakliem, Diagnostics of Silane Glow Discharges Using Probes and Mass Spectroscopy L. Ghtrmun, Relation between the Atomic and the Electronic Structures A. Chenevas-Paule, Experiment Determination of Structure S. Minomura, Pressure Effects on the Local Atomic Structure D.Adler, Defects and Density of Localized States

Part B J I. Pankove, Introduction G. D.Cody, The Optical Absorption Edge of a-Si:H N. M. Amer and W. B. Jackson, Optical Properties of Defect States in a-Si: H P. J. Zunzucchi, The Vibrational Spectra of a-Si: H Y. Hamakawa, Electroreflectance and Electroabsorption J. S. Lannin, Raman Scattering of Amorphous Si, Ge, and Their Alloys R. A. Street, Luminescence in a-Si: H R. S. Crandall, Photoconductivity J. Tauc, Time-Resolved Spectroscopy of Electronic Relaxation Processes P. E. Vanier, IR-Induced Quenching and Enhancement of Photoconductivity and Photo luminescence H.Schude, Irradiation-lnduced Metastable Effects L. Ley, Photoelectron Emission Studies

Part C J. I. Pankove, Introduction J. D. Cohen, Density of States from Junction Measurements in Hydrogenated Amorphous Silicon P. C. Taylor, Magnetic Resonance Measurements in a-SL H X Morigaki, Optically Detected Magnetic Resonance J. Dresner, Carrier Mobility in a-Si: H T. Tiedje, Information about band-Tail States from Time-of-Flight Experiments A. R. Moore, DilTusion Length in Undoped a-Si: H W. Beyer and J. Overhof, Doping Effects in a-Si: H H. Fritzche, Electronic Properties of Surfaces in a-Si:H C. R. Wronski, The Staebler-Wronski Effect R. . I Nemanich, Schottky Barriers on a-Si: H B. Abeles and T. Tzedje, Amorphous Semiconductor Superlattices

Part 11 J. I. Pankove, Introduction D. E. Curlson, Solar Cells G. A . Swarfz, Closed-Form Solution of I-V Characteristic for a a-Si: H Solar Cells I. Shimizu, Electrophotography S. Ishioka, Image Pickup Tubes

CONTENTS OF VOLUMES IN THISSERIES

319

P. G. LeComber and W E. Spear, The Development o f the a-Si: H Field-Effect Transistor and Its Possible Applications D. G. Ast, a-Si:H FET-Addressed LCD Panel S. Kaneko, Solid-state Image Sensor M. Matsumura, Charge-Coupled Devices M. A . Bosch, Optical Recording A. D’Amico and G. Fortunato, Ambient Sensors H. Kukimoto, Amorphous Light-Emitting Devices R. J. Phelun, Jr., Fast Detectors and Modulators J. I. Pankove, Hybrid Structures P. G. LeComber, A. E. Owen, W. E. Spear, J. Hajto, and W. K. Choi, Electronic Switching in Amorphous Silicon Junction Devices

Volume 22 Lightwave Communications Technology Part A K. Nakajima, The Liquid-Phase Epitaxial Growth of InGaAsP W. T. Tsang, Molecular Beam Epitaxy for 111-V Compound Semiconductors G. E. Stringfellow, Organometallic Vapor-Phase Epitaxial Growth of 111-V Semiconductors G. Beuchet, Halide and Chloride Transport Vapor-Phase Deposition of InGaAsP and GaAs M. Razeghi, Low-Pressure Metallo-Organic Chemical Vapor Deposition of Ga,In, _,ASP, - y Alloys P. M. Petrox Defects in 111-V Compound Semiconductors

Part B J. P. van der Ziel, Mode Locking of Semiconductor Lasers K Y. Lau and A. Yariv, High-Frequency Current Modulation of Semiconductor Injection Lasers C. H. Henry, Special Properties of Semiconductor Lasers Y. Suematsu, K. Kishino, S. Arai, and F. Koyama, Dynamic Single-Mode Semiconductor Lasers with a Distributed Reflector W T. Tsang, The Cleaved-Coupled-Cavity (C3) Laser

Part C R. J. Nelson and N. K Dutta, Review of InGaAsP InP Laser Structures and Comparison of Their Performance N. Chinone and M. Nakamura, Mode-Stabilized Semiconductor Lasers for 0.7-0.8- and 1.1-1.6-pm Regions Y. Horikoshi, Semiconductor Lasers with Wavelengths Exceeding 2 pm B. A. Dean and M. Dixon, The Functional Reliability of Semiconductor Lasers as Optical Transmitters R. H. Saul, T. P. Lee, and C. A. Burus, Light-Emitting Device Design C. L. Z i p f l , Light-Emitting Diode-Reliability T.P. Lee and T. Li, LED-Based Multimode Lightwave Systems K. Ogawa, Semiconductor Noise-Mode Partition Noise

320

CONTENTS OF VOLUMESIN THIS SERIES

Part D F. Capasso, The Physics of Avalanche Photodiodes T. P. Pearsull and M. A. Pollack, Compound Semiconductor Photodiodes T. Kancda, Silicon and Germanium Avalanche Photodiodes S. R. Forrest, Sensitivity of Avalanche Photodetector Receivers for High-Bit-Rate LongWavelength Optical Communication Systems J. C Campbell, Phototransistors for Lightwave Communications

Part E S. Wang, Principles and Characteristics of Integrable Active and Passive Optical Devices S. Margalit and A. Yariv, Integrated Electronic and Photonic Devices T. Mukui, Y. Yumamoto, and T. Kimuru, Optical Amplification by Semiconductor Lasers

Volume 23 Pulsed Laser Processing of Semiconductors R. F. Wood, C. W. White, and R. T. Young, Laser Processing of Semiconductors: An Overview C. W. White, Segregation, Solute Trapping, and Supersaturated Alloys G. E. Jellison, Jr., Optical and Electrical Properties of Pulsed Laser-Annealed Silicon R. F. Wood and G. E. Jellison, Jr., Melting Model of Pulsed Laser Processing R. E Wood and F. W: Young, Jr., Nonequilibrium Solidification Following Pulsed Laser Melting D. H. Lowndes and G. E. Jellison, Jr., Time-Resolved Measurement During Pulsed Laser Irradiation of Silicon D. M. Zebner, Surface Studies of Pulsed Laser Irradiated Semiconductors D.H.Lowndes, Pulsed Beam Processing of Gallium Arsenide R. B. James, Pulsed CO, Laser Annealing of Semiconductors R. T. Young and R. F. Wood, Applications of Pulsed Laser Processing

Volume 24 Applications of Multiquantum Wells, Selective Doping, and Superlattices C. Weisbuch, Fundamental Properties of 111-V Semiconductor Two-Dimensional Quantized Structures: The Basis for Optical and Electronic Device Applications H. Morkoc and H . Unlu, Factors Affecting the Performance of (Al, Ga)As/GaAs and (Al, Ga)As/InGaAs Modulation-Doped Field-Effect Transistors: Microwave and Digital Applications N. T. Linh, Two-Dimensional Electron Gas FECTs: Microwave Applications M. Abe er al., Ultra-High-speed HEMT Integrated Circuits D.S. Chemla, D.A. B. Miller, und P. W Smith, Nonlinear Optical Properties of Multiple Quantum Well Structures for Optical Signal Processing F. Cupasso, Graded-Gap and Superlattice Devices by Band-Gap Engineering W. T. Tsung, Quantum Confinement Heterostructure Semiconductor Lasers G. C. Osbourn et al., Principles and Applications of Semiconductor Strained-Layer Superlattices

CONTENTS OF VOLUMES IN THISSERIES

321

Volume 25 Diluted Magnetic Semiconductors W Giriai and J. K. Furdyna, Crystal Structure, Composition, and Materials Preparation of Diluted Magnetic Semiconductors W. M. Eecker, Band Structure and Optical Properties of Wide-Gap AF-xMn,B,v Alloys at Zero Magnetic Field S. Oseroff and P. H. Keesom, Magnetic Properties: Macroscopic Studies T. Giebultowicz and T. M. Holden, Neutron Scattering Studies of the Magnetic Structure and Dynamics of Diluted Magnetic Semiconductors J. Kossut, Band Structure and Quantum Transport Phenomena in Narrow-Gap Diluted Magnetic Semiconductors C. Riquaux, Magnetooptical Properties of Large-Gap Diluted Magnetic Semiconductors J. A . Guj, Magnetooptical Properties of Large-Gap Diluted Magnetic Semiconductors J. Mycielski, Shallow Acceptors in Diluted Magnetic Semiconductors: Splitting, Boil-off, Giant Negative Magnetoresistance A . K Ramadas and R. Rodriquez, Raman Scattering in Diluted Magnetic Semiconductors P. A. Wol& Theory of Bound Magnetic Polarons in Semimagnetic Semiconductors

Volume 26 111-V Compound Semiconductors and Semiconductor Properties of Superionic Materials 2. Yuunxi, 111-V Compounds

H. V. Winston, A. T. Hunter, H. Kimura, and R. E. Lee, InAs-Alloyed GaAs Substrates for Direct Implantation P. K. Ehattacharya and S. Dhar, Deep Levels in IIILV Compound Semiconductors Grown by MBE Y. Ya. Gurevich and A . K Ivanov-Shits, Semiconductor Properties of Supersonic Materials

Volume 27 High Conducting Quasi-One-DimensionalOrganic Crystals E. M. Conwell, Introduction to Highly Conducting Quasi-One-Dimensional Organic Crystals I. A. Howard, A Reference Guide to the Conducting Quasi-One-Dimensional Organic

Molecular Crystals J. P. Pouquet, Structural Instabilities E. M . Conwell, Transport Properties C. S.Jacobsen, Optical Properties J. C. Scott, Magnetic Properties L. Zuppiroli, Irradiation Effects: Perfect Crystals and Real Crystals

Volume 28 Measurement of High-speed Signals in Solid State Devices J. Frey and D. Ioannou, Materials and Devices for High-speed and Optoelectronic Applications H. Schumacher and E. Strid, Electronic Wafer Probing Techniques D. H. Auston, Picosecond Photoconductivity: High-speed Measurements of Devices and Materials J. A. Valdmanis, Electro-Optic Measurement Techniques for Picosecond Materials, Devices, and Integrated Circuits. J. M. Wiesenfeld and R K. Jain, Direct Optical Probing of Integrated Circuits and High-speed Devices G. Plows, Electron-Beam Probing A. M. Weiner and R. E. Marcus, Photoemissive Probing

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CONTENTS OF VOLUMES IN THISSERIES

Volume 29 Very High Speed Integrated Circuits: Gallium Arsenide LSI M. Kuzuhara and T Nazaki, Active Layer Formation by Ion Implantation H. Hasimoio, Focused Ion Beam Implantation Technology T. Nozaki and A . Higashisaka, Device Fabrication Process Technology M. In0 and T. Takada, GaAs LSI Circuit Design M. Hirayama, M. Ohmnri, and K. Yamasaki, GaAs LSI Fabrication and Performance

Volume 30 Very High Speed Integrated Circuits Heterostructure H. Watanabe, T. Mizutani, and A . Usui, Fundamentals of Epitaxial Growth and Atomic Layer Epitaxy

S. Hiyamizu, Characteristics of Two-Dimensional Electron Gas in 111-V Compound Heterostructures Grown by MBE T. Nukanisi, Metalorganic Vapor Phase Epitaxy for High-Quality Active Layers T. Nimura, High Electron Mobility Transistor and LSI Applications T. Sugeta and T. Ishibashi, Hetero-Bipolar Transistor and LSI Application I€. Matsueda, T. Tanaka, and M. Nakamura, Optoelectronic Integrated Circuits

Volume 3 1 Indium Phosphide: Crystal Growth and Characterization J. P. Forges, Growth of Discoloration-free InP M. J. McCollum and G. E. Stillman, High Purity InP Grown by Hydride Vapor Phase Epitaxy T. Inada and T. Fukuda, Direct Synthesis and Growth of Indium Phosphide by the Liquid Phosphorous Encapsulated Czochralski Method 0. Oda, K. Katagiri, K. Shinohara, S. Katsura, Y. Takahashi, K. Kainosho, K. Kohiro, and R. Hirano, InP Crystal Growth, Substrate Preparation and Evaluation K. Tada, M. Tatsumi, M. Morioka, T. Araki, and T. Kawase, InP Substrates: Production and Quality Control M. Razeghi, LP-MOCVD Growth, Characterization, and Application of InP Material T. A. Kennedy and P. J. Lin-Chung, Stoichiometric Defects in InP

Volme 32 Strained-Layer Superlattices: Physics T. P. Pearsall, Strained-Layer Superlattices F. H. Pollack, Effects of Homogeneous Strain on the Electronic and Vibrational Levels in Semiconductors J. Y Marzin, J. M. Gerhrd, P. Voisin, and J. A. Brum, Optical Studies of Strained 111-V Heterolayers R. People and S. A . Jackson, Structurally Induced States from Strain and Confinement M. Jaros, Microscopic Phenomena in Ordered Superlattices

Volume 33 Strained-Layer Superlattices: Materials Science and Technology R. Hull and J. C. Bean, Principles and Concepts of Strained-Layer Epitaxy W. J. Schaff; P. J. Tasker, M. C. Foi,sy, and L. F. Eastman, Device Applications of Strained-Layer Epitaxy

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323

S. T. Picraux, B. L. Doyle, and J. Y. Tsao, Structure and Characterization of Strained-Layer Superlattices E. Kasper and F. Schaffer, Group IV Compounds D. L. Martin, Molecular Beam Epitaxy of IV-VI Compounds Heterojunction R. L. Gunshor, L. A. Kolodziejski, A. V. Nurmikko, and N. Otsuka, Molecular Beam Epitaxy of 11-VI Semiconductor Microstructures

Volume 34 Hydrogen in Semiconductors J. 1. Pankove and N. M. Johnson, Introduction to Hydrogen in Semiconductors C. H. Seager, Hydrogenation Methods J. I. Pankove, Hydrogenation of Defects in Crystalline Silicon J. W. Corbett, P. Deak, U . V. Desnica, and S. J . Pearton, Hydrogen Passivation of Damage Centers in Semiconductors S. J. Pearton, Neutralization of Deep Levels in Silicon J. I. Pankove, Neutralization of Shallow Acceptors in Silicon N . M. Johnson, Neutralization of Donor Dopants and Formation of Hydrogen-Induced Defects in n-Type Silicon M . Stavola and S. J. Pearton, Vibrational Spectroscopy of Hydrogen-Related Defects in Silicon A. D.Marwick, Hydrogen in Semiconductors: Ion Beam Techniques C. Herring and N. M. Johnson, Hydrogen Migration and Solubility in Silicon E. E. Huller, Hydrogen-Related Phenomena in Crystalline Germanium J. Kakalios, Hydrogen Diffusion in Amorphous Silicon J. Chevalier, B. Clerjaud, and B. Pajot, Neutralization of Defects and Dopants in 111-V Semiconductors G. G. DeLeo and W. B. Fowler, Computational Studies of Hydrogen-Containing Complexes in Semiconductors R. F. Kiep and T. L. Estle, Muonium in Semiconductors C. G. Van de Walk, Theory of Isolated Interstitial Hydrogen and Muonium in Crystalline Semiconductors

Volume 35 Nanostructured Systems M. Reed, Introduction H. van Houten, C. W. J. Beenakker. and B. J. van Wees, Quantum Point Contacts G. Timp, When Does a Wire Become an Electron Waveguide? M . Biittiker, The Quantum Hall Effects in Open Conductors W Hansen, J. P. Kotthaus, and U. Merkt, Electrons in Laterally Periodic Nanostructures

Volume 36 The Spectroscopy of Semiconductors D.Heirnan, Spectroscopy of Semiconductors at Low Temperatures and High Magnetic Fields A. K Nurmikko, Transient Spectroscopy by Ultrashort Laser Pulse Techniques A. K. Ramdas and S. Rodriguez, Piezospectroscopy of Semiconductors 0. J. Glembocki und B. V. Shanabrook, Photoreflectance Spectroscopy of Microstructures D. G. Seiler, C. L. Littler, and M. H. Wiler, One- and Two-Photon Magneto-Optical

Spectroscopy of InSb and Hg, -xCd,Te

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CONTENTS OF VOLUMES IN THISSERIES

Volume 37 The Mechanical Properties of Semiconductors A.-E. Chen, A. Sher and W. T. Yost, Elastic Constants and Related Properties of Semiconductor Compounds and Their Alloys D. R. Clarke, Fracture of Silicon and Other Semiconductors H. Siethof, The Plasticity of Elemental and Compound Semiconductors S. Guruswamy, K. T. Faber and J. P. Hirth, Mechanical Behavior of Compound Semiconductors S. Mahajan, Deformation Behavior of Compound Semiconductors J. P. Hirth, Injection of Dislocations into Strained Multilayer Structures D. Kendull, C. E. Fleddermunn, and K. J. Malfoy,Critical Technologies for the Micromachining of Silicon I. M a m b a and K. Mokuyu, Processing and Semiconductor Thermoelastic Behavior

Volume 38 Imperfections in IlIlV Materials U.Scherz and M. Schefler, Density-Functional Theory of sp-Bonded Defects in Ill/V Semiconductors

M. Kaminska and E. R. Weber, El2 Defect in GaAs D. C. Look, Defects Relevant For Compensation in Semi-Insulating GaAs

R. C. Newman, Local Vibrational Mode Spectroscopy of Defects in lll/V Compounds A . M. Hennel, Transition Metals in IIl/V Compounds K. J. Malloy and K. Khachaturyan, DX and Related Defects in Semiconductors F Swaminuthan and A. S. Jordan, Dislocations in lll/V Compounds K. W. Nauka, Deep Level Defects in the Epitaxial IIIjV Materials

Volume 39 Minority Carriers in 111-V Semiconductors: Physics and Applications N. K. Dutta, Radiative Transitions in GaAs and Other ITI.-V Compounds R. K Ahrenkiel, Minority-Carrier Lifetime in 111-V Semiconductors 7: Furuta, High Field Minority Electron Transport in p-CaAs M. S. Lundsfrom,Minority-Carrier Transport in 111-V Semiconductors R. A. Abram, Effects of Heavy Doping and High Excitation on the Band Structure of GaAs D. Yevick and W Bardyszewski, An Introduction to Non-Equilibrium Many-Body Analyses of Optical Processes in Ill-V Semiconductors

Volume 40 Epitaxial Microstructures E. F. Schubert, Delta-Doping of Semiconductors: Electronic, Optical, and Structural Properties of Materials and Devices A. Gossard, M. Sundaram, and P. Hopkins, Wide Graded Potential Wells P. Petrox Direct Growth of Nanometer-Size Quantum Wire Superlattices E. Kapun, Lateral Patterning of Quantum Well Heterostructures by Growth of Nonplanar Substrates H. Temkin, D. Gershoni. and M. Punish, Optical Properties of Ga, -.In,As/InP Quantum Wells

CONTENTS OF VOLUMES IN THIS SERIES

325

Volume 41 High Speed Heterostructure Devices F. Capasso, E Eeltram, S. Sen, A. Pahlevi, and A. Y. Cho, Quantum Electron Devices: Physics and Applications P. Solomon, D. J. Frank. S. L. Wright, and F. Canora, GaAs-Gate Semiconductor-lnsulatorSemiconductor FET M. H. Hashemi and U.K. Mishra, Unipolar InP-Based Transistors R. Kiehl, Complementary Heterostructure FET Integrated Circuits T. Ishibashi, GaAs-Based and InP-Based Heterostructure Bipolar Transistors H. C. Liu and T. C. L. G. Sollner, High-Frequency-Tunneling Devices H. Ohnishi, T. More, M. Takatsu, K. Imamura, and N. Yokoyama, Resonant-Tunneling Hot-Electron Transistors and Circuits

Volume 42 Oxygen in Silicon F. Shimura, Introduction to Oxygen in Silicon W. Lin, The Incorporation of Oxygen into Silicon Crystals T. . I Schaffner and D. X Schroder, Characterization Techniques for Oxygen in Silicon W. M. Bullis, Oxygen Concentration Measurement S. M. Hu, Intrinsic Point Defects in Silicon E. Pujot, Some Atomic Configurations of Oxygen J. Michel and L. C. Kimerling, Electical Properties of Oxygen in Silicon R. C. Newman and R. Jones, Diffusion of Oxygen in Silicon T. Y. Tan and W. J. Taylor, Mechanisms of Oxygen Precipitation: Some Quantitative Aspects M. Schrems, Simulation of Oxygen Precipitation X Simino and I. Yonenaga, Oxygen Effect on Mechanical Properties W. Bergholz, Grown-in and Process-Induced Effects F. Shimura, Intrinsic/Internal Gettering H. Tsuya, Oxygen Effect on Electronic Device Performance

Volume 43 Semiconductors for Room Temperature Nuclear Detector Applications R. B. James and T. E. Schlesinger, Introduction and Overview L. S. Darken and C. E. Cox, High-Purity Germanium Detectors A. Burger, D. Nason, L. Van den Berg, and M. Schieber, Growth of Mercuric Iodide X . J. Bao, T. E. Schlesinger, and R. E. James, Electrical Properties of Mercuric Iodide X J. Eao, R. B. James, and T. E. Schlesinger, Optical Properties of Red Mercuric Iodide M. Hage-Ali and P. Szferf, Growth Methods of CdTe Nuclear Detector Materials M. Hage-Ali and P Srfert, Characterization of CdTe Nuclear Detector Materials M. Hage-Ali and P. Szfert, CdTe Nuclear Detectors and Applications R B. James, T. E. Schlesinger, J. Lund, and M. Schieber, Cd, -xZn,Te Spectrometers for Gamma and X-Ray Applications D. S. McGregor, J. E. Kammeraad, Gallium Arsenide Radiation Detectors and Spectrometers J. C. Lund. F. Olschner, and A. Burger, Lead Iodide M . R. Squillante, and K. S. Shah, Other Materials: Status and Prospects V. M. Gerrish, Characterization and Quantification of Detector Performance J. S. Iwanczyk and B. E. Putt, Electronics for X-ray and Gamma Ray Spectrometers M. Schieber, R E. James, and T. E. Schlesinger, Summary and Remaining Issues for Room Temperature Radiation Spectrometers

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CONTENTS OF VOLUMES IN THISSERIES

Volume 44 11-IV BluelGreen Light Emitters: Device Physics and Epitaxial Growth J. Han and R. L. Gunshor, MBE Growth and Electrical Properties of Wide Bandgap ZnSe-based 11-VI Semiconductors S. Fujita and S. Fujita, Growth and characterization of ZnSe-based 11- VI Semiconductors by MOVPE E. Ho and L. A. Kolodziejski, Gaseous Source UHV Epitaxy Technologies for Wide Bandgap 11-VI Semiconductors C. G. Van de Walk, Doping of Wide-Band-Gap 11-VI Compounds--- Theory R. Cingolani, Optical Properties of Excitons in ZnSe-Based Quantum Well Heterostructures A. Ishibashi and A. V. Nurmikko, 11-VI Diode Lasers: A Current View of Device Performance and Issues S. Guha and J. Petruzello, Defects and Degradation in Wide-Gap 11-VI-based Structures and Light Emitting Devices

Volume 45 Effect of Disorder and Defects in Ion-ImplantedSemiconductors: Electrical and Physiochemical Characterization IT. Ryssel, Ion Implantation into Semiconductors: Historical Perspectives You-Nian Wang and Teng-Cai Ma, Electronic Stopping Power for Energetic Ions in Solids S. 7'. Nakagawa, Solid Effect on the Electronic Stopping of Crystalline Target and Application to Range Estimation G.Miiller, S. Kalbitzer and G. N. Greaves, Ion Beams in Amorphous Semiconductor Research J. Boussey-Said, Sheet and Spreading Resistance Analysis of Ion Implanted and Annealed Scmiconductors M. L. Polignano and C. Queirolo, Studies of the Stripping Hall Effect in Ion-Implanted Silicon J. Stoemenos, Transmission Electron Microscopy Analyses R. Nipoti and M. Servidori, Rutherford Backscattering Studics of Ion Implanted Semiconductors P. Zaumseil, X-ray Diffraction Techniques

Volume 46 Effect of Disorder and Defects in Ion-Implanted Semiconductors: Optical and Photothermal Characterization M. Fried, i? Lohner and J. Gyulai, Ellipsometric Analysis A. Seas and C. Christofides, Transmission and Reflection Spectroscopy on Ion Implanted Semiconductors A. Othonos and C. Christofides, Photoluminescence and Raman Scattering of Ion Implanted

Semiconductors. Influence of Annealing C. Christofides, Photomodulated Thermoreflectance Investigation of Implanted Wafers. Annealing Kinetics of Defects U. Zammit, Photothermal Deflection Spectroscopy Characterization of Ion-Implanted and Annealed Silicon Films A . Mandelis, '4. Budirnan and M. Vargas, Photothermal Deep-Level Transient Spectroscopy of Impurities and Defects in Semiconductors R. Kalish and S. Charbonneau, Ion Implantation into Quantum-Well Structures A. M. Myasnikov and N. N. Gerasimenko, Ion Implantation and Thermal Annealing of 111-V Compound Semiconducting Systems: Some Problems of HI-V Narrow Gap Semiconductors

CONTENTS OF VOLUMES IN THISSERIES

327

Volume 47 Uncooled Infrared Imaging Arrays and Systems R. G. Buser and M. P. Tompsett, Historical Overview P. W Kruse, Principles of Uncooled Infrared Focal Plane Arrays R. A. Wood, Monolithic Silicon Microbolometer Arrays C. M . Hanson, Hybrid Pyroelectric-Ferroelectric Bolometer Arrays D. L. Polla and J . R. Choi, Monolithic Pyroelectric Bolometer Arrays N . Teranishi, Thermoelectric Uncooled Infrared Focal Plane Arrays M . F. Tompsett, Pyroelectric Vidicon ?: W Kenny, Tunneling Infrared Sensors J . R. Vig,R. L. Filler and Y: Kim, Application of Quartz Microresonators to Uncooled Infrared Imaging Arrays P. W Kruse, Application of Uncooled Monolithic Thermoelectric Linear Arrays to Imaging Radiometers

Volume 48 High Brightness Light Emitting Diodes G. B. Stringfellow, Materials Issues in High-Brightness Light-Emitting Diodes M . G. Craford, Overview of Device issues in High-Brightness Light-Emitting Diodes F. M . Steranka, AlGaAs Red Light Emitting Diodes C. H. Chen, S. A. Stockman, M . J . Peanasky, and C. P. Kuo, OMVPE Growth of AlGaInP for High Efficiency Visible Light-Emitting Diodes F. A. Kish and R. M . Fletcher, AlGaInP Light-Emitting Diodes M . W Hodapp, Applications for High Brightness Light-Emitting Diodes I. Akasaki and H. Amano, Organometallic Vapor Epitaxy of GaN for High Brightness Blue Light Emitting Diodes S. Nakamura, Group 111-V Nitride Based Ultraviolet-Blue-Green-Yellow Light-Emitting Diodes and Laser Diodes

Volume 49 Light Emission in Silicon: from Physics to Devices D.J. Lockwood, Light Emission in Silicon G. Abstreiter, Band Gaps and Light Emission in Si/SiGe Atomic Layer Structures T. G. Brown and D. G. Hall, Radiative Isoelectronic Impurities in Silicon and Silicon-Germanium Alloys and Superlattices J. Michel, L. V. C. Assali, M. T. Morse, and L. C. Kimerling, Erbium in Silicon Y. Kanemitsu, Silicon and Germanium Nanoparticles P. M. Fauchet, Porous Silicon: Photoluminescence and Electroluminescent Devices C. Delerue, G. Allan, and M. Lannoo, Theory of Radiative and Nonradiative Processes in Silicon Nanocrystallites L. Brus, Silicon Polymers and Nanocrystals

Volume 50 Gallium Nitride (GaN) J. I. Pankove and T. D. Muustukas, Introduction S. P. DenBaars and S. Keller, Metalorganic Chemical Vapor Deposition (MOCVD) of Group 111 Nitrides W A. Bryden and T. J. Kistenmacher, Growth of Group 111-A Nitrides by Reactive Sputtering N. Newman, Thermochemistry of 111-N Semiconductors S. J. Pearton and R. J. Shul, Etching of 111 Nitrides

328

CONTENTS OF VOLUMES IN

THISSERIES

S. M. Bedair, Indium-based Nitride Compounds A. Trampert, 0. Brandt, and X H. Ploog, Crystal Structure of Group 111 Nitrides H. Morkoc, F. Hamdani, and A. Salvador, Electronic and Optical Properties of 111-V Nitride based Quantum Wells and Superlattices K. Doverspike and J. I. Pankove, Doping in the 111-Nitrides T. Suski and P. Perlin. High Pressure Studies of Defects and Impurities in Gallium Nitride B. Munemar, Optical Properties of GaN W R. L. Lambrecht, Band Structure of the Group 111 Nitrides N. E. Christensen and P. Perlin, Phonons and Phase Transitions in GaN S. Nakamura, Applications of LEDs and LDs I. Akusuki and H. Ammo, Lasers J. A. Cooper, Jr., Nonvolatile Random Access Memories in Wide Bandgap Semiconductors

Volume 5 I A

Identification of Defects in Semiconductors

C. D. Warkins, EPR and ENDOR Studies of Defects in Semiconductors J.-M. Spaefh. Magneto-Optical and Electrical Detection of Paramagnetic Resonance in Semiconductors T. A. Kennedy and E. R. Glaser, Magnetic Resonance of Epitaxial Layers Detected by Photoluminescence K. H. Chow. B. Hitti, and R. F. KieJ, pSR on Muonium in Semiconductors and Its Relation to Hydrogen K. Saarinen, P. Hautojiirvi, and C. Corbel, Positron Annihilation Spectroscopy of Defects in Semiconductors R. Jones and P. R. Brzddon, The Ah lnitio Cluster Method and the Dynamics of Defects in Semiconductors

Volume 51B Identification of' Defects in Semiconductors G. Davies, Optical Measurements of Point Defects P. M. Muuney, Defect identification Using Capacitance Spectroscopy M. Sfavola, Vibrational Spectroscopy of Light Element Impurities in Semiconductors P. Schwander, W. D. Rau, C. Kisielowski, M. Gribelyuk, and A. Ourmazd, Defect Processes in Semiconductors Studied at the Atomic Level by Transmission Electron Microscopy N. D. Jager and E. R. Weber, Scanning Tunneling Microscopy of Defects in Semiconductors

Volume 52 Sic Materials and Devices K. Jarrendahl and R. F. Davis, Materials Properties and Characterization of Sic V. A. Dmitriev and M. G. Spencer, Sic Fabrication Technology: Growth and Doping V. Saxena and A. J. Steckl, Building Blocks for Sic Devices: Ohmic Contacts, Schottky Contacts, and p-n Junctions M. S. Shur, Sic Transistors C. D. Brandt. R C. Clarke, R. R Siergiej, J. B. Casady, A. W. Morse, S. Sriram, and A. K. Agarwal, Sic for Applications in High-Power Electronics R. J. Trew, Sic Microwave Devices

CONTENTS OF VOLUMES IN THIS SERIES

329

J. Edmond, H. Kong, G. Negley, M. Leonard, K Doverspike, W Weeks, A. Suvorov, D. Waltz, and C. Carter, Jr., Sic-Based UV Photodiodes and Light-Emitting Diodes H. Morkoc, Beyond Silicon Carbide! 111-V Nitride-Based Heterostructures and Devices

Volume 53 Cumulative Subject and Author Index Including Tables of Contents for Volume 1-50

Volume 54 High Pressure in Semiconductor Physics I W. Paul, High Pressure in Semiconductor Physics: A Historical Overview N. E. Christensen, Electronic Structure Calculations for Semiconductors under Pressure R. J. Neimes and M. I. McMahon, Structural Transitions in the Group IV, 111-V and 11-VI Semiconductors Under Pressure A. R. Goni and X Syassen, Optical Properties of Semiconductors Under Pressure P. Trautman, M. Baj, and J. M. Baranowski, Hydrostatic Pressure and Uniaxial Stress in Investigations of the EL2 Defect in GaAs M. Li and P. Y. Yu,High-pressure Study of DX Centers Using Capacitance Techniques T. Suski, Spatial Correlations of Impurity Charges in Doped Semiconductors N. Kuroda, Pressure Effects on the Electronic Properties of Diluted Magnetic Semiconductors

Volume 55 High Pressure in Semiconductor Physics I1 D. K Maude and J. C. Portal, Parallel Transport in Low-Dimensional Semiconductor Structures P. C. Klipstein, Tunneling Under Pressure: High-pressure Studies of Vertical Transport in Semiconductor Heterostructures E. Anastassakis and M. Cardona, Phonons, Strains, and Pressure in Semiconductors F. H. Pollak, Effects of External Uniaxial Stress on the Optical Properties of Semiconductors and Semiconductor Microstructures A. R Adarns, M. Silver, and . I Allam, Semiconductor Optoelectronic Devices S. Porowski and I. Grzegory, The Application of High Nitrogen Pressure in the Physics and Technology of 111-N Compounds M. YousuJ Diamond Anvil Cells in High Pressure Studies of Semiconductors

Volume 56 Germanium Silicon: Physics and Materials J. C. Bean, Growth Techniques and Procedures D. E. Savage, F. Liu, V. Zielasek, and M. G. Lagally, Fundamental Crystal Growth Mechanisms R. Hull, Misfit Strain Accommodation in SiGe Heterostructures M. J. Shaw and M. Jaros, Fundamental Physics of Strained Layer GeSi: Quo Vadis? F. Cerdeira, Optical Properties S. A. Ringel and P. N. Grillot. Electronic Properties and Deep Levels in Germanium-Silicon J. C. Campbell, Optoelcctronics in Silicon and Germanium Silicon K. Eberl, K Brunner, and 0.G. Schmidt, Si, -yCy and Si, -~s-yGexCy Alloy Layers

330

CONTENTS OF VOLUMHIN THISSERIES

Volume 57 Gallium Nitride (GaN) I1 R. J. Molnar, Hydride Vapor Phase Epitaxial Growth of 111-V Nitrides T. D.Moustakas, Growth o f 111-V Nitrides by Molecular Beam Epitaxy Z. Liliental- Weber, Defects in Bulk GaN and Homoepitaxial Layers C. G. Van de Walle and N. M. Johnson, Hydrogen in 111-V Nitrides W. Gotz and N. M. Johnson, Characterization of Dopants and Deep Level Defects in Gallium Nitride B. Gil, Stress Effects on Optical Properties C. Kisieluwski, Strain in GaN Thin Films and Heterostructures J. A. Miragliotta and D.K. Wickenden, Nonlinear Optical Properties of Gallium Nitride B. K. Meyer, Magnetic Resonance Investigations on Group 111-Nitrides M. S. Shur and M. A s f K h a n , GaN and AlGaN Ultraviolet Detectors C. H. Qiu, J. I. Pankove, and C. Rossington, Ill-V Nitride-Based X-ray Detectors

Volume 58 Nonlinear Optics in Semiconductors I A. Kost, Resonant Optical Nonlinearities in Semiconductors E. Garmire, Optical Nonlinearities in Semiconductors Enhanced by Carrier Transport D. S. Chemla, Ultrafast Transient Nonlinear Optical Processes in Semiconductors M. Sheik-Bahae and E. W. Van Strylmd, Optical Nonlinearities in the Transparency Region of Bulk Semiconductors J. E. Millerd, M. Ziuri, and A. Parfovi, Photorefractivity in Semiconductors

Volume 59 Nonlinear Optics in Semiconductors I1 J. B. Khurgin, Second Order Nonlinearities and Optical Rectification K. L. Hall, E. R. Thorn. and E. P. Ippen, Nonlinearities in Active Media E. Hanamura, Optical Responses of Quantum Wires/Dots and Microcavities U. Keller, Semiconductor Nonlinearities for Solid-state Laser Modelocking and Q-Switching A. Miller, Transient Grating Studies of Carrier Diffusion and Mobility in Semiconductors

Volume 60 Self-Assembled InGaAdGaAs Quantum Dots Mitsuru Sugawara, Theoretical Bases of the Optical Properties of Semiconductor Quantum Nano-Structures Yoshiaki Nakata, Yoshihiro Sugiyama, and Mitsuru Sugawara, Molecular Beam Epitaxial Growth of Self-Assembled InAs/GaAs Quantum Dots Kohki Mukai, Mitsuru Sugawara, Mitsuru Egawu, and Nobuyuki Ohtsuka, Metalorganic Vapor Phase Epitaxial Growth of Self-Assembled InGaAs/GaAs Quantum Dots Emitting at 1.3 pm Kohki Mukai and Mitsuru Sugawara, Optical Characterization of Quantum Dots Kohki Mukai and Mitsuru Sugawara, The Photon Bottleneck Effect in Quantum Dots Hujirne Shuji, Self-Assembled Quantum Dot Lasers Hiroshi Ishikawa. Applications of Quantum Dot to Optical Devices Mitsuru Sugawara, Kohki Mukai, Hiroshi Ishikawa, Koji Otsubo, and Yoshiaki Nakata, The Latest News

CONTENTS OF VOLUMES IN THISSERIES

331

Volume 61 Hydrogen in Semiconductors I1 Norbert H. Nickel, Introduction to Hydrogen in Semiconductors I1 Noble M. Johnson and Chris G. Van de Walle, Isolated Monatomic Hydrogen in Silicon Yurij V. Gorelkinskii, Electron Paramagnetic Resonance Studies of Hydrogen and HydrogenRelated Defects in Crystalline Silicon Norbert H. Nickel, Hydrogen in Polycrystalline Silicon Wolfiard Beyer, Hydrogen Phenomena in Hydrogenated Amorphous Silicon Chris G. Vat7 de Walle, Hydrogen Interactions with Polycrystalline and Amorphous Silicon-Theory Karen M. McNamara Rutledge, Hydrogen in Polycrystalline CVD Diamond Roger L. Lichti, Dynamics of Muonium Diffusion, Site Changes and Charge-State Transitions Matshew D. McCluskey and Eugene E. Halter, Hydrogen in 111-V and 11-VI Semiconductors S. J. Pearton and J. W. Lee, The Properties of Hydrogen in GaN and Related Alloys Jorg Neugebauer and Chris G. Van de Walle, Theory of Hydrogen in GaN

Volume 62 Intersubband Transitions in Quantum Wells: Physics and Device Applications I Manjred Helm, The Basic Physics of Intersubband Transitions Jerome Faist, Carlo Sirtori, Federico Capasso, Loren N. Pfeifer, Ken W. West, Deborah L. Sivco, and Alfred Y. Cho, Quantum Interference Effects in Intersubband Transitions H. C. Liu, Quantum Well Infrared Photodetector Physics and Novel Devices S. D. Gunapala and S. I? Bandara, Quantum Well Infrared Photodetector (QWIP) Focal Plane Arrays

Volume 63 Chemical Mechanical Polishing in Si Processing Frank B. Kaufmun, Introduction Thomas Bibby and Karey Holland, Equipment John P. Bare, Facilitization Duane S. Boning and Okumu Ouma, Modeling and Simulation Shin Hwa Li, Bruce Tredinnick, and Mel Hoffman, Consumables I: Slurry Lee M. Cook, CMP Consumables 11: Pad FranGois Tardij; Post-CMP Clean Shin Hwa Li, Tara Chhatpur, and Frederic Robert, CMP Metrology Shin Hwa Li, Visun Bucha, and Kyle Wooldridge, Applications and CMP-Related Process Problems

ISBN 0-32-752373-9

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Conducting Organic Materials and Devices, Volume 81 (Semiconductors and Semimetals) (Semiconductors and Semimetals)

Semiconductors and Semimetals, Volume 21, Part A: Hydrogenated Amorphous Silicon : Preparation and Structure (Semiconductors and Semimetals)

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Semiconductors and Semimetals. Volume 27: Highly Conducting Quasi-One-Dimensional Organic Crystals (Semiconductors and Semimetals)

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Semiconductors and Semimetals, Volume 70: Recent Trends in Thermoelectric Materials Research, Part Two (Semiconductors and Semimetals)

Semiconductors and Semimetals, Volume 69: Recent Trends in Thermoelectric Materials Research, Part One (Semiconductors and Semimetals)

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Semiconductors and Semimetals. Volume 22: Lightwave Communications Technology, Part B, Semiconductor Injection Lasers, I (Semiconductors and Semimetals)

Electroluminescence I, Volume 64 (Semiconductors and Semimetals)

Electroluminescence II, Volume 65 (Semiconductors and Semimetals)

Semiconductors and Semimetals Volume 72 Silicon Epitaxy (Semiconductors and Semimetals)

Semiconductors and Semimetals Volume 72 Silicon Epitaxy (Semiconductors and Semimetals)

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Semiconductors and Semimetals. Volume 25: Diluted Magnetic Semiconductors (Semiconductors and Semimetals)

Semiconductors and Semimetals, Volume 67: Ultrafast Physical Processes in Semiconductors (Semiconductors and Semimetals)

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Semiconductors and Semimetals, Volume 14 Cadmium Telluride

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Semiconductors and Semimetals, Volume 06 Injection Phenomena

Semiconductors and Semimetals, Volume 51A Identification of Defects in Semiconductors

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Semiconductors and Semimetals, Volume 71: Recent Trends in Thermoelectric Materials Research: Part Three (Semiconductors and Semimetals)

Conducting Organic Materials and Devices, Volume 81 (Semiconductors and Semimetals) (Semiconductors and Semimetals)

Semiconductors and Semimetals, Volume 21, Part A: Hydrogenated Amorphous Silicon : Preparation and Structure (Semiconductors and Semimetals)

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Semiconductors and Semimetals, Volume 70: Recent Trends in Thermoelectric Materials Research, Part Two (Semiconductors and Semimetals)

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Semiconductors and Semimetals. Volume 22: Lightwave Communications Technology, Part B, Semiconductor Injection Lasers, I (Semiconductors and Semimetals)

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Electroluminescence I, Volume 64 (Semiconductors and Semimetals)

Electroluminescence II, Volume 65 (Semiconductors and Semimetals)

Semiconductors and Semimetals Volume 72 Silicon Epitaxy (Semiconductors and Semimetals)

Semiconductors and Semimetals Volume 72 Silicon Epitaxy (Semiconductors and Semimetals)

SEMICONDUCTORS & SEMIMETALS V38, Volume 38 (Semiconductors and Semimetals)

Semiconductors and Semimetals, Volume 67: Ultrafast Physical Processes in Semiconductors (Semiconductors and Semimetals)

Semiconductors and Semimetals. Volume 25: Diluted Magnetic Semiconductors (Semiconductors and Semimetals)

Semiconductors and Semimetals, Volume 67: Ultrafast Physical Processes in Semiconductors (Semiconductors and Semimetals)

Nanostructured Systems, Volume 35 (Semiconductors and Semimetals)

Semiconductors and Semimetals, Volume 14 Cadmium Telluride

Semiconductors and Semimetals Volume 50: Gallium Nitride (Gan) I

Semiconductors and Semimetals. Volume 11: Solar Cells

Semiconductors and Semimetals, Volume 13 Cadmium Telluride

The Spectroscopy of Semiconductors: Volume 36 (Semiconductors and Semimetals)

Semiconductors and Semimetals, Volume 06 Injection Phenomena

Semiconductors and Semimetals, Volume 51A Identification of Defects in Semiconductors

Hydrogen in Semiconductors II, Volume 61 (Semiconductors and Semimetals)

Semiconductors and Semimetals, Volume 59 Nonlinear Optics in Semiconductors II

Hydrogen in Semiconductors (Semiconductors and Semimetals, Volume 34)

Semiconductors and Semimetals, Volume 71: Recent Trends in Thermoelectric Materials Research: Part Three (Semiconductors and Semimetals)

Conducting Organic Materials and Devices, Volume 81 (Semiconductors and Semimetals) (Semiconductors and Semimetals)

Semiconductors and Semimetals, Volume 21, Part A: Hydrogenated Amorphous Silicon : Preparation and Structure (Semiconductors and Semimetals)

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