SEMICONDUCTORS AND SEMIMETALS VOLUME 17 CW Beam Processing of Silicon and Other Semiconductors
Semiconductors and Semimetals A Treatise
Edited by R . K . WILLARDSON
ALBERT C . BEER
WILLARDSON CONSULTING
BATlFLLE COLUMBUS LABORATORIES
SPOKANE, WASHINGTON
COLUMBUS, OHIO
SEMICONDUCTORS AND SEMIMETALS VOLUME 17 CW Beam Processing of Silicon and Other Semiconductors
Volume Editor JAMES F. GIBBONS DEPARTMENT OF ELECTRICAL ENGINEERING STANFORD ELECTRONICS LABORATORIES STANFORD UNIVERSITY STANFORD, CALIFORNIA
1984
ACADEMIC PRESS, INC. (Harcourt Brace Jovanovich, Publishers)
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ACADEMIC PRESS,INC. Orlando. Florida 32887
United Kin dom Editionpublished by ACADEM~CPRESS. INC. (LONDON)LTD. 24/28 Oval Road, London NWl
7DX
Library of Congress Cataloging in Publication Data Ilaln entry under t i t l e :
Semiconductors and semlmtals. Vol. 13 edited by K. Zanlo; v. 17 edited by James F. Gibbons. Includes bibllographlcal references and Indexes. 1. Sealconduetors--Collected w r k r . 2. S a l m t a l s - Collected works. I. Ylllardson, Robart I(. 11. Beer, Albert C. QCb10.9eS47 537.6'22 65-26048 ISM 0-12-752117-8 (v. 17)
PRIhTED INTHECSlTEDS~AlESOPAMERlCA
84858687
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Contents LISTOF CONTRIBUTORS . . . . . . . . . . . . . . . . . . . . . . PREFACE . . . . . . . . . . . . . . . . . . . . . . . . . .
ix xi
Chapter 1 Beam Processing of Silicon James F . Gibbons 1.1 1.2 1.3 1.4 1.5 1.6
Central Features of Pulsed Beam Annealing of Ion Implanted Silicon Central Features of CW Beam Annealing of Ion Implanted Silicon . Annealing with Large Diameter Scanning CW Sources . . . . CW Beam Recrystallization of Polysilicon Films . . . . . . Silicide Formation with Scanning Beams . . . . . . . . Rapid Thermal Processing with Stationary Incoherent Sources . . References . . . . . . . . . . . . . . . . . .
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Chapter 2 Temperature Distributions and Solid Phase Reaction Rates Produced by Scanning CW Beams Arto Lietoila. Richard B . Gold. James F. Gibbons. and Lee A . Christel 2.1 2.2 2.3 2.4 2.5 2.6
Introduction . . . . . . . . . . . . . . . . . . . The Kirchhoff Transform . . . . . . . . . . . . . . . Temperature Calculations for CW Lasers . . . . . . . . . Temperature Calculations for Wide Area Optical Sources: The Arc Lamp Temperature Calculations for Scanning CW Electron Beams . . . Reaction Rate Calculations . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . .
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71 72 74 89 93 101 105
Chapter 3 Applications of CW Beam Processing to Ion Implanted Crystalline Silicon Arto Lietoila and James F. Gibbons 3.1 Introduction .
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107
vi 3.2 3.3 3.4 3.5 3.6 3.7
CONTENTS
Experimental Methods . . . . . . . . . . . . . CW Laser Induced Crystallization of Amorphous Silicon . Annealing of “As Implanted Si: C < C.. . . . . . . Annealing of “As Implanted Si: C > C,, . . . . . . High Resolution Selective E-Beam Annealing of As Implanted Annealing of B Implanted Si . . . . . . . . . . . . . . . . . . . . . . . . References .
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Silicon
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165
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173
Chapter 4 Electronic Defects in CW Transient
Thermal Processed Silicon N . M . Johnson 4.1 4.2 4.3 4.4 4.5 4.6 4.7 4.8
Introduction . . . . . . . . . . . . . Device Processing and Measurement Techniques . CW Beam Processed Bulk Silicon . . . . . Lateral Nonuniformities in Scanned-Beam Annealing Isothermal Transient Annealing . . . . . . Post-Recrystallization Processing . . . . . Beam-Crystallized Silicon Thin Films . . . . Summary and Conclusions . . . . . . . References . . . . . . . . . . . . .
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Chapter 5 Beam Recrystallized Polycrystalline Silicon:
Properties. Applications. and Techniques K . F. Lee . T. J . Stultz. and James F . Gibbons . I Recrystallization of Thin Polycrystalline Films with a Scanning CW Laser . .2 Resistivity Reduction in Beam-Recrystallized Polysilicon Films . . . . .3 Electronic Properties of Laser Recrystallized Films Intended for MOS Device Fabrication . . . . . . . . . . . . . . . . . . . . . 5.4 Characteristics of MOS Devices and Integrated Circuits Fabricated on Laser Recrystallized Polysilicon Films . . . . . . . . . . . . . . 5.5 Improvements in the Recrystallization Process . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . .
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250
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Chapter 6 Metal-Silicon Reactions and Silicide
Formation T. Shibatu. A . Wakita. 7: W. Sigrrioii. and James F. Gibbons 6.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . 341 6 . 2 Experimental Techniques . . . . . . . . . . . . . . . . . . . . 344 6.3 Experimental Results: Laser Processing . . . . . . . . . . . . . . . . 357
vii
CONTENTS
6.4 Experimental Results: Electron Beam Processing . . . . . . . . . . . . . 371 6.5 Thermal Stability and Oxidation Properties of CW Beam Reacted Silicides . . . . . 375 6.6 Laser Processing of Nb3X Superconductors . . . . . . . . . . . . . . 385 References . . . . . . . . . . . . . . . . . . . . . . . . 393
Chapter 7 CW Beam Processing of Gallium Arsenide Yves I . Nissim and James F . Gibbons 7.1 7.2 7.3 7.4 7.5
Introduction . . . . . . . . . . . . . . . CW Laser Annealing of Ion Implanted GaAs . . . . CW Laser Processing of Thin Films Deposited on GaAs . Rapid Thermal Processing . . . . . . . . . Novel Applications of Rapid Thermal Processing in GaAs References . . . . . . . . . . . . . . .
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INDEX . . . . . . . . . . . . . . . . . . . . . CONTENTS OF PREVIOUS VOLUMES . . . . . . . . . . . . .
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397
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List of Contributors Numbers in parentheses indicate the pages on which the authors' contributions begin.
LEEA. CHRISTEL',Stanford Electronics Laboratories, Stanford University, Stanford, California 94305 (71) JAMES F. GIBBONS, Department of Electrical Engineering, Stanford Electronics Laboratories, Stanford University, Stanford, California 94305 (1, 71, 107, 227, 341, 397) RICHARDB . GOLD', Stanford Electronics Laboratories, Stanford University, Stanford, California 94305 (71) N. M. JOHNSON,Xerox Corporation, Palo Alto Research Center, Palo Alto, California 94304 (177) K . F. LEE3, Stanford Electronics Laboratories, Stanford University, Stanford, California 94305 (227) ARTOLIETOILA4, Stanford Electronics Laboratories, Stanford University, Stanford, California 94305 (71, 107) YvES I. NISSIM', Stanford Electronics Laboratories, Stanford University, Stanford, California 94305 (397) T. SHIBATA', Stanford Electronics Laboratories, Stanford University, Stanford, California 94305 (341) T. W. SIGMON, Stanford Electronics Laboratories, Stanford University, Stanford, California 94305 (341) T. J. STULTZ', Stanford Electronics Laboratories, Stanford University, Stanford, California 94305 (227) A. WAKITA,Stanford Electronics Laboratories, Stanford University, Stanford, California 94305 (341) 'Present address: SERA Solar Corporation, Santa Clara, California 95054. 'Present address: Adams-Russell Company, Burlington, Massachusetts 01803. 3Present address: AT&T Bell Laboratories, Holmdel, New Jersey 07733. 4Present address: Micronas, Inc., 00101 Helsinki, Finland. *Present address: CNET, 92220 Bagneux, France. 6Present address: VLSI Research Center, Toshiba Research and Development Center, Kawasaki City, Kanagawa, 201 Japan. 'Present address: TS Associates, San Jose, California 95128.
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Preface This book had its beginnings at the 1979 winter meeting of the American Physical Society, at which I was asked to review the basic features of cw laser processing of ion implanted silicon. The review covered techniques for calculating laser-induced temperatures, experimental results on the annealing of ion implanted layers in silicon, and enough material on structural (TEM) and electronic (DLTS) defects to suggest that the process was reasonably well behaved and perhaps had a future in materials preparation and processing. A1 Beer, who was attending the meeting, thought the subject might be suitable for a short research monograph that would be useful to those entering the field. We agreed that, as things seemed reasonably tidy, the project could easily be finished within a year. However, the field was just then entering a period of very rapid growth. The fabrication of MOSFETS on polysilicon films recrystallized by a cw scanning laser, reported in July 1979, opened up the field of silicon-on-insulators (SOI). There followed a succession of significant contributions to this technology at a number of laboratories, dealing with (1) such things as thermal profile management, beam shaping, seeding, and zone recrystallization on the materials side and (2) a number of new device configurations, including the general technique of three-dimensional integration, on the applications side. Simultaneously, the use of a cw scanning laser beam to react metals with silicon to form single-phase silicide layers was being explored; and its potential consequences for materials, device, and integrated circuit technology required that it be included if the book was to offer a broad, balanced view of the field. Finally, as these areas began to settle down enough to be reviewed with some reasonable perspective, the rapid thermal processing technique was reported in Japan. Its natural relationship to the cw scanning laser process and its “throughput” advantage from a manufacturing point of view then required a further modification of many chapters of the book. The result is that 5 years have passed since A1 Beer and I shook hands on the project. The cw beam technology as it relates to semiconductor processing is now much more mature and better understood than it was in 1979-1980, though its applications in the semiconductor industry are very likely still in their infancy. It is on this hopeful note that my coauthors and I now offer this volume. Chapter 1 provides an overview of the general field of beam processing, with a brief discussion of the main topics to be developed in depth in the succeeding chapters. Chapter 2 deals with the calculation of cw beam-induced temperature
xi
xii
PREFACE
profiles in semiconductor substrates under a variety of conditions that are found in practical applications. A review of the use of cw beam processing for annealing ion implanted silicon is presented in Chapter 3, and the question of electronic defects remaining after various implantation and cw beam annealing cycles is reviewed in Chapter 4. In Chapter 5 we discuss the recrystallization of thin polysilicon films by cw beam systems and summarize its applications in device fabrication. The use of a cw beam to react a thin metal film with a silicon (crystalline or polycrystalline) substrate to form a silicide is reviewed in Chapter 6 , together with oxidation characteristics and other properties of the silicide layers that are important for practical applications. Finally, in Chapter 7 we review the current state of the art regarding the use of cw beams for processing of GaAs. It will be apparent from this chapter that much remains to be done in the beam processing of compound semiconductors. To conclude this preface, I would like to express my most sincere appreciation to the authors of these chapters. We were all privileged to share the experience and excitement of working in a field during the early stages of its development, both as colleagues at Stanford and as members of the larger community of “beam annealers” whose pioneering work is described here in terms that we hope will reflect the respect we have for all of them. My coauthors and I would also like to acknowledge our enormous indebtedness to Ms. Mary Cloutier, both for her extraordinary skill in preparing this manuscript and for the patience and unfailing good humor that she maintained through its several revisions. Finally, 1 want to thank Dick Reynolds and Sven Roosild at DARPA, Ben Wilcox at NSF, and Horst Wittman at ARO for the support that enabled the Stanford group to participate in the development of the field.
JAMES F. GIBBONS
CHAPTER 1
Beam Processing of Silicon James F. Gibbons STANFORD ELECTRONICS LABORATORIES STANFORD UNIVERSITY STANFORD, CALIFORNIA
1.1 CENTRAL FEATURES OF PULSED BEAM
ANNEALINGOF ION IMPLANTED SILICON . . . . . . . . . . . . . . . 1.2 CENTRAL FEATURES OF C W BEAMANNEALING OF ION IMPLANTED SILICON
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1.3 ANNEALING WITH LARGEDIAMETER SCANNING Cw SOURCES ............................. 1.4 CW BEAM RECRYSTALLIZATION OF POLYSILICON FILMS. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.5 SILICIDE FORMATION WITH SCANNING BEAMS . . . . . . . . . . . 1.6 RAPID THERMAL PROCESSING WITH STATIONARY INCOHERENT SOURCES
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REFERENCES. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
2 30 40 46 51
55 66
High power lasers have been used i n t h e semiconductor indust r y f o r many y e a r s , f o r a p p l i c a t i o n s such as wafer s c r i b i n g , r e s i s t o r trimming and t h e d r i l l i n g of ceramics, where t h e c l e a n removal of a s p e c i f i c amount of material i s r e q u i r e d ; and f o r c o n t a c t a l l o y i n g , where permanent, - r e l a t i v e l y e x t e n s i v e chemical D e s p i t e t h e mature development of t h e s e changes are sought [1.1]. a p p l i c a t i o n s , however, lasers and e l e c t r o n beams have only rec e n t l y been considered f o r t h e p r o c e s s i n g o p e r a t i o n s t h a t are required during s i l i c o n device f a b r i c a t i o n . The e a r l i e s t work w i t h Q-switched lasers w a s performed by Russian workers [1.2-1.41, who showed t h a t i o n implanted s i l i c o n could be annealed s u c c e s s f u l l y w i t h laser p u l s e s of s u f f i c i e n t energy and d u r a t i o n . The importance of t h i s work was q u i c k l y a p p r e c i a t e d by workers a t s e v e r a l d i f f e r e n t l a b o r a t o r i e s i n Europe [1.5-1.71 and t h e United S t a t e s 11.8-1.101, and a large volume of published work has now appeared on t h e use of both Qswitched lasers and pulsed e l e c t r o n beams f1.111 f o r semiconductor A t t h e same t i m e , t h e use of scanning cw laser and processing. e l e c t r o n beams f o r a n n e a l i n g i o n implanted s i l i c o n was being e x p l o r e d a t S t a n f o r d [1.12-1.141, w i t h a p a r t i c u l a r f o c u s on i t s a p p l i c a t i o n t o semiconductor device p r o c e s s i n g . More r e c e n t l y , both scanning and s t a t i o n a r y i n c o h e r e n t l i g h t s o u r c e s have been used t o a c h i e v e a n n e a l i n g of i o n implanted s i l i c o n w i t h a wafer throughput t h a t i s compatible w i t h t h e needs of t h e semiconductor industry. SEMICONDUCTORS AND SEMIMETALS, VOL 17
1
Copynght 0 1984 by Academic Press, Inc All nghts of reproductlon m any form reserved ISBN 0-12-752117-8
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J . F. GIBBONS
A s a r e s u l t of t h e s e e f f o r t s , beam a n n e a l i n g of i o n implanted s i l i c o n has now been r a t h e r thoroughly s t u d i e d and shown t o possess s e v e r a l p o t e n t i a l advantages over t h e f u r n a c e a n n e a l i n g process. I n a d d i t i o n , l a s e r s , e l e c t r o n beams and a r c sources have been used t o r e c r y s t a l l i z e vapor d e p o s i t e d p o l y s i l i c o n , w i t h s u b s t a n t i a l improvements i n i t s e l e c t r o n i c p r o p e r t i e s [1.15,1.16] ; t o f a c i l i t a t e t h e formation of metal s i l i c i d e s [1.17-1.201; and t o perform a number of o t h e r p r o c e s s i n g f u n c t i o n s t h a t are of i n c r e a s i n g importance i n t h e f a b r i c a t i o n of f i n e geometry i n t e g r a t e d c i r c u i t s and h i g h speed d e v i c e s . I n t h i s c h a p t e r w e w i l l review t h e b a s i c mechanisms by which beam a n n e a l i n g proceeds and summari z e t h e b a s i c m e t a l l u r g i c a l and e l e c t r o n i c p r o p e r t i e s of beamannealed material.
Two d i s t i n c t l y d i f f e r e n t beam a n n e a l i n g mechanisms have been i d e n t i f i e d , depending on t h e d u r a t i o n of t h e beam exposure. For Q-switched l a s e r s o r pulsed e l e c t r o n b e a m s , exposure times a r e t y p i c a l l y i n t h e range of 5 ns - 500 ns and t h e a n n e a l i n g p r o c e s s then i n v o l v e s t h e formation of a t h i n s u r f a c e l a y e r of molten s i l i c o n t h a t r e c r y s t a l l i z e s on t h e u n d e r l y i n g s u b s t r a t e when t h e r a d i a t i o n i s removed. If t h e i r r a d i a t e d sample i s an i o n implant e d s i n g l e c r y s t a l and t h e depth of t h e molten l a y e r i s s u f f i c i e n t t o envelop t h e i m p l a n t a t i o n damaged region, t h e molten l a y e r regrows by a l i q u i d phase e p i t a x i a l process on t h e c r y s t a l l i n e s u b s t r a t e t o produce material w i t h a very high degree of s t r u c t u r a l p e r f e c t i o n and very s u p e r i o r e l e c t r o n i c p r o p e r t i e s . For cw systems, on t h e o t h e r hand, t h e s i l i c o n s u r f a c e i s t y p i c a l l y exposed t o t h e beam f o r 0.1-10 ms and, i n some c a s e s , f o r d u r a t i o n s of s e v e r a l t e n s of seconds. The a n n e a l i n g of i o n implanted m a t e r i a l can then proceed by a s o l i d phase e p i t a x i a l regrowth process a t temperatures t h a t a r e w e l l below t h e m e l t i n g p o i n t . A s i n t h e pulsed beam case, a very h i g h degree of cryst a l l i n e p e r f e c t i o n and very s u p e r i o r e l e c t r o n i c p r o p e r t i e s can be obtained under a p p r o p r i a t e a n n e a l i n g c o n d i t i o n s . However, t h e absence of m e l t i n g proves t o be of some i n t e r e s t s i n c e no r e d i s t r i b u t i o n of t h e implanted impurity p r o f i l e then occurs during t h e a n n e a l i n g p r o c e s s , whereas s i g n i f i c a n t impurity r e d i s t r i b u t i o n occurs when a n n e a l i n g i s e f f e c t e d by a p u l s e d laser o r e l e c t r o n beam. These and o t h e r d i f f e r e n c e s i n t h e two a n n e a l i n g In processes make i t convenient t o d i s c u s s them s e p a r a t e l y . what f o l l o w s w e f i r s t c o n s i d e r t h e pulsed beam a n n e a l i n g p r o c e s s and then t a k e up t h e cw a l t e r n a t i v e .
1.1
CENTRAL SILICON
FEATURES OF PULSED
BEAM ANNEALING OF ION IMPLANTED
W e w i l l begin w i t h a b r i e f review of experimental d a t a t h a t c h a r a c t e r i z e t h e a n n e a l i n g of i o n implanted s i l i c o n u s i n g a Qswitched l a s e r and then t u r n t o a b r i e f mathematical a n a l y s i s of the process.
1. BEAM PROCESSING OF SILICON
3
Laser a n n e a l i n g has been demonstrated u s i n g a v a r i e t y of s o u r c e s and p u l s e d u r a t i o n s . I n most of t h e experiments r e p o r t e d s o f a r , implanted samples have been annealed d i r e c t l y i n t h e l a b o r a t o r y ambient; i.e., no s p e c i a l p r e c a u t i o n s have been taken t o immerse t h e wafer i n an i n e r t environment d u r i n g annealing. The area i r r a d i a t e d by t h e beam i s , i n most cases, l a r g e compared t o b o t h t h e d i f f u s i o n l e n g t h f o r h e a t i n t h e s o l i d and t h e t h i c k n e s s of t h e w a f e r , s o t h e p r o c e s s can be t r e a t e d as being b a s i c a l l y one dimensional.
Experimental r e s u l t s from v a r i o u s l a b o r a t o r i e s are reasonably c o n s i s t e n t and are summarized i n Table 1.1. For convenience, t h e d a t a are s e l e c t e d t o i l l u s t r a t e t h e a n n e a l i n g of As+-implanted s i l i c o n w i t h t h e As+ i m p l a n t a t i o n c o n d i t i o n s chosen t o provide amorphous r e g i o n s of d i f f e r i n g t h i c k n e s s e s a t t h e sample s u r f a c e . For The p r i n c i p a l f e a t u r e s of t h e d a t a are as follows. p u l s e s of 25-100 ns d u r a t i o n ( t y p i c a l of a Q-switched l a s e r ) :
1. Surface m e l t i n g is i n i t i a t e d at a t h r e s h o l d energy of 0.2-3.5 J/cm2, depending on t h e laser wavelength and t h e t h i c k n e s s of t h e amorphous l a y e r . 2. F u l l a n n e a l i n g of a n i o n implanted l a y e r ( a s judged by c r y s t a l recovery and e l e c t r i c a l a c t i v i t y ) r e q u i r e s a d d i t i o n a l energy i n an amount t h a t a l s o depends on both t h e laser wavelength and t h e t h i c k n e s s of t h e amorphous l a y e r . TABLE 1.1. of As+-Implanted
Laser
Experimental R e s u l t s f o r Pulsed Laser Annealing Silicon AS+
d'
Melt Initiated
Imp. P a r . Ruby 3 50 n s
5x1Ol5, 400 keV
43008
Ruby 9 60 n s
1.4x1Ol6, 100
17008
Fully Recovered
2 J/cm2 0.64
Ref.
[5]
1.4
Nd :YAG ,
110 ns
8x1Ol5, 100 keV
40 ns
130 keV
468
3.5
4.5
[211
30 keV
468
0.2
0.8
[211
Nd:YAG, Doubled 40 n s
-15008
6 J/cm2
[7]
4
J. F. GIBBONS
E a r l y evidence s u g g e s t i n g t h e s u r f a c e m e l t i n g t h e o r y of pulsed laser a n n e a l i n g w a s provided by Auston e t a l . [1.21], who measured t h e o p t i c a l r e f l e c t i v i t y of samples d u r i n g t h e a n n e a l i n g c y c l e w i t h a low power (2mW) continuous He-Ne laser. F i g u r e 1.1 shows a t y p i c a l t i m e r e s o l v e d s u r f a c e r e f l e c t i v i t y measurement . with a when a Nd:YAG l a s e r o p e r a t e d a t a wavelength of 0.53 m p u l s e w i d t h of 30 n s and a t o t a l energy d e n s i t y energy of 2.75 J/cm2 i s used t o i r r a d i a t e a
s i l i c o n sample implanted w i t h As' a t 30 keV t o a dose of 1015/cm2. A s t h e a n n e a l i n g l a s e r t u r n s on, t h e i n i t i a l r e f l e c t i v i t y of t h e amorphous l a y e r , Ra, rises w i t h i n c r e a s i n g sample t e m p e r a t u r e t o a v a l u e It;, where s u r f a c e The r e f l e c t i v i t y t h e n rises a b r u p t l y t o melting i s i n i t i a t e d . t h e r e f l e c t i v i t y of molten s i l i c o n and remains a t t h i s v a l u e a s t h e l i q u i d - s o l i d i n t e r f a c e p e n e t r a t e s more deeply i n t o t h e sample. The p e n e t r a t i o n ceases a t a t i m e a f t e r t h e a n n e a l i n g laser p u l s e has t e r m i n a t e d t h a t depends on t h e t o t a l energy absorbed d u r i n g the pulse.
A s t h e l i q u i d - s o l i d i n t e r f a c e r e v e r s e s i t s d i r e c t i o n and proceeds back toward t h e s u r f a c e , t h e molten zone regrows on t h e u n d e r l y i n g s i l i c o n c r y s t a l . The r e f l e c t i v i t y u l t i m a t e l y drops t o RG, t h e r e f l e c t i v i t y of t h e h o t s o l i d . The s u r f a c e is m e l t e d The f a l l t i m e , T f , i s t h e time r e q u i r e d f o r a t o t a l time 5. f o r t h e regrowth f r o n t t o move through one o p t i c a l s k i n d e p t h and i s a measure of t h e regrowth v e l o c i t y . The r e f l e c t i v i t y d e c r e a s e s t o i t s room t e m p e r a t u r e v a l u e f o r c r y s t a l l i n e s i l i c o n , Rc, as t h e s u r f a c e cools.
100 nrldiv. FIGURE 1.1. Ref. 1.21).
T i m e r e s o l v e d s u r f a c e r e f l e c t i v i t y trace ( a f t e r
1. BEAM PROCESSING OF SILICON 1.1.1
5
Qualitative Analysis of the Melt Threshold
The initial interaction of a light beam with a material is always with its electrons. Energy absorbed by the electrons is ultimately shared with the atoms of the material as the electronic excitation is transformed into heat. Under most conditions the transfer of electronic excitation to heat is accomplished with relaxation times on the order of 1 ps or less and may therefore be considered to be instantaneous. The basic process of pulsed beam annealing then consists of the absorption of sufficient energy to melt a layer of silicon having a thickness at least equal to the thickness of the implantation-damaged region, followed by liquid phase epitaxial regrowth of the melted layer. The laser energy required to bring the surface of the sample to its melting temperature is therefore a critical parameter in the process. To estimate the threshold energy required to produce surface melting, we consider the sample geometry shown in Fig. 1.2 [1.22]. The sample consists of an implantation-damaged layer of thickness xd and optical absorption coefficient ad resting on a crystalline solid with optical absorption coefficient aC. A laser pulse of intensity 1, and duration T~ illuminates the front surface of the sample at normal incidence. The reflection coefficient at this surface is R. The laser intensity is assumed to be sufficiently low that ordinary one-photon optical absorption processes are dominant.
LASERPULSE ABSORBING MEDIUM
Xd
FIGURE 1.2. Intensity and temperature profiles in laserirradiate silicon. (a) Penetration depth of light small compared (b) Penetration depth of light large compared to (DT )If2. to (DT;)~’~ (after Ref. 1.22).
J . F. GIBBONS
6
The r e l e v a n t thermal p r o p e r t i e s of t h e material are i t s s p e c i f i c h e a t C v , i t s mass d e n s i t y p and i t s t h e r m a l c o n d u c t i v i t y K. The h e a t d i f f u s i v i t y i s t h e n g i v e by D = K/pCv and d e t e r T h i s l e n g t h g i v e s an mines a c h a r a c t e r i s t i c l e n g t h (DT ) P e s t i m a t e of t h e d i s t a n c e o v e r which t h e t e m p e r a t u r e p r o f i l e i s s p r e a d by h e a t d i f f u s i o n d u r i n g t h e laser p u l s e .
”.
To estimate t h e s u r f a c e t e m p e r a t u r e r i s e produced i n a l i m i t i n g c a s e of p r a c t i c a l i n t e r e s t , we suppose t h a t t h e laser e n e r g y i s absorbed i n a the diffusion length f o r h e a t [a-l < Under t h e s e c o n d i t i o n s [ i l l u s t r a t e d i n Fig. 1.2(a)] can be c o n s i d e r e d t o h e a t a s l a b of t h i c k n e s s ( D T ~ ) The c o r r e s p o n d i n g temperat u r e r i s e AT i s t h e n o b t a i n e d from
.
( I o ~ p ) ( l - R ) = PC,AT(DT~)1 / 2
(1.1)
T h i s formula s u g g e s t s t h a t t h e s u r f a c e t e m p e r a t u r e i s p r o p o r t i o n a l t o t h e energy i n t h e laser p u l s e (Io T ~ and ) predicts a t h r e s h o l d e n e r g y of
a-1 6 ( D T ~ ) ~ / Z
(1.2)
where Tm i s t h e m e l t i n g t e m p e r a t u r e and To i s t h e ambient temperature. The extreme t e m p e r a t u r e change e n v i s a g e d (- 14OOOC) of c o u r s e makes i t n e c e s s a r y t o a c c o u n t f o r t h e t e m p e r a t u r e dependence of t h e t h e r m a l and o p t i c a l p a r a m e t e r s a p p e a r i n g i n E q . (1.2) i n o r d e r t o a r r i v e a t a meaningful estimate of E& f o r p r a c t i c a l u s e . The p a r a m e t e r s t h a t a r e c u s t o m a r i l y used f o r p r o c e s s a n a l y s i s are g i v e n i n Table 1.2. The t h e r m a l c o n d u c t i v i t y i s an e s p e c i a l l y s e n s i t i v e f u n c t i o n of t e m p e r a t u r e and i t s t e m p e r a t u r e dependence must be t a k e n i n t o account t o o b t a i n a c c u r a t e t h e o r e t i c a l p r e d i c t i o n s . For n u m e r i c a l e s t i m a t e s w e u s e t h e a v e r a g e v a l u e of K c a l c u l a t e d from T
(1.3a) LO
w i t h K(T) f o r s i l i c o n g i v e n by
(W/ cm°K)
(1.3b)
1.
BEAM PROCESSING OF SILICON
7
TABLE 1.2. Thermal and Optical Parameters for Solid Silicon. The optical absorption coefficients for implantation-amorphized silicon have not been measured; values quoted for ad are for sputtered silicon. Parameter Cv, J/gm-OK P , gm/cm3 K, W/cm-OK R ad ( h = 0.69 ad ( h = 1.06 ac (1 = 0.69 ac ( h = 1.06
Value at T = 300°K
pm) pm) pm) pm)
Value at T = 1685'K
0.65 2.33 1.45 0.35 3x104 3-6~10~ 3x10 10
Ref.
0.95 2.33 0.25 0.7
Substitution of numerical values from Table 1.2 leads to values of E& that are in satisfactory agreement with experiment for short wavelength irradiation such as that from ruby or frequency-doubled Nd:YAG lasers, where the basic assumption of heat transport made in t e analysis is at least approximately valid The validity of this assumption is [i.e., a-1 < (DT )lh]. improved if an Bmplantation-amorphized sample is irradiated, of course, since the optical absorption coefficient is then increased. For Q-switched Nd:YAG lasers operated at the fundamental frequency, however, the optical absorption depth (a-l) in crystalline silicon is substantially greater than the diffusion length for heatZ An order of magnitude estimate of the melt threshold energy Em can be made for this case by assuming that the thickness lab that is heated by the radiation is (a-l). Replacing a-1 in Eq. (1.2), we obtain
with
a-1
<
(DT~)~/~
In practical circumstances, however, the melt threshold predicted by Eq. 1.4 proves to be too large for a variety of reasons. For the particular case of ion implanted silicon with damage depths of 0.1 - 0.3 pm, free carrier absorption in the underlying crystalline material and bandgap narrowing during the heatting cycle have a profound effect on the surface temperature. The possible effect of a dense free carrier plasma is also potentially important and needs to be included in the analysis. In the next section we describe a model that has been developed to calculate surface temperatures when these and other effects are included.
J. F. GIBBONS
8 1.1.2
Q u a n t i t a t i v e A n a l y s i s of t h e M e l t Threshold Energy f o r Ion-Implanted C r y s t a l l i n e S i
In t h i s s e c t i o n w e p r e s e n t t h e r e s u l t s of computer modeling of t h e t e m p e r a t u r e and c a r r i e r c o n c e n t r a t i o n v e r s u s t i m e i n p u l s e d l a s e r i r r a d i a t e d S i b e f o r e m e l t i n g o c c u r s [1.28,1.29]. The r e s u l t s a r e o b t a i n e d by s i m u l t a n e o u s l y s o l v i n g t h e d i f f e r e n t i a l e q u a t i o n s f o r t e m p e r a t u r e , c a r r i e r c o n c e n t r a t i o n ( i n c l u d i n g d i f f u s i o n ) and laser i n t e n s i t y w i t h i n t h e sample. Free carrier absorption, Auger r e c o m b i n a t i o n and t h e t e m p e r a t u r e dependence of t h e t h e r m a l conductivity a r e included. I n a d d i t i o n , t h e model a c c o u n t s f o r t h e t e m p e r a t u r e dependence of t h e bandgap and l a t t i c e a b s o r p t i o n c o e f f i c i e n t ; t h e e l e c t r o n i c c o n t r i b u t i o n t o t h e t h e r m a l conduct i v i t y ( t h i s t u r n s o u t t o be i n s i g n i f i c a n t ) ; t h e f a c t t h a t when a photon creates an e l e c t r o n - h o l d p a i r , t h e photon energy i s n o t immediately t r a n s f e r r e d t o t h e l a t t i c e ; and e l e c t r o n - h o l e s c a t t e r i n g when c a l c u l a t i n g t h e ambipolar d i f f u s i v i t y of c a r r i e r s . In a d d i t i o n , a c c u r a t e v a l u e s f o r t h e r e f l e c t i v i t y and a b s o r p t i o n c o e f f i c i e n t of t h e amorphous s i l i c o n a r e employed. With r e g a r d t o t h e k i n e t i c s f o l l o w i n g assumptions are made:
of
the
energy coupling,
the
1. E l e c t r o n and h o l e c o n c e n t r a t i o n s a r e e q u a l ; c a r r i e r d i f f u s i o n i s c h a r a c t e r i z e d by t h e ambipolar d i f f u s i o n c o e f f i c i e n t . 2 . The dominant recombination mechanism i s Auger recombinat i o n , where t h e gap energy i s t r a n s f e r r e d t o an a l r e a d y e x i s t i n g f r e e c a r r i e r ; t h e r e v e r s e p r o c e s s , impact i o n i z a t i o n , i s not included.
3 . The r e l e a s e of e n e r g y from t h e h o t c a r r i e r s t o t h e l a t t i c e i s c h a r a c t e r i z e d by an energy r e l a x a t i o n t i m e , T ~ . 4 . The d i a m e t e r of t h e l a s e r s p o t i s l a r g e compared t o t h e t h i c k n e s s of t h e amorphous l a y e r ; i . e . , t h e problem i s one dimensional. Using t h e s e assumptions t h e d i f f e r e n t i a l e q u a t i o n s f o r t e m p e r a t u r e T , c a r r i e r c o n c e n t r a t i o n p, and l a s e r i n t e n s i t y I , a r e f o r m u l a t e d and s o l v e d f o r t h e s u r f a c e t e m p e r a t u r e as a f u n c t i o n of t i m e p r i o r t o m e l t i n g . The r e s u l t s of t h e s e c a l c u l a t i o n s can be summarized as f o l l o w s .
1.1.2.1
Er,
Comparison of C a l c u l a t e d Melt Threshold w i t h E x p e r i m e n t a l Data
Auston e t a l . L1.21) have determined t h e t h r e s h o l d e n e r g y , r e q u i r e d t o reach an enhanced r e f l e c t i v i t y of -70% w i t h
1.
BEAM PROCESSING OF SILICON
9
a 30-40 ns p u l s e a t .53 pm and 1.06 um; both c r y s t a l l i n e S i and samples w i t h a 468 amorphous s u r f a c e l a y e r were s t u d i e d .
a
The computer model was used t o s t u d y t h e s e a n n e a l i n g cases. A comparison between e x p e r i m e n t a l and c a l c u l a t e d t h r e s h o l d energ i e s i s given i n Table 1.3. A s can be seen, t h e agreement i s r e a s o n a b l e , and t h e shape of t h e p u l s e does n o t appear t o be a c r u c i a l p a r a m e t e r . Note a l s o t h a t t h e melt t h r e s h o l d s f o r cryst a l l i n e S i and i m p l a n t a t i o n amorphized s i l i c o n are very s i m i l a r a t a laser wavelength of 1.06 pm. The g r e a t e s t d i f f e r e n c e between c a l c u l a t i o n and e x p e r i m e n t a l d a t a i s a t a laser wavelength The d i f f e r e n c e may be due t o inaccuracy i n t h e of .53 pm. formula used t o c a l c u l a t e o p t i c a l a b s o r p t i o n c o e f f i c i e n t a t a laser wavelength X =.53 pm a t h i g h temperatures (see R e f . 1 . 2 8 ) . TABLE 1.3. A Comparison Between Experimentally Determined Thresholds, Er, f o r Achieving t h e Enhanced R e f l e c t i v i t y (Ref. 1.21), and C a l c u l a t e d E n e r g i e s , Em, Required t o Reach S i l i c o n M e l t i n g Temperature f o r Nanosecond P u l s e s .
Em, Calc. h (pm)
Amorphous Layer (nm) ~~
1.06 1.06 0.53 0.53
1.1.2.2
Pulse Length (ns)
E r , EF. (J/cm )
Constant Intensity
Gauss i a n Pulse
40 30 30
5.5 3 .7 0.35 0 .2
3.8 3 .3 0.47 0.27
4.8 4.1 0.56 0.35
Absorption
(FCA)
~~~
0 46.8 0 46 .8
Free
(J/cm 2)
40
Carrier
and
Bandgap Narrowing
The importance of f r e e carrier a b s o r p t i o n and t h e v a r i a t i o n of t h e semiconductor bandgap w i t h temperature i n determining t h e m e l t t h r e s h o l d are shown i n F i g . 1.3, which i l l u s t r a t e s t h e s u r f a c e temperature v s t i m e of a c r y s t a l l i n e S i sample being i r r a d i a t e d w i t h a c o n s t a n t i n t e n s i t y of 160 MW/cm2 a t X = 1.06pm. The f i g u r e shows t h a t t h e p r e d i c t e d Em i s doubled i f t h e t e m p e r a t u r e dependence of Eg i s o m i t t e d . I f f r e e carier a b s o r p t i o n i s o m i t t e d , t h e n t h e behavior changes d r a s t i c a l l y : a f t e r 70 ns of r a d i a t i o n , which corresponds t o a t o t a l energy d e n s i t y of 11.2 J/cm2, t h e s u r f a c e temperature has r i s e n only t o about 70'C. The e v o l u t i o n of t h e l a t t i c e and f r e e carrier a b s o r p t i o n c o e f f i c i e n t s ( f o r 1.06 pm, 160 mW/cm2) is shown i n F i g . 1.4. Using t h e r e s u l t s p r e s e n t e d i n F i g . 1.4, t h e temperature behavior shown i n F i g . 1.3 can now be e x p l a i n e d as f o l l o w s . A t t h e o u t s e t
J. F. GIBBONS
10
ENERGY ( J / c m 2 )
4
2
1: 1 . 0 k p m 1400 -
6
P:'160 MW/:m2 Em- 3 . 8 J
i
10
8
CR&ALLINE'Si
Em=7.9J
-
t
-
w 1000Lz
b E o = CONSTANT
W
a
600-
-
W t-
400-
-
5
30 T I ME
FIGURE crystalline pendence of lines); X =
40 (ns)
5b
so
'70
1.3. The s u r f a c e t e m p e r a t u r e v s . p u l s e d u r a t i o n i n S i , when f r e e c a r r i e r a b s o r p t i o n and t e m p e r a t u r e debandgap a r e i n c l u d e d ( s o l i d l i n e ) o r i g n o r e d (dashed 1.06 Urn and P = 160 mW/cm2.
of t h e p u l s e , l a t t i c e a b s o r p t i o n (- 10 cm-l) i s s u f f i c i e n t t o c r e a t e some f r e e c a r r i e r s . Those c a r r i e r s t h e n a b s o r b r a d i a t i o n T h i s c a u s e s t h e bandgap t o s h r i n k which i s c o n v e r t e d t o h e a t . and l a t t i c e a b s o r p t i o n t o i n c r e a s e . Thus more c a r r i e r s a r e g e n e r a t e d , which means s t r o n g e r f r e e c a r r i e r a b s o r p t i o n , and s o on. The p r o c e s s i s t h u s s e l f - r e g e n e r a t i v e , so t h a t a f t e r about 20 n s t h e t e m p e r a t u r e rises r a p i d l y t o t h e m e l t i n g p o i n t . The two a b s o r p t i o n c o e f f i c i e n t s are of t h e same o r d e r of magnitude most of t h e t i m e , but t h e s e l f - r e g e n e r a t i v e p r o c e s s could n e v e r s t a r t without f r e e carrier absorption. I n c o n t r a s t t o t h e b e h a v i o r a t 1.06 urn, t h e f r e e c a r r i e r a b s o r p t i o n p r o c e s s i s found t o be c o m p l e t e l y i n s i g n i f i c a n t a t a l a s e r wavelength of - 5 3 um. Also, t h e e f f e c t of bandgap narrowing i s much less profound, as i s i l l u s t r a t e d i n F i g . 1.5. N o t i c e t h a t f o r t h i s f i g u r e a h i g h e r c o n s t a n t i n t e n s i t y ( 5 0 mW/cm2) h a s been used t h a n t h a t used f o r com a r i s o n w i t h t h e e x p e r i m e n t a l d a t a . The r e a s o n i s t h a t 50 MW/cmLP i s more t y p i c a l of t h e range used f o r p r a c t i c a l annealing. The consequence of t h i s i n c r e a s e i n t h e power d e n s i t y i s a s l i g h t d e c r e a s e i n t h e m e l t t h r e s h o l d . F i g u r e 1.5 a l s o shows t h e r a t h e r minute e f f e c t of t h e e l e c t r o n i c heat conduction, K ~ .
1.
BEAM PROCESSING OF SILICON
I
I
1
1
11
A=1.06pm P=160MW/cm2 CRYSTALLINE Si
, 5
10
15
20
5
T I M E (ns)
FIGURE 1.4. L a t t i c e and f r e e c a r r i e r a b s o r p t i o n c o e f f i c i e n t s vs. t i m e under same c o n d i t i o n s a s i n Fig. 1.3 (temperature ' dependence of bandgap i s i n c l u d e d ) . For t h e .53 pm r a d i a t i o n , t h e temperature rise w i t h t i m e does The not e x h i b i t t h e e x p l o s i v e c h a r a c t e r observed a t 1.06 pm. reason i s t h a t only l a t t i c e a b s o r p t i o n i s important, and i t i s a l r e a d y s t r o n g at room temperature. The time d e r i v a t i v e of t h e s u r f a c e temperature i n c r e a s e s , however, during t h e p u l s e , s i n c e l a t t i c e a b s o r p t i o n i n c r e a s e s w i t h temperature. 1.1.2.3
Energy Relaxation T i m e
A very imaginative proposal r e g a r d i n g t h e mechanism of pulsed l a s e r a n n e a l i n g h a s been made by Van Vechten i1.301, who argues t h a t i n t e n s e laser i r r a d i a t i o n may produce such l a r g e carrier c o n c e n t r a t i o n s t h a t c a r r i e r s c r e e n i n g e f f e c t s would i n c r e a s e t h e I f t h i s were t h e c a s e , energy r e l a x a t i o n t i m e s u b s t a n t i a l l y . electron-hole plasma e f f e c t s r a t h e r than s i m p l e s u r f a c e m e l t i n g would be r e q u i r e d t o e x p l a i n t h e annealing phenomena observed under pulsed l a s e r i r r a d i a t i o n . To s t u d y t h i s p o s s i b i l i t y , t h e program energy r e l a x a t i o n t i m e T~ was v a r i e d i n t h e computer
J. F. GIBBONS
12
0.2
0.1 1
l4o0 1200
-
1
0.3
0.4
I 0.26J 0.31J
I
0.39J
t /
A 0.53p m
-
/
P = 50 M W / c m 2
/
-
CRYSTALLINE Si
V
-
1000-
-
IT 3 I-
800
-
-
0
w
a
IT
W
t-
-
--- E g = CONSTANT
I
0
1
2
3
4
5
6
7
.
6
FIGURE 1.5. The e f f e c t of bandgap narrowing and e l e c t r o n i c h e a t conduction on t h e temperature e v o l u t i o n d u r i n g a n n e a l i n g w i t h 0.53 pm r a d i a t i o n . Free c a r r i e r a b s o r p t i o n does not have any a p p r e c i a t i o n e f f e c t a t t h i s wavelength. The from 10 p s up t o 10 n s f o r t h e .53 pm, 50 mW/cm2 r a d i a t i o n . The r e s u l t s are g i v e n i n Fig. 1.6, where t h e s u r f a c e temperature vs. t i m e i s p l o t t e d f o r s e v e r a l v a l u e s of re. For r e of l e s s than 0.1 n s , t h e r e i s no e f f e c t of r e on t h e r e s u l t s . I f T~ i s i n c r e a s e d t o 1 n s , t h e r e i s a s l i g h t i n c r e a s e i n Em. I n c r e a s i n g re t o 10 ns c a u s e s Em t o be about t h r e e t i m e s as much as f o r t h e low r e v a l u e s , but m e l t i n g t e m p e r a t u r e i s s t i l l reached w i t h less than 1 J / c m 2 . These r e s u l t s i n d i c a t e t h a t u n l e s s t h e energy r e l a x a t i o n t i m e i s a t l e a s t comparable t o t h e l e n g t h of t h e laser p u l s e , t h e c o u p l i n g between t h e carriers and t h e l a t t i c e i s s t r o n g enough t o a s s u r e t h a t t h e sample s u r f a c e r e a c h e s t h e m e l t i n g p o i n t i n t y p i c a l p u l s e d laser p r o c e s s i n g cases. Of c o u r s e , e x t r e m e l y l a r g e v a l u e s of r e would delay l a t t i c e h e a t i n g so much t h a t c a r r i e r d i f f u s i o n would cause t h e energy t o be d e p o s i t e d i n a l a r g e enough volume t o p r e v e n t s u r f a c e m e l t i n g , as suggested by Van Vechten. C a l c u l a t i o n s show t h a t , f o r t h e p u l s e used above ( 1 J/cm2 i n 20 n s ) , r e would have t o be about 100 ns b e f o r e m e l t i n g could be avoided L1.281.
1. BEAM PROCESSING OF SILICON
13
ENERGY (J/crnZ)
0.2
FIGURE 1.6.
0.4
0.6
0.8
1.0
The effect of the energy relaxation time, ' c ~ ,
on the annealing results for 0.53 pm laser radiation.
1.1.3
Calculation of Carrier and Lattice in Si by Picosecond Laser Pulses
Temperature Induced
While the energy relaxation time is a noncritical parameter for nanosecond pulse annealing, the situation may be different when picosecond laser pulses are used. To study this case, the previous equations were used together with the assumption that excited carriers instantaneously establish a Maxwell-Boltzmann distribution characterized by a carrier temperature Te (equal for holes and electrons) 11.291. The computer model was used to study annealing with laser pulses of 20 ps duration in the energy range 0.1 - 0.5 J/cm2 at 0.53 um. Both carrier-concentrationa laser wavelength of dependent and constant values of the energy relaxation time re were used with little difference in the basic results. Table 1.4 shows that the choice of T~ has a profound effect on both the carrier temperature and the melt threshold. For comparison with experiment, Liu et al. [1.31] have determined that the melt threshold for 20 ps pulses at a wavelength of 0.53 urn is 0.2 J/cm2 for <111> crystalline Si. Liu et al. also performed in situ electron emission measurements which indicated As that the maximum electron temperature achieved is -5000'K. can be seen from Table 1.4, these data fit the calculations extremely well for a constant energy relaxation time of 1 ps.
14
J. F. GIBBONS
TABLE 1.4. i z a t i o n Time bottom l i n e . Energy densisy (Jfcm 1
R e s u l t s of C a l c u l a t i o n s f o r t h e Constant ThermalThe m e l t t h r e s h o l d s are i n d i c a t e d on t h e
T ~ .
Maximum carrier temperature Te
0.1 pS
=
0.1 0.2 0.4 0.5
950 2200 2600 2900
Em
0.16 J/cm2
1.1.4
K K K K
1 PS 2600 5500 7500 8500
10 ps
K K K K
0.19 J/cm 2
6 9 21 23
400 500 000 500
K K K K
0.4 J/cm2
C a l c u l a t i o n of Melt Depth v s P u l s e Energy
In p r a c t i c a l a n n e a l i n g s i t u a t i o n s it i s n e c e s s a r y t o m e l t t h e s i l i c o n t o a d e p t h t h a t exceeds t h e depth of t h e damage produced by t h e i o n i m p l a n t a t i o n . T h i s r e q u i r e s i n c r e a s i n g t h e p u l s e energy beyond t h e m e l t t h r e s h o l d v a l u e by an amount t h a t i s , v e r y roughly, H p Xd/(l-Rg), where H i s t h e h e a t of f u s i o n and Rg is t h e r e f l e c t i o n c o e f f i c i e n t of l i q u i d s i l i c o n . To s t u d y t h i s s i t u a t i o n , c a l c u l a t i o n s of t h e m e l t d e p t h v e r s u s t i m e have been c a r r i e d o u t by a number of a u t h o r s (Baeri e t a l . [1.32] and Wang e t a l . i1.331) w i t h t h e g e n e r a l r e s u l t shown i n Fig. 1.7. These c a l c u l a t i o n s were performed [1.32] f o r i r r a d i a t i o n w i t h a ruby l a s e r w i t h a p u l s e d u r a t i o n of 50 ns and t h e l a s e r e n e r g i e s as shown i n t h e f i g u r e . A s u r f a c e amorphous l a y e r of s i l i c o n having a t h i c k n e s s of 0.2 urn is assumed t o approximate a t y p i c a l damage d e p t h produced by i o n i m p l a n t a t i o n . A s an example w e see t h a t f o r a p u l s e energy of 2.5 J / c m 2 , t h e c a l c u l a t i o n s p r e d i c t a maximum m e l t d e p t h of a b o u t 6000 A which o c c u r s about 100 n s a f t e r t h e laser p u l s e is i n i t i a t e d . The p r o p a g a t i o n r a t e of t h e t h e l i q u i d - s o l i d i n t e r f a c e toward t h e s u r f a c e can be e s t i m a t e d by o b s e r v i n g t h a t t h e melt f r o n t moves from a d e p t h of -0.6 MI t o t h e s u r f a c e i n about 300 ns, which l e a d s t o an estimate of t h e regrowth r a t e of -2 m/sec. The f a c t t h a t t h i s growth p r o c e s s i s a c t u a l l y e p i t a x i a l r e c r y s t a l l i z a t i o n i s of course not p r e d i c t e d by t h e model and must be e s t a b l i s h e d by t r a n s m i s s i o n e l e c t r o n microscopy, i o n c h a n n e l i n g o r o t h e r c r y s t a l l o g r a p h i c measurement technique. ~
1 . BEAM PROCESSING OF SILICON
15
0 time( ns)
FIGURE 1.7. Calculations of the penetration and retreat of the melt front in a-Si (200 nm>/Si structure for ruby laser pulses of various energies. 1.1.5
Impurity Profiles in Pulse Annealed Silicon
From a device applications viewpoint, one of the most important features of the process has to do with impurity redistribution effects which occur during the annealing process. Those effects have been thoroughly catalogued by the extensive early work of White and co-workers [1.9,1.34]. Figure 1.8 shows B profiles in as-implanted and laser annealed silicon samples taken from their work. The measurements were made by secondary ion mass spectroscopy (SIMS) and therefore yield the total concentration of boron without regard for lattice location. The samples were implanted at 35 keV to a dose of 1 x 1016/cm2. The as-implanted profile is approximately Gaussian as expected, and is only marginally changed by conventional furnace thermal annealing at 900°C for 30 minutes. In contrast, the profile measured after laser annealing (1.6 J/cm2, 60 ns pulse duration) exhibits a substantial redistribution of boron. In particular, after the pulsed laser annealing cycle, the impurity profile becomes almost uniform from the surface down to a depth of approximately 0.2 pm in the crystal, and significant quantities of boron are observed at depths of approximately 0.5 pm.
J . F. GIBBONS
16
I
A S IMPLANTED LASER ANNEALED
o
,02i
1
I
I
_db
i
-E 0
z I020 0
0
0 0
L
u
0
0
0
0
z
e
4
-1
0 0
0 0 0 0
0
0'8
L
I
0
0.4
I
0.3 DEPTH (gL)
02
0.4
0.5
FIGURE 1.8. I l l u s t r a t i n g t h e r e d i s t r i b u t i o n of boron t h a t o c c u r s as a r e s u l t of p u l s e d l a s e r a n n e a l i n g . The r e d i s t r i b u t i o n of t h e i m p l a n t e d boron i s found t o be b o t h p u l s e e n e r g y d e n s i t y and p u l s e number dependent. F i g u r e 1.9 shows e x p e r i m e n t a l r e s u l t s measured a f t e r laser a n n e a l i n g w i t h d i f f e r e n t At a p u l s e e n e r g y d e n s i t i e s i n t h e range 0.64 J/cm2 t o 3.1 J / c m 2 . laser energy d e n s i t y of a p p r o x i m a t e l y 0.6 J/cm2, o n l y s u r f a c e m e l t i n g i s i n i t i a t e d and t h e boron p r o f i l e i s i n d i s t i n g u i s h a b l e from t h a t of t h e as-implanted sample. H a l l e f f e c t measurements and t r a n s m i s s i o n e l e c t r o n microscopy show about 30% of t h e exp e c t e d e l e c t r i c a l a c t i v i t y and s i g n i f i c a n t damage remaining i n t h e form of d i s l o c a t i o n l o o p s under t h i s p u l s e energy. A t e n e r g y d e n s i t i e s 1.1 J/cm2 and g r e a t e r , t h e i m p u r i t y p r o f i l e s are a l m o s t f l a t - t o p p e d i n t h e s u r f a c e r e g i o n , and t h e boron s p r e a d s d e e p e r i n t h e sample as t h e e n e r g y d e n s i t y i s i n c r e a s e d . The s u b s t a n t i a l r e d i s t r i b u t i o n of boron induced by p u l s e d laser a n n e a l i n g cannot be e x p l a i n e d by t h e r m a l d i f f u s i o n i n t h e s o l i d phase because t h e t i m e d u r a t i o n i s t o o s h o r t . However, t h e
1.
40"
BEAM PROCESSING OF SILICON
L UJ
U 0
0.4
17
0.2
0.3
0.4 DEPTH ( p )
0.5
0.6
0.1
FIGURE 1.9. Illustrating the redistribution of boron in pulse annealed silicon using pulses of different energies.
melt depth calculations just discussed show that a region several thousand angstroms deep can be melted for pulse energies greater than about 1 J/cm2. The dopant atoms have a very high diffusion coefficient in the melt and the implanted profile can diffuse rapidly through the liquid while the crystal is molten. For boron in silicon, the experimental profile can be fit by using a liquid diffusion coefficient of 2.4 f 0 . 7 ~ 1 0cm2/sec ~ [1.35] and a diffusion time of 180 ns. These parameters provide the calculated fit shown in Fig. 1.10 and adequately justify the conclusion that normal diffusion processes in the molten state account for the spreading of the implanted profile during pulsed laser annealing.
J. F. GIBBONS
18 5
2
402'
5
f
-E,
-
2
it 40'"
0
sa I2
5
W
0
z
sz 2
g 40" 0
m
5
2
lo'* 5
FIGURE 1.10. Comparison of t h e o r y pulse annealed ion implanted s i l i c o n .
and e x p e r i m e n t f o r ruby
H a l l e f f e c t and t r a n s m i s s i o n e l e c t r o n microscopy measurements made on i m p l a n t e d material a n n e a l e d i n t h i s way show 100% e l e c t r i c a l a c t i v i t y and no d e f e c t s t o a r e s o l u t i o n of a t l e a s t 50 A , p r o v i d e d t h a t t h e amorphous l a y e r i s s u f f i c i e n t l y t h i n . However, i t s h o u l d be mentioned t h a t , s i n c e t h e r e f l e c t i o n coe f f i c i e n t R a t t h e s u r f a c e of a n i r r a d i a t e d sample changes from 0.35 t o 0.7 when t h e s u r f a c e becomes m o l t e n , o n l y 30% of t h e p u l s e over and above t h e m e l t i n g t h r e s h o l d i s a b s o r b e d by t h e s i l i c o n . The r e s u l t i s t h a t i t i s d i f f i c u l t t o m e l t l a y e r s t h a t a r e i n e x c e s s of a p p r o x i m a t e l y 1 pm w i t h o u t h e a t i n g t h e s u r f a c e t o t h e b o i l i n g p o i n t , which l e a d s t o v e r y poor s u r f a c e morphology. A s a r e s u l t , pulsed laser annealing i s g e n e r a l l y u s e f u l only f o r t h e a n n e a l i n g of r e l a t i v e l y s h a l l o w implanted l a y e r s .
1. 1.1.6
BEAM PROCESSING OF SILICON
19
Supersaturation and Segregation
In normal crystal growth from the melt, which is carried out under near-equilibrium conditions (and therefore at very low growth rates), the solid which forms from the liquid at any moment has a concentration of dopant given by Cs = koCL, where ko is the equilibrium distribution coefficient and CL is the dopant concentration in the liquid. It follows that, unless ko = 1, the dopant will segregate into the liquid as crystal growth proceeds under near-equilibrium conditions. However the recrystallization process for pulsed laser annealing is far from equilibrium; growth rates are very high (on the order of meterslsec) and there is then the possiblity that dopants can be incorporated into the solid at concentrations that far exceed the solid solubility limit for near-equilibrium growth processes, thus producing a substitu3-tional supersaturated alloy. Supersaturation will require, or be characterized by, a nonequilibrium distribution coefficient k' that is appropriate for the high speed, nonequilibrium recrystallization process.
If we assume as a limiting case that the diffusion coefficient of the dopant in the liquid phase is so large that the dopant is always uniformly distributed in the liquid that remains at any stage of the recrystallization process, then the impurity distribution following a pulsed laser annealing process will be controlled completely by a "nonequilibrium" distribution coefficient k'. In particular, if the maximum melt depth is denoted by xm, then the concentration incorporated into the solid, C,(x), is given by Cs(x) = k'
@ Xm
[$]
(1-k')
where @ is the dose of the implanted impurity and x is measured from toward the surface. Equation (1.5) shows that, if k' = 1, the implanted dopant will be distributed uniformly over a depth x,,. However if k is less than 1, dopant will be swept toward the surface as recrystallization proceeds, and the dopant concentration will approach infinity at the surface (for the simple model under consideration). Figure 1.11 shows profiles obtained from Eq. (1.5) fork' = 1, 0.1 and 0.01, respectively. Under realistic conditions, of course, it is necessary to include both diffusion and dopant segregation into the liquid to calculate the final impurity profile.
J . F. GIBBONS
20
\\---I
k'= 0.01I
0.2
F I G U R E 1.11.
0.4
0.6
Q8
1.0
Illustrating dopant segregation during resolid-
ification. White et al. [1.361 have studied pulsed laser annealing of high dose implants in Si, using both diffusion and distribution coefficient phenomena in the theoretical analysis of their experimental data. An example of their analysis for a high dose As Rutherford backscattering implant into Si is shown in Fig. 1.12. and channeling analysis shows that all of the As atoms were incorporated into substitutional sites by the pulsed laser anneal. Furthermore, as indicated on Fig. 1.12, (1) the shape of the impurity profile after annealing can be be ex lained entirely by diffusion, so the effective value of k ' for 7$As in Si is unity for these conditions. This is to be compared to an equilibrium value of ko for As in Si of -0.3.
(2) 75As has been incorporated substitutionally at a concentration that is approximately four times the reported maximum equilibrium solubility.
1.
BEAM PROCESSING OF SILICON
o
n
21
\
L o IMPLANTED 0
LASER ANNEALED
-
CALCULATED t k = 1.0)
'I
o
2
10n
1
I
0
I
I I 0.1 DEPTH (pm)
I 0.2
I
I 0.3
FIGURE 1 . 1 2 . Dopant p r o f i l e s f o r 75As (100 keV, 6.4x1Ol6/ em2) i n s i l i c o n compared t o model c a l c u l a t i o n s . The h o r i z o n t a l l i n e i n d i c a t e s t h e e q u i l i b r i u m s o l u b i l i t y l i m i t ( a f t e r Ref. 1.36). Very s u b s t a n t i a l i n c r e a s e s i n k ' over ko were observed by White e t a l . [1.36] f o r a v e r y w i d e range of i m p u r i t i e s , though k ' was found t o be u n i t y f o r only B , P and A s i n S i . For i m p u r i t i e s w i t h low v a l u e s of k ' , such a s Cu i n S i , f o r example, s e g r e g a t i o n of t h e impurity t o t h e s u r f a c e i s t h e dominant p r o c e s s , a s shown i n Fig. 1.13. This process i s c l e a r l y s i m i l a r t o t h e i d e a l i z e d process d e s c r i b e d by E q . (1.5) and shown i n Fig. 1.11 f o r a very l o w value of k. White e t a l . [36] have a l s o e s t a b l i s h e d maximum s u b s t i t u t i o n a l c o n c e n t r a t i o n s f o r dopant i n c o r p o r a t i o n following pulsed l a s e r a n n e a l i n g , f o r dopants t h a t can be r e a d i l y s t u d i e d by backs c a t t e r i n g and channeling a n a l y s i s . The r e s u l t s of t h e i r s t u d y a r e given i n Table 1.5. f o r both k ' and Cs,,ax
J. F. GIBBONS
22
2
a
V
S
2 10''
-1
I-
4000
0
3000
20!0 DEPTH ( A )
6 3 C ~ +(450 heV. 6.9 r40'5/Cm2)
in (ii 1 )
si
FIGURE 1.13. Profile of 63Cu (150 keV, 6.9 x 1015/cm2) implanted into <111> Si before and after laser annealing (after R e f . 1.36).
TABLE 1.5 Distribution Coefficient and Maximum Substitutional Solubility for Dopants in Si Following Pulsed Laser Annealing.
Dopant
:c
Cmax
k'
kO
As
1.0
0.3
Sb Bi
0.1 0.4
0.023 0.0007
Ga
0.2
0.008
4.5~1019
4 .5x1020
In
0.15
0.0004
8x10I7
1. 5x1020
1.5x102
S
6x1021
7 ~ 1 0 ~ 1~.3x1021 4x1OZ0 8x1Ol7
1 . BEAM PROCESSING OF SILICON
23
With regard t o device f a b r i c a t i o n , these supersaturated a l l o y s may prove t o be of some i n t e r e s t , though i s should be mentioned t h a t p r e c i p i t a t i o n of these supersaturated dopants occurs r a p i d l y i n subsequent thermal processing s t e p s , with the development of d i s l o c a t i o n loops, rods and p r e c i p i t a t i o n networks t h a t can be expected t o have d e l e t e r i o u s e f f e c t s on device properties. 1.1.7
Crystallographic Analysis of Pulsed Implanted S i l i c o n
Laser Annealed
Ion
The damage remaining i n ion implanted samples a f t e r an annealing cycle is t r a d i t i o n a l l y studied by two techniques. Transmission e l e c t r o n microscopy (TEM) i s used t o i d e n t i f y cryst a l l o g r a p h i c d e f e c t s such a s d i s l o c a t i o n loops, rods and precip i t a t e s t r u c t u r e s ; and deep l e v e l t r a n s i e n t spectroscopy (DLTS) i s used t o measure the concentration and d i s t r i b u t i o n of r e s i d u a l point d e f e c t s . In t h i s s e c t i o n w e summarize t h e TEM r e s u l t s f o r pulse annealed samples obtained by Narayan and co-orkers (1.371.
Ion Implanted Samples
1.1.7.1
Figure 1.14 shows transmission e l e c t r o n micrographs which compare d i r e c t l y the damage t h a t remains i n the l a t t i c e a f t e r l a s e r annealing and conventional thermal annealing f o r s i l i c o n c r y s t a l s implanted by B, P, and As. A f t e r thermal annealing a t temperatures up t o l l O O ° C , s i g n i f i c a n t damage remains i n the implanted region i n t h e form of d i s l o c a t i o n loops as shown i n t h e r i g h t column of micrographs. By c o n t r a s t , a f t e r l a s e r annealing (- .5 J/cmZ, 60 nsec) no damage ( d i s l o c a t i o n loops or s t a c k i n g f a u l t s ) remains i n the implanted region down t o the r e s o l u t i o n of t h e microscope (- 10 A). D i f f r a c t i o n p a t t e r n s from these same c r y s t a l s have no i r r e g u l a r i t i e s , and t h e o r i e n t a t i o n of t h e r e c r y s t a l l i z e d region is the same a s t h a t of t h e underlying s u b s t r a t e , (001). Upon subsequent heating t o 900°C, a low conc e n t r a t i o n (< 5 x 1014/cm3) of small d e f e c t s (average diameter 30 A develops) i n t h e implanted region.
-
Figure 1.15 shows t h r e e d i f f e r e n t types of microstructure and c r y s t a l s t r u c t u r e i n t h e implanted region which was observed a f t e r annealing with d i f f e r e n t l a s e r energy d e n s i t i e s . These r e s u l t s were obtained using s i l i c o n c r y s t a l s implanted with 75As (100 keV), 1.4 x 10l6/cm2). Both ion channeling and transmission 1600 A dee was made e l e c t r o n microscopy showed t h a t a region amorphous by implantation. Annealing with 0.63 J/cmf caused a 1000 A t h i c k t o become p o l y c r y s t a l l i n e a s indicated by region the d i f f r a c t i o n pattern. This suggests t h a t the c r y s t a l was 1000 A The regrowth was not e p i t a x i a l melted t o a depth of
-
-
-
.
J. F. GIBBONS
24
FIGURE 1 . 1 4 . Comparison of l a s e r ( l e f t column) and t h e r m a l ( r i g h t column) a n n e a l i n g of i o n i m p l a n t e d s i l i c o n , (001) o r i e n t a 11B(35 t i o n . I m l a n t e d s p e c i e s , energy and dose were: tl"p ro? keV, 3x10 5/cm2); middle row - 31P(80 keV, 1x10 / c m ); bottom
-
P
row - 75As(100 keV, 1 x 10l6/cm2). C o n d i t i o n s f o r t h e r m a l anneal30 m i n s ) . Conditions i n g were B,P(llOO°C - 30 m i n s ) ; As(900"C f o r l a s e r a n n e a l i n g were 1.5 J / c m 2 , 60 n s e c ( a f t e r Ref. 1.37).
-
because t h e m e l t e d r e g i o n d i d n o t e x t e n d i n t o t h e undamaged subs t r a t e . A f t e r a n n e a l i n g w i t h 1.07 J / c m Z , which i s v e r y n e a r t h e t h r e s h o l d f o r complete e l e c t r i c a l a c t i v a t i o n , t h e e l e c t r o n d i f f r a c t i o n p a t t e r n showed t h e r e c r y s t a l l i z e d r e g i o n t o be s i n g l e c r y s t a l , but a t h i n s u r f a c e l a y e r w a s observed containing r e s i d u a l damage i n t h e form of d i s l o c a t i o n l o o p s . Stereomicroscopy showed t h a t t h e r e s i d u a l damage was c o n f i n e d t o a d e p t h of 200 8 . The o r i g i n of t h i s damaged l a y e r a t t h e s u r f a c e i s n o t w e l l u n d e r s t o o d ,
-
1.
BEAM PROCESSING OF SILICON
25
75As( 100 keV, 1. ~ x ~ O ~ ~ / C U I ~ )
FIGURE 1.15. M i c r o s t r u c t u r e developed i n s i l i c o n by annealThe energy d e n s i t y i n g w i t h d i f f e r e n t laser energy d e n s i t i e s . used f o r a n n e a l i n g i s i n d i c a t e d b e s i d e e a c h micrograph and corr e s p o n d i n g d i f f r a c t i o n p a t t e r n . P u l s e d u r a t i o n was 25 n s e c ( a f t e r Ref. 1.37). b u t i t may be r e l a t e d t o r a d i a n t h e a t l o s s e s i n t h e s u r f a c e r e g i o n d u r i n g t h e regrowth p r o c e s s . R e s i d u a l damage a t t h e s u r f a c e h a s been observed f o r B , P and A s i m p l a n t s a f t e r i r r a d i a t i o n w i t h l a s e r energy d e n s i t i e s v e r y n e a r t h e t h r e s h o l d f o r complete a n n e a l i n g (- 1.1 J/cm2 f o r o u r implant c o n d i t i o n s ) .
26
J. F. GIBBONS
1.1.7.2
Thermally-Diffused Samples
In t h e r m a l l y d i f f u s e d specimens, Narayan and White 11.371 have shown t h a t d i s l o c a t i o n l o o p s and p r e c i p i t a t e s formed d u r i n g phosphorus o r boron d i f f u s i o n a t l l O O ° C f o r 30 minutes can be d i s solved by t h e u s e of p u l s e d laser i r r a d i a t i o n . For a phosphorusd i f f u s e d sample, one p u l s e of energy d e n s i t y 1.5 3 / c m 2 produced an a n n e a l e d , d e f e c t - f r e e r e g i o n (i.e., c o n t a i n i n g no phosphorous The p r e c i p i t a t e s and d i s l o c a t i o n l o o p s ) t h a t was 0.29 pm deep. depth of t h e d e f e c t and p r e c i p i t a t e f r e e r e g i o n w a s found t o be 0.38 pm a f t e r two such p u l s e s ; and a f t e r seven p u l s e s t h e depth was determined t o be 0.58 um. A s a t u r a t i o n phenomenon i n t h e depth of t h e annealed r e g i o n w a s observed a f t e r approximately f i v e The r e s u l t s of 1.5 J / c m 2 , 50 n sec p u l s e are summarized pulses. i n Table 1.6. TABLE 1.6 Ruby Laser (X = 0.694 pm) Annealing Parameters i n Phosphorous D i f f u s e d S i l i c o n .
P u l s e Energy D n s i t y E(Jou1e cm )
5
Pulse Duration (Nano secs)
50
Number of Pulses
Depth of Annealing (pm)
1.5 1.5 1.5 1.5
50 50 50
1 2 5 7
0.29 0.38 0.56 0.58
2 .o 2 .o 2 .o 2 .o 3 .O
20 20 20 20 20
1 3 5 14 1
0.57 0.70 0.75 0.77 0.95
The d e p t h s of t h e annealed r e g i o n s were s i m i l a r l y determined f o r 1, 3, 5 and 1 4 laser p u l s e s of energy d e n s i t y 2.0 J/cm2 and p u l s e d u r a t i o n 20 n sec w i t h t h e r e s u l t s summarized i n Table 1.6. Again s a t u r a t i o n i n t h e a n n e a l i n g d e p t h w a s observed a f t e r about f i v e laser p u l s e s . I t s h o u l d be mentioned t h a t i n t h e specimens i r r a d i a t e d w i t h t h r e e o r more laser p u l s e s , a small amount of r e p r e c i p i t a t i o n of phosphorous n e a r t h e s u r f a c e was observed. T h i s i s probably r e l a t e d t o t h e s e g r e g a t i o n c o e f f i c i e n t of phosFor 2.0 and phorous i n s i l i c o n being s l i g h t l y less t h a n u n i t y . 3.0 J/cm2 (20 n s e c ) l a s e r p u l s e s , t h e i n c r e a s e i n t h e a n n e a l i n g depth (from 0.57 m t o 0.95 m) is approximately p r o p o r t i o n a l t o t h e p u l s e energy d e n s i t y . A complete a n n e a l i n g of d e f e c t s up t o one micron was p o s s i b l e w i t h one p u l s e of E = 3.3 J / c m 2 and T = 20 n s e c .
1. 1.1.8
BEAM PROCESSING OF SILICON
27
Deep Level T r a n s i e n t Spectroscopy
Capacitance t r a n s i e n t spectroscopy provides a very powerful technique f o r measuring t h e energy l e v e l s and d e f e c t s t a t e dens i t i e s i n processed s i l i c o n . The p r i n c i p l e of t h e method was f i r s t d e s c r i b e d by Lang L1.381 and r e c e n t l y summarized f o r l a s e r a n n e a l i n g a p p l i c a t i o n s by Johnson e t a l . [1.39]. B r i e f l y , carrier t r a p p i n g c e n t e r s surrounding a pn j u n c t i o n o r a Schottky b a r r i e r are charged by a p p l i c a t i o n of an a p p r o p r i a t e v o l t a g e pulse. T h e i r d i s c h a r g e i s t h e n encouraged by scanning t h e sample over a tempera t u r e range from l i q u i d n i t r o g e n temperature t o room temperature. Each t r a p p i n g c e n t e r i s i d e a l l y c h a r a c t e r i z e d by an energy l e v e l and t h e r e f o r e a s p e c i f i c temperature a t which c a r r i e r emission from t h e t r a p w i l l be maximized. T h i s c a r r i e r emission w i l l change t h e charge d e n s i t y i n t h e d e p l e t i o n r e g i o n of a r e v e r s e b i a s e d pn j u n c t i o n o r Schottky b a r r i e r , t h u s p r o v i d i n g a c a p a c i t i v e t r a n s i e n t s i g n a l t h a t can be used t o d e t e c t t h e t r a p p i n g c e n t e r . F i g u r e 1.16 shows a t y p i c a l d e f e c t s t a t e spectrum observed i n a p u l s e l a s e r annealed, boron-doped, Czochralski grown S i subThe dominant d e f e c t s t a t e i s s t r a t e , a f t e r Benton e t a l . [1.40]. a h o l e t r a p [H(0.36)1 l o c a t e d approximately a t an energy l e v e l t h a t i s 0.36 e V above t h e v a l e n c e band. By measuring t h e concent r a t i o n of t h i s d e f e c t vs. laser power and i m p l a n t a t i o n dose, Benton e t a l . were a b l e t o e s t a b l i s h t h a t : 1. The t h r e s h o l d power f o r r e s i d u a l minimum d e f e c t d e n s i t y a f t e r pulsed laser a n n e a l i n g d e c r e a s e s w i t h i n c r e a s i n g implant dose, due presumably t o t h e enhanced coupling of t h e near-bandedge l i g h t from t h e Nd:YAG 1.06 pm laser.
2 . A h i g h e r d e f e c t s t a t e c o n c e n t r a t i o n was observed n e a r np' j u n c t i o n s than i n t h e s u b s t r a t e ( t h e pn j u n c t i o n s were f a b r i c a t e d by A s implanted followed by laser a n n e a l i n g a t a wavelength of 1.06 pm). 3 . The average d e f e c t s t a t e c o n c e n t r a t i o n appears t o show a minimum when t h e laser power i s j u s t s u f f i c i e n t t o produce a m e l t depth t h a t p r e c i s e l y matches t h e range of t h e i o n damage. Deeper m e l t i n g may l e a v e quenched-in d e f e c t s beyond t h e j u n c t i o n . The c o n c e n t r a t i o n of t h e s e d e f e c t s can be reduced by a subsequent low temperature f u r n a c e anneal (80OOC f o r 30 min). 1.1.9
Pulsed E l e c t r o n Beam Annealing of Ion Implanted S i l i c o n
The u s e of pulsed e l e c t r o n beams f o r a n n e a l i n g i o n implanted s i l i c o n w a s f i r s t developed f o r t h e purpose of a n n e a l i n g i o n i m p l a n t e d s o l a r c e l l s [1.411. The d e t a i l s of t h e a n n e a l i n g mechani s m were explored l a t e r by Kennedy e t a l . [1.421, u s i n g Rutherford
J. F. GIBBONS
28
sb I M K A N T LZ 62MWIcm2 T, = 182rns
; L HlO53)
5c
100
150
250
200
TEMPERATURE
(OK1
FIGURE 1.16. Capacitance t r a n s i e n t d e f e c t s t a t e s p e c t r a of l a s e r annealed, boron doped C z o c h r a l s k i grown s i l i c o n ( a f t e r R e f .
1.40). backscatteringlchanneling s p e c t r o s c o p y and TEM t o s t u d y t h e annealing process. The s i l i c o n c r y s t a l s used f o r t h e s t u d y were (100) and (111) i n o r i e n t a t i o n , and were implanted w i t h e i t h e r 28Si+ o r 75As+ o r b o t h . The i m p l a n t s were c a r r i e d o u t a t 25, 50, 100 and 200 keV. The c o n d i t i o n s used f o r t h e p u l s e d e l e c t r o n beam were: an energy d e n s i t y of 0.8-1.2 J/cm2 (over a 50 c m diameter beam) and a p u l s e l e n g t h of 100 n s . The r e s u l t s w e r e as f o l l o w s .
1.1.9.1
75As+-implanted Samples
Pulsed e l e c t r o n beam a n n e a l i n g a t an energy of 0.8 J/cm2 succeeded i n completely a n n e a l i n g t h e samples f o r a l l of t h e implant e n e r g i e s . S u b s t a n t i a l r e d i s t r i b u t i o n of t h e 75As w a s observed, l e a d i n g t o an i m p u r i t y p r o f i l e s i m i l a r t o t h a t o b t a i n e d w i t h p u l s e d laser a n n e a l i n g . Kennedy e t a l . e s t i m a t e d t h e d i f f u s i o n c o e f f i c i e n t f o r As r e q u i r e d t o e x p l a i n t h e A s p r o f i l e d e p t h t o be l o 4 cm2/sec, which i s t y p i c a l of l i q u i d s t a t e d i f f u s i v i t i e s . SIMS measurements of a h i g h d o s e , 75As implanted s i l i c o n sample annealed w i t h a pulsed e l e c t r o n beam a t an i n c i d e n t energy of 1.1 JIcm2 were o b t a i n e d by Wilson e t a l . [1.42] and are shown i n F i g . 1 . 1 7 . These d a t a are very similar t o t h e r e s u l t s f o r p u l s e d l a s e r a n n e a l i n g . They s u p p o r t t h e m e l t i n g h y p o t h e s i s and l e a d t o t h e c o n c l u s i o n t h a t p u l s e d laser and e l e c t r o n beam a n n e a l i n g of an i o n implanted sample proceed by t h e same b a s i c mechanism, though of c o u r s e t h e c o u p l i n g of t h e beam t o t h e s o l i d i s d i f f e r e n t f o r t h e two c a s e s .
1. BEAM PROCESSING OF SILICON
29
2
00
102’
00
**
L-
0
lo L 5
*5
0
*.
_I
**.
0
o A s IMPLANTED ELECTRON BEAM ANNEALED
**
2i
10’9 0
0.1
DEPTH (pl
0.2
0.3
7 5 A s + ( f 0 0 keV, 1.f x f0’6/crn2)
FIGURE 1.17. 75As profile in the as-implanted electron beam annealed condition. 1.1.9.2
and pulsed
28Silicon-implanted Samples
A similar experiment carried out by Kennedy et al. [ 1 . 4 1 ] on Si implanted samples (i.e., Si into Si) showed that full annealing was possible for the implants at 25 and 50 keV only. For higher implant energies the melt depth was insufficient to envelop the implantation damaged layer so that the recrystallization did not start from a single crystal.
Such a result is also obtained for pulsed laser annealing when the pulse energy is insufficient to melt the sample to a depth that exceeds the implantation-damaged region. The central conclusions from the studies of Wilson et al. [1.42] and Kennedy et al. [1.41] are that, when laser and electron
beam systems are optimized, they give essentially the same results. In particular, the do ant redistribution is the same; the fraction of dopant atoms (p5As) annealed into substitutional sites is the same; and TEM shows an absence of defects to a resolution of -50 fi in both cases.
30
J. F. GIBBONS
Summary of Results for Pulsed Beam Annealing
1.1.10
To summarize the foregoing, the mechanism of annealing for both Q-switched lasers and pulsed electron beams involves melting of a surface layer followed by a liquid phase epitaxial regrowth on the underlying substrate. “he critical parameter for this process is the pulse energy density. The principal electrical characteristics of the annealed layer are:
1. 100% substitutionality of the implanted dopant, even for concentrations that exceed the solid solubility. 2.
No residual defects in TEM to 50 A resolution.
3 . Redistribution of the implanted dopant via diffusion in a liquid layer during the recrystallization process.
4 . Residual point defects that may require a subsequent low temperature ( 8 O O 0 C , 30 min) anneal for their removal.
1.2
CENTRAL FEATURES OF CW
BEAM ANNEALING OF ION IMPLANTED
SILICON
In contrast to the pulsed annealing process, a scanning laser or electron beam provides an extremely convenient and highly controllable means for heating the surface of a semiconductor to a given temperature while holding the body of the material at a convenient low temperature (typically 3.50-500°C). Recognition of this fact leads naturally to a comparison of the annealing mechanism with the annealing process for furnace annealing. In what follows we first give a brief description of furnace annealing as it applies to implantation-amorphized layers, and then proceed to a discussion of the scanning beam alternative.
1.2.1
Furnace Annealing Process
Conventional furnace annealing of implantation-amorphized silicon proceeds by what is called a solid phase epitaxial (SPE) recrystallization process. Atoms at the amorphous-crystalline interface rearrange themselves into crystalline locations by a process that appears to involve vacancy diffusion in both the crystalline and the disordered layers. The process is described quantitatively by an activation energy Ea and proceeds at a rate r(t> = ro exp(-Ea/kT)
1.
BEAM PROCESSING OF SILICON
31
where r(t) is normally measured in 8/second. Measurements of r(t) made by Cspregi et al. i1.431 in the temperature range 400'-600'C undoped silicon (amorphized by the implantation of Si+) show an activation energy of E = 2.35 eV and a preexponential multiplier of 3.22 x lOI4 8l:ec. Using these values at a temperature of 800'C gives a furnace regtowth rate of of about 0.3 pm/sec for undoped Si. In other words, a 0.5 pm thick amorphized layer can recrystallwe by the solid phase epitaxial process process on the underlying substrate in less than two seconds if the layer is held at a temperature of 800'C. The preexponential factor is found to be larger if dopants such as B y P and As are added, and smaller if oxygen is incorporated l1.431. Using Rutherford backscatteringlchanneling analysis, the implanted dopants are found to be normally incorporated into substitional sites during this recrystallization process, leading to high electrical activity. However, point defects and trapping centers form during the recrystallization process (related to the vacancy motion by which the process proceeds), and these defects must be annealed to obtain good pn junction characteristics; i.e., high carrier mobilities and reasonable carrier lifetimes. The motion and agglomeration of defects during the annealing is especially troublesome if the implant dose and energy are not sufficient to produce an amorphous layer that envelops the implanted dopant. Such a condition is most often obtained when B is implanted into Si, a case which has been thoroughly studied. A 30 minute anneal in the temperature range 800°C-10000C then produces a combination of dislocation loops, rods and precipitates that are nearly imune to further furnace annealing [1.44, 1.451. Partly as a result of these problems it is customary to use a so-called "two-stage" annealing process (1.461. The first stage is carried out at -600°C for 112 hour and is intended only to recrystallize the implanted region. Point defects and small damage clusters remain in the regrown layer after this stage, which often lead to low electrical activity and poor carrier mobility and lifetime. The low temperature anneal is then typically followed by a 1000°C, 10 minute anneal to remove point defects, increase carrier mobility and lifetime, and provide for some impurity diffusion so that pn junctions formed by the process are located outside the region of residual damage. Such a process is especially important in the annealing of implanted <111> Si; i.e., for most bipolar applications. In contrast a single stage anneal can be sufficient for less demanding applications in Si.
32
J. F. GIBBONS
1.2.2
B a s i c CW L a s e r Annealing Systems
The b a s i c cw l a s e r a n n e a l i n g p r o c e s s i s i d e n t i c a l t o t h e f u r nace a n n e a l i n g p r o c e s s j u s t d e s c r i b e d , e x c e p t t h a t t h e sample i s maintained a t t h e a n n e a l i n g t e m p e r a t u r e f o r such a s h o r t time t h a t o n l y t h e f a s t e s t t h e r m a l p r o c e s s e s , i n c l u d i n g s o l i d phase e p i t a x y , can be c a r r i e d t o completion. Dopant p r e c i p i t a t i o n and t h e f o r m a t i o n of d i s l o c a t i o n l o o p s and r o d s a r e g e n e r a l l y n o t observed s i n c e t h e r e i s i n s u f f i c i e n t time f o r them t o form. A b a s i c system t h a t can be used f o r scanned c w l a s e r a n n e a l i n g i s shown i n F i g . 1.18. I t c o n s i s t s of an A r ( o r o t h e r ) c w l a s e r t h a t i s p a s s e d t h r o u g h a s e r i e s of l e n s e s , beam-shaping opt i c s and f i l t e r s and t h e n d e f l e c t e d o n t o a sample t h a t i s mounted i n t h e f o c a l p l a n e of t h e f i n a l l e n s [ 1 . 4 7 ] . The beam s h a p i n g o p t i c s a r e used f o r t h e r e c r y s t a l l i z a t i o n s t u d i e s d e s c r i b e d i n Chapter 5 of t h i s volume. For t h e a n n e a l i n g e x p e r i m e n t s r e p o r t e d The sample i s mounted on a h e r e no beam s h a p i n g i s r e q u i r e d . computer d r i v e n , h e a t e d XY t a b l e which p e r m i t s t h e sample t o be scanned u n d e r n e a t h t h e beam. A s an a l t e r n a t i v e , t h e beam can be swept w i t h X and Y m i r r o r s . The X m i r r o r can be mounted on a galvonometer t h a t i s d r i v e n w i t h a t r i a n g u l a r waveform and t h e Y m i r r o r mounted on a galvonometer t h a t i s d r i v e n by a s t a i r c a s e waveform. T h i s arrangement p e r m i t s t h e beam t o be scanned a c r o s s
2OX EXPANDER WITH SPATIAL
18-w
CONDENSING
SCAN
CONTROL
STAGE
FIGURE 1.18. A g e n e r a l s c h e m a t i c of t h e a n n e a l i n g a p p a r a t u s , i n c l u d i n g A r + l a s e r , l e n s p e r p e n d i c u l a r X and Y m i r r o r s , and a vacuum sample h o l d e r .
1.
BEAM PROCESSING OF SILICON
33
t h e t a r g e t i n t h e X d i r e c t i o n , stepped by a c o n t r o l l e d Y increment and then scanned back a c r o s s t h e t a r g e t i n r e v e r s e X d i r e c t i o n . I n d i v i d u a l s c a n l i n e s can be overlapped o r not by a p p r o p r i a t e adjustments of t h e Y s t e p . The samples a r e mounted on a s t a g e t h a t can be heated t o C o n t r o l of t h e annealing ambient can be o b t a i n e d by about 500". p l a c i n g a c y l i n d r i c a l q u a r t z j a c k e t around t h e sample h o l d e r and pumping a p p r o p r i a t e gases i n t o t h i s j a c k e t . A v a r i e t y of l e n s e s , l a s e r powers, scan r a t e s , and sample temperatures have been found t o produce e s s e n t i a l l y p e r f e c t a n n e a l i n g of i o n implanted semicond u c t o r s . I f t h e samples are h e l d a t room temperature, a t y p i c a l s e t of a n n e a l i n g c o n d i t i o n s c o n s i s t s of a l a s e r output (Ar, m i l t i l i n e mode) of 7 W focused through a 79 mm l e n s i n t o a 38 micron s p o t on t h e t a r g e t . The s p o t i s t y p i c a l l y scanned a c r o s s t h e t a r g e t a t a r a t e of approximately 2.5 c m p e r second. Adjacent scan l i n e s are overlapped by approximately 30% t o produce f u l l a n n e a l i n g i n t h e overlapped a r e a s . The l a s e r power f o r f u l l a n n e a l i n g can be reduced (and t h e width of t h e annealed l i n e i n c r e a s e d ) by i n c r e a s i n g t h e sample temperature.
1.2.3
Surface Temperature P r o f i l e s
The beam geometry and scanning c o n d i t i o n s l i s t e d above l e a d This i s t o be comt o a s p o t dwell t i m e on t h e o r d e r of 1 m s . pared t o a thermal t i m e constant ( f o r a 40 u m cube of S i ) of approximately 10 u s , from which it follows t h a t t h e semicond u c t o r s u r f a c e has adequate t i m e t o come t o thermal e q u i l i b r i u m w i t h t h e scanning heat source. It is t h e r e f o r e p o s s i b l e t o c a l c u l a t e e x a c t l y t h e s u r f a c e temperature t h a t w i l l be produced a t the c e n t e r of the s p o t by simply s o l v i n g t h e s t e a d y s t a t e e q u a t i o n f o r h e a t flow. C a l c u l a t i o n s of t h i s type have been carr i e d out f o r both c i r c u l a r and e l l i p t i c a l scanning beams [1.48]. The d e t a i l s of t h e c a l c u l a t i o n s a r e p r e s e n t e d i n Chapter 2 of t h i s volume f o r a number of d i f f e r e n t sources and scanning geometries. The r e s u l t s f o r a beam of c i r c u l a r c r o s s s e c t i o n i r r a d i a t i n g a The c a l c u l a t i o n s s i l i c o n s u b s t r a t e a r e shown i n Fig. 1.19. l e a d i n g t o t h e s e r e s u l t s account f o r t h e temperature dependence of t h e thermal c o n d u c t i v i t y and s p e c i f i c h e a t , and have proven t o g i v e very a c c u r a t e e s t i m a t e s of the s u r f a c e temperatures achieved w i t h t h e l a s e r beam.
Two f e a t u r e s of t h e s e curves a r e worth p a r t i c u l a r a t t e n t i o n . F i r s t , t h e h o r i z o n t a l a x i s i s i n u n i t s of power p e r u n i t r a d i u s of t h e beam, t h i s being due t o t h e f a c t t h a t , f o r a small beam diameter, the heat flow problem has e s s e n t i a l l y hemispherical symmetry. A s a r e s u l t t h e s u r f a c e temperature i s a f u n c t i o n of power divided by s p o t r a d i u s r a t h e r than power p e r u n i t a r e a of t h e beam. Secondly, t h e s i z e a b l e decrease of thermal c o n d u c t i v i t y
34
J. F. GIBBONS
p = '(I-R) (kW/cm)
R,
FIGURE 1.19. The t r u e maximum temperature (X=Y=Z=V=O) i n for S i is p l o t t e d v e r s u s t h e normalized power p (p=P(l-R)/Rx) d i f f e r e n t s u b s t r a t e backsurface temperatures. of S i w i t h temperature l e a d s t o a s i t u a t i o n i n which t h e temperat u r e a t t h e i r r a d i a t e d s u r f a c e is a very s e n s i t i v e f u n c t i o n of t h e backsurf ace temperature. For example, a s u r f ace temperature a t the c e n t e r of a 40 micron s p o t of 1 0 0 0 ° C can be o b t a i n e d w i t h a 5 w a t t laser o u t p u t i f t h e b a c k s u r f a c e temperature ( o r thermal If t h e backsurface temperature i s reduced t o b i a s ) is 35OOC. 150°C, t h e temperature i n t h e c e n t e r of t h e i r r a d i a t e d a r e a Hence i t i s very important t o c o n t r o l drops by n e a r l y 50OoC. t h e backsurface temperature a c c u r a t e l y , and i t is a l s o p o s s i b l e t o use r e l a t i v e l y low laser powers i f s u b s t a n t i a l thermal b i a s can be employed.
1.2.4
Basic Mechanism of CW Beam Annealing
A s mentioned e a r l i e r , t h e b a s i c mechanism of cw beam annealA v e r y convincing i l l u s t r a t i o n of i n g i s s o l i d phase e p i t a x y . t h i s f a c t i s provided by t h e work of Roth e t a l . [1.461 and i l l u s Here t h e regrowth rate of both implantationt r a t e d i n Fig. 1.20. amorphized and UHV d e p o s i t e d S i l a y e r s on a S i s u b s t r a t e i s p l o t t e d a s a f u n c t i o n of t h e s u r f a c e temperature produced by the laser. The growth rate was determined from time-resolved r e f l e c t i v i t y measurements. The d a t a show e x a c t l y t h e form expected f o r a process d e s c r i b e d by Eq. (1.6), and i n f a c t t h e growth r a t e a t 800°C i s found t o be 0.3 pm/sec, i n e x c e l l e n t agreement w i t h t h e furnace a n n e a l i n g r e s u l t . Hence t h e regrowth p r o c e s s i s q u a n t i t a t i v e l y i d e n t i f i e d a s s o l i d phase r e c r y s t a l l i z a t i o n . A d d i t i o n a l i n f o r m a t i o n on t h e growth mechanism i s provided i n Chapter 3 of t h i s volume.
1. BEAM PROCESSING OF SILICON
35
TEMPERATURE, OC
IWO
sw
BOO
1w
600
I
I
I
-OL-S,~10Ol
(MULTIPLE IMPLANT1 E,=26BeV v o = 301" d c m l s e c
V UHV-OEPOSITEOSn FILM ON CLEAN L 11001 SUBSTRATE SPE PERFORMEDIN AIR AMBIENT E. = 2 11 rV
-
104
I I
g
\
1
I
1
I
I
10
11
12
13
14
\
1hT. *v-'
FIGURE 1.20. SPE growth rates for amorphous Si layers formed by ion-implantation and U W evaporation, determined from timeresolved reflectivity measurements (after Ref. 1.46). 9 9 8 1 Elsevier Science Publishing Co.
1.2.5
Impurity Profile
Figure 1.21 shows the impurity profiles obtained by SIMS under as-implanted, laser annealed and thermally annealed conditions for As'implanted into (100) Si under typical conditions. The most striking feature of this figure is that the laser annealed profile is identical to the as-implanted profile. In other words, there has been no diffusion of the implanted species during the laser anneal. Furthermore, the as-implanted impurity distribution is matched exactly by the Pearson type IV distribution function using LSS moments 11.501, so the experimental and theoretical profiles are in excellent agreement. The thermal anneal shows the well-catalogued impurity redistribution given by the open triangles in Fig. 1.21. Majority carrier profiles and carrier mobilities were obtained by sheet resistance and Hall effect measurements and are shown in Fig. 1.22. As can be seen the carrier concentration profile also fits the Pearson type IV distribution quite well. Under very high dose conditions (> 10l6/cm2), 75As precipitation can occur near the peak of the profile, leading to a residual, nonsubstitutional As content of approximately 5% 11.481. As in the case of pulsed beam annealing, dopant concentrations can
J . E GIBBONS
36
As IMPLANTED oLASER ANNEAL ATHERUM ANNEAL (tOOO°C 30min) 0
DISTRIBUTION WITH LSS RANGE STATISTICS
SPUTTERING TIME (set3
FIGURE 1.21. C a r r i e r c o n c e n t r a t i o n and m o b i l i t y p r o f i l e s obtained on As-implanted samples annealed w i t h a scanning A r cw l a s e r beam. exceed t h e e q u i l i b r i u m s o l i d s o l u b i l i t y l i m i t s when implanted samples a r e annealed u s i n g a scanning beam, though p r e c i p i t a t i o n of t h e s e dopants occurs when t h e samples a r e s u b j e c t e d t o f u r t h e r high temperature p r o c e s s i n g s t e p s . The s u b j e c t of dopant i n c o r p o r a t i o n and p r e c i p i t a t i o n i s thoroughly d i s c u s s e d i n Chapter 3 of t h i s volume.
1.2.6.
Transmission E l e c t r o n Microscopy
T y p i c a l r e s u l t s of TEM performed on t h e r m a l l y annealed and The t h e r m a l l y l a s e r annealed samples are shown i n Fig. 1.23. annealed sample (Fig. 1.23a) shows a s i n g l e c r y s t a l d i f f r a c t i o n p a t t e r n w i t h a b r i g h t f i e l d micrograph c o n t a i n i n g t h e u s u a l d e f e c t c l u s t e r s and dislocation v a r i e t y of -2008, diameter loops. The l a s e r annealed sample shown i n Fig. (1.23b) i s essent i a l l y f r e e of any d e f e c t o b s e r v a b l e i n TEM except near t h e It should boundary between c r y s t a l l i z e d and amorphous r e g i o n s . be emphasized a l s o t h a t no d e f e c t s a r e observed i n TEM i n t h e region where a d j a c e n t s c a n l i n e s overlap.
1.
BEAM PROCESSING OF SILICON
37
As+. 100 keV, 3~10’4/cm2
200
400 600 800 1000 1200 DEPTH
(8)
FIGURE 1.22. C a r r i e r c o n c e n t r a t i o n and m o b i l i t y p r o f i l e s o b t a i n e d on As-implanted samples annealed w i t h a scanning A r cw laser beam.
1.2.7
B-implanted S i n g l e C r y s t a l S i
A s i m i l a r s e t of experiments has been performed i n B-implanted Si. The c e n t r a l r e s u l t s a r e i d e n t i c a l t o t h o s e d e s c r i b e d I n p a r t i c u l a r , 100% e l e c t r i c a l a c t i above f o r As-implanted S i . v i t y can be obtained w i t h no d i f f u s i o n of t h e implanted species from i t s as-implanted p r o f i l e . R e c r y s t a l l i z a t i o n i s a l s o p e r f e c t 20 A. These r e s u l t s a r e a s judged by TEM t o a r e s o l u t i o n of independent of whether t h e B i s implanted i n t o pre-amorphized S i or directly i n t o single crystal material.
-
1.2.8
Residual Defects Following CW Laser Annealing
A s mentioned above, cw l a s e r annealing of ion-implanted s i l i c o n i s t y p i c a l l y c a r r i e d o u t under c o n d i t i o n s t h a t produce a s u r f a c e temperature of approximately 1000-llOO°C w i t h a backs u r f a c e temperature of 35OoC. This thermal g r a d i e n t can produce a r a p i d d i f f u s i o n of vacancies and o t h e r p o i n t d e f e c t s from t h e s u r f a c e i n t o t h e c o o l e r i n t e r i o r of t h e sample and a consequent r e d u c t i o n of m i n o r i t y carrier l i f e t i m e .
38
J . F. GIBBONS
FIGURE 1.23. E l e c t r o n micrographs of As-implanted s i l i c o n s u b j e c t e d t o thermal a n n e a l of 1000°C 30 min ( a ) and laser anneali n g ( b ) ; i n s e r t s show d i f f r a c t i o n p a t t e r n s which are t y p i c a l f o r t h e i r regions. P o i n t d e f e c t d e n s i t y p r o f i l e s i n b o t h cw l a s e r annealed and cw e l e c t r o n beam annealed ion-implanted s i l i c o n have been measured These r e s u l t s and o t h e r s of importance by Johnson e t a l . [1.51]. f o r u n d e r s t a n d i n g and applying cw beam p r o c e s s i n g are p r e s e n t e d and thoroughly d i s c u s s e d i n Chapter 4 of t h i s volume. A s an example of t h e r e s u l t s , we show i n Fig. 1.24 t h e d e f e c t d e n s i t y s i g n a l o b t a i n e d from a c o n s t a n t c a p a c i t a n c e (CC)-DLTS measurement made on cw l a s e r annealed s i l i c o n . From t h i s f i g u r e w e see t h a t a s i g n i f i c a n t d e n s i t y of p o i n t d e f e c t s remains immedia t e l y a f t e r t h e a n n e a l i n g cycle. The f i g u r e a l s o shows t h a t t h e s e d e f e c t s can be reduced t o low l e v e l s (comparable t o f u r n a c e ann e a l i n g ) by a subsequent h e a t t r e a t m e n t a t 85OOC f o r 1 1 2 hour.
1. BEAM PROCESSING OF SILICON
2c
Epitaxial Silicon N-Tw, (100) Salt Implanted e,, = 347 sac-’
(- 0.49 eV)
I\
AWD : 5920 -* 6610 A
E 15
1I
-AWD : 7160-+ A 8540A
I\ I \
-->
39
/\10.iavj
.I
v)
5 10
3
(0.19 SV)
/pro>
0
CW L m r Annealed
5
Furnace Annealed (1000 C. 30 m i d
0 -1
I 100
I
150
I
200 T(K)
1
I
I
250
300
330
FIGURE 1.24. Defect density signal obtained from DLTS measurements of cw laser annealed silicon. The density profile for a particular defect can also be measured by a sophisticated application of the CC-DLTS technique. An example is shown in Fig. 1.25, where we see that a substantial density of defects can be produced at depths up to -1 pm immediately following cw beam processing under typical conditions. As mentioned above, the defect density can be removed by subsequent thermal processing, sometimes at temperatures as low as 450°C (see Ref. 1.51 and Chapter 4 of this volume). 1.2.9
Summary
To summarize the foregoing, cw laser annealing proceeds by solid phase epitaxial recrystallization. For “small“ irradiated areas (40 urn spot), the critical parameter is the beam power per unit radius. The principal results of the annealing are: 1. 100% substitutionality of the implanted dopant even for concentrations that exceed the solid solubility.
2.
No residual defects in TEM to 50 A resolution.
3.
No dopant redistribution during annealing.
J. F. GIBBONS
'z
C
oapnl [lo' A1
FIGURE 1.25.
Defect density vs depth for a cw laser annealed silicon sample. 4 . Residual point defects remain below the recrystallized layer that require a subsequent low temperature anneal for their removal (8OOOC for 10 minutes).
For VLSI applications, the absence of dopant redistribution may he a significant feature of this annealing process. In any case it is an ideal companion for the implantation process because the dopants are annealed into their as-implanted sites. Since the implanted impurity profile can be calculated with considerable accuracy, this process provides a means of assuring that computer simulations of impurity profiles based on the implantation process can he realized with precision.
1.3
AXNEALING WITH LARGE DIAMETER SCANNING CW SOURCES
A very attractive alternative for cw laser annealing is the use o f cw arc sources. Arc sources can be configured to produce both large diameter annealing spots and ribbon-shaped beams. These beams are scanned at speeds that produce dwell times on the order of seconds, which substantially reduces the severity of the thermal shock that i s often obtained with a cw laser. Such an annealing time i s , however, still short enough to prevent excessive dopant redistribution during annealing of i o n implanted semiconductors. Furthermore, the throughput of an arc lamp or
1.
BEAM PROCESSING OF SILICON
41
ribbon e l e c t r o n beam annealing system can a l s o be very h i g h (hundreds of 4" wafers p e r hour) and, i n a d d i t i o n , both of t h e s e a n n e a l i n g techniques a r e i n s e n s i t i v e t o a n t i - r e f l e c t i v e d i e l e c t r i c coatings, unlike t h e i r l a s e r counterpart. Both a r c lamps and e l e c t r o n beams can be used f o r a n n e a l i n g i n two p o s s i b l e modes, an i s o t h e r m a l mode and a h e a t sunk mode. I n t h e i s o t h e r m a l mode t h e e n t i r e wafer reaches t h e a n n e a l i n g t e m perature. The dominant h e a t l o s s mechanism i n t h i s case i s r a d i a t i o n ; t h e t i m e r e q u i r e d t o reach a given temperature i s t y p i c a l l y on t h e order of a few seconds. This method has t h e advantage of r e q u i r i n g low beam power. However, i n c e r t a i n cases i t i s d e s i r a b l e c o n f i n e t h e h e a t i n g t o t h e sample s u r f a c e . The h e a t s i n k mode must then be employed i n which t h e backsurface The p r i n c i of t h e sample i s kept a t a s p e c i f i e d temperature To. p a l h e a t l o s s mechanism under t h e s e c o n d i t i o n s i s conduction through t h e wafer. C a l c u l a t i o n s of sample temperature vs beam power f o r both t h e i s o t h e r m a l and h e a t sunk modes of o p e r a t i o n , i n c l u d i n g t h e estimated s p e c t r a l response of t h e arc s o u r c e , a r e Here w e d i s c u s s t h e b a s i c experimental p r e s e n t e d i n Chapter 2. r e s u l t s o b t a i n e d w i t h such systems. 1.3.1
Rapid Annealing of
S i l i c o n With a Scanning cw Hg Lamp
A p r a c t i c a l " l a r g e diameter" scanned annealing system has been b u i l t by S t u l t z e t a l . 11.521 u s i n g a 3" long mercury a r c lamp w i t h an e l l i p t i c a l r e f l e c t o r . The l i g h t from t h e lamp i s focused i n t o a narrow (<5mm wide) ribbon by a 4" long r e f l e c t o r w i t h an e l l i p t i c a l c r o s s s e c t i o n . By u s i n g t h i s shape of r e f l e c t o r t o c o l l e c t and focus t h e l i g h t , a n i n t e n s e l i n e a r h e a t s o u r c e can be c r e a t e d while s t i l l maintaining a reasonable working d i s t a n c e between t h e sample and t h e lamp. The r e f l e c t o r has This major and minor axes of 1 0 cm and 8.6 cm, r e s p e c t i v e l y . c o n f i g u r a t i o n g i v e s a magnification f a c t o r of about 3.2 and a working d i s t a n c e of about 26 mm between t h e edge of t h e r e f l e c t o r and t h e f o c a l plane. The wafers are placed on a h e a t e d sample s t a g e , which can be t r a n s l a t e d beneath t h e lamp a t speeds up t o 1 0 cm/s. A schematic r e p r e s e n t a t i o n of t h e a r c lamp a n n e a l i n g system i s shown i n Fig. A v a r i e t y of a r c sources can be used, though a mercury 1.26. lamp i s recommended because t h e s p e c t r a l d i s t r i b u t i o n of t h i s lamp i s most h e a v i l y weighted i n t h e uv range. Consequently, t h e mean a b s o r p t i o n depth f o r t h i s source i s shallower t h a n t h a t f o r t h e o t h e r lamps. S p e c i f i c a l l y , t h e f r a c t i o n of l i g h t absorbed versus depth i s given by: F = I(A)[l-R(A)I(l
-
exp[-a(A)xl
(1.15)
J . F. GIBBONS
42
/-ELLlPITlCAL
REFLECTOR
HEATED TRANSLA T l O N STAGE ( 0 10 cm/s 1 +
FOCUSED LINEAR
’\-SILICON
/ BEAM J
WA FER
FIGURE 1.26. Schematic r e p r e s e n t a t i o n of scanning cw arc lamp a n n e a l i n g system. where I ( A ) i s t h e wavelength dependent r a d i a t i o n i n t e n s i t y ; R(A) i s t h e wavelength dependent r e f l e c t i v e l y ; a(A> i s t h e wavelength dependent a b s o r p t i o n c o e f f i c i e n t .
Using published a b s o r p t i o n and r e f l e c t i v i t y d a t a f o r s i l i c o n and t h e s p e c t r a l d i s t r i b u t i o n s f o r xenon, krypton and mercury lamps, t h e p e r c e n t of i n c i d e n t l i g h t absorbed v e r s u s d e p t h i n s i l i c o n These d a t a a r e given i n was c a l c u l a t e d f o r each of t h e lamps. I n t h i n f i l m p r o c e s s e s such as a n n e a l i n g and r e c r y s Fig. 1.27. t a l l i z a t i o n , t h e region of i n t e r e s t i s g e n e r a l l y less than 1 wrn t h i c k . As shown i n Fig. 1.27, t h e f r a c t i o n of l i g h t absorbed from a mercury lamp i n t h i s r e g i o n i s more t h a n t w i c e t h a t f o r t h e o t h e r lamps. 100,
_ _ _ _ MERCURY
,
I
4
I0
XENON KRYPTOk
%K
4 2
70
50
50
I-
z 43
8 3 0 K
O
, b
:
2
3
1
5 DEPTH b m )
i
+
8
FIGURE 1.27. Percent of i n c f d e n t l i g h t absorbed v e r s u s depth i n t o s i l i c o n f o r t h r e e d i f f e r e n t sources.
1. BEAM PROCESSING OF SILICON
43
FIGURE 1.28. Photograph of a p a r t i a l l y annealed s i l i c o n wafer, and TEN d i f f r a c t i o n p a t t e r n s from t h e annealed and unann e a l e d regions. e l 9 8 3 E l s e v i e r Science Publishing Co. Annealing experiments have been c a r r i e d out u s i n g 2 and 3 i n c h 1-8 R-cm p-type (100) s i l i c o n wafers which were implanted a t lOOkeV t o 1 x 1015cm-2. Several combinations of w i t h 75As' a r c lamp power, scan rate and s u b s t r a t e temperature can be used t o achieve a wafer temperature s u f f i c i e n t f o r good a n n e a l i n g , In g e n e r a l , scan r a t e s from 5mm/s t o lcm/s, s u b s t r a t e -1000°C. temperatures of 400°C t o 600°C and a r c lamp i n p u t power of 1Kw/ i n c h t o 1.5Kw/ i n c h g i v e good r e s u l t s . Because of t h e l e n g t h of t h e arc lamp, an e n t i r e wafer can be annealed i n a s i n g l e scan. Figure 1.28 i s a photograph of a 3" wafer i n which t h e s c a n was stopped a t t h e c e n t e r of t h e wafer. Because of t h e d i f f e r e n c e i n r e f l e c t i v i t i e s between s i n g l e c r y s t a l and amorphous s i l i c o n , t h e demarcation between t h e annealed and unannealed r e g i o n s i s r e a d i l y seen. Transmission e l e c t r o n microscope ( E M ) d i f f r a c t i o n p a t t e r n s made from s a m p l e s taken from t h e unannealed and annealed r e g i o n s c l e a r l y i n d i c a t e t h e amorphous and completely r e c r y s t a l l i z e d n a t u r e of t h e r e g i o n s , r e s p e c t i v e l y . (The p i c t u r e i s somewhat d i s t o r t e d due t o t h e a n g l e a t which t h e photo was taken.) The r e c r y s t a l l i z e d l a y e r was a l s o analyzed by Rutherford B a c k s c a t t e r i n g Spectroscopy (RBS) and d i r e c t l y compared t o a
J. F. GIBBONS
44
*+ a-Si
c-si
2 . 2 MeV
1
0.5
AS IMPLANTED
< u vr Y
V
<
m
0.5
1.0
ENERGY (MeV)
FIGURE 1.29. Rutherford backscattering spectra from asimplanted, arc lamp annealed and thermally annealed samples. 01983 E l s e v i e r S c i e n c e P u b l i s h i n g Co.
sample which was thermally annealed at 850°C for 30 minutes followed by a 10 second 1000°C anneal. Figure 1.29 shows the backscattering spectra from an as-implanted, arc lamp annealed and thermally annealed sample. A s shown, there is no detectable difference between the quality of the regrown material annealed by the scanned arc lamp and the thermally processed sample. Carrier concentration and mobility of the scanning arc lamp annealed material as a function of depth were determined by means of differential sheet resistivity and H a l l effect together with an anodic oxidation stripping technique. These results are shown in Fig. 1 . 3 0 , along with the as-implanted profile calculated using LSS theory and published bulk silicon mobilities for the respective impurity concentrations. A s shown, no measurable dopant redistribution occurred during the annealing process and the free carrier mobility in the annealed region i s as good as found i n bulk silicon. Finally, to assess the uniformity of the anneal, four point probe sheet resistivity measurements were made on a wafer which was completely annealed in a single scan. The variation across the wafer from side to side and top to bottom (with respect to the scanning beam) showed no significant variation i n either direction, indicating a very uniform annealing of the implanted wafer. These data are shown in Fig. 1 . 3 1 .
1. BEAM PROCESSING OF SILICON
45
-LSS THEORETICAL DISTRIBUTION
0 CARRIER CONCENTRATION
DEPTH (nm)
FIGURE 1.30.
Carrier concentration and mobility versus depth from an arc lamp annealed sample. O1983 Elsevier Science Publishing Co.
I
c:
140-
-
130-
1 PERPENDICULAR TO SCAh II PARALLEL TO SCAN
1
2- 1 2 0 u)
1
w fY 110-
I l l
I l l 1
Ill
IW
w
1100-
u)
901 1
I
I
I
0.5
0
0.5
DISTANCE FROM CENTER OF WAFER ( i n )
Sheet resistivity profiles across an arc lamp FIGURE 1.31. annealed wafer. 01983 E l s e v i e r Science Publishing CO.
J. E GIBBONS
46
The s c a n n i n g mercury arc lamp t h u s p r o v i d e s a t o o l f o r a n n e a l i n g t h a t i s e q u i v a l e n t t o t h e cw l a s e r i n r e s u l t s and h a s t h e added f e a t u r e t h a t t h e w a f e r throughput c a p a b i l i t y can be v e r y high.
1.4
CW BEAM RECRYSTALLIZATION OF POLYSILICON FILMS
H e a v i l y doped p o l y c r y s t a l l i n e s i l i c o n ( p o l y s i l i c o n ) i s a m a t e r i a l t h a t i s w i d e l y used i n p r e s e n t day s i l i c o n i n t e g r a t e d c i r c u i t technology f o r g a t e s and i n t e r c o n n e c t i o n l i n e s i n MOS integrated circuits. I n i t i a l i n t e r e s t i n t h e beam a n n e a l i n g of p o l y s i l i c o n a r o s e because of i t s p o t e n t i a l a s a p r o c e s s t h a t could be used t o promote g r a i n growth and hence o b t a i n r e s i s t i v i t y red u c t i o n i n h e a v i l y doped f i l m s . Experiments u n d e r t a k e n t o e x p l o r e t h a t p o t e n t i a l are d e s c r i b e d i n Chapter 5. These experiments showed t h a t s i g n i f i c a n t g r a i n growth can be produced when a t h i n p o l y s i l i c o n f i l m i s i r r a d i a t e d by a s c a n n i n g h e a t s o u r c e under c o n d i t i o n s t h a t produce m e l t i n g t h r o u g h o u t t h e f i l m . The m a j o r i t y c a r r i e r e l e c t r o n i c p r o p e r t i e s of t h e s e r e c r y s t a l l i z e d f i l m s c l o s e l y approximate t h o s e of bulk s i l i c o n . This r e s u l t has l e d t o s u c c e s s f u l a t t e m p t s t o f a b r i c a t e MOS t r a n s i s t o r s and i n t e g r a t e d c i r c u i t s d i r e c t l y i n t h e beam r e c r y s t a l l i z e d m a t e r i a l . A s u b s t a n t i a l i n t e r e s t now e x i s t s i n t h e p o t e n t i a l of beam r e c r y s t a l l i z e d s i l i c o n - o n - i n s u l a t o r s (SOI) as a s u b s t r a t e f o r i n t e g r a t e d c i r c u i t fabrication. Both t h e r e c r y s t a l l i z a t i o n p r o c e s s and i t s a p p l i c a t i o n t o d e v i c e f a b r i c a t i o n are d i s c u s s e d a t l e n g t h i n Chapter 5. We review h e r e o n l y t h e h i g h l i g h t s of t h e s u b j e c t .
1.4.1
Sample P r e p a r a t i o n
The p o l y s i l i c o n f i l m s u s e d i n t h i s work are t y p i c a l l y 0.50.6um t h i c k f i l m s d e p o s i t e d by low p r e s s u r e chemical v a p o r depos i t i o n (LPCVD) on a n o x i d i z e d s i l i c o n s u b s t r a t e . The g r a i n s i z e i n t h e a s - d e p o s i t e d f i l m s i s -500A. These f i l m s are t h e n capped w i t h an S i O p l a y e r f o l l o w e d by an Si3N4 l a y e r t o p r o v i d e good s u r f a c e morphology i n t h e r e c r y s t a l l i z e d f i l m . Details on encapsul a t i o n methods and r e s u l t s a r e p r o v i d e d i n S e c t i o n s 5.1 and 5.4 and the r e l a t e d references. Beam p r o c e s s i n g can be performed w i t h any s c a n n i n g s y s t e m t h a t i s c a p a b l e of d e l i v e r i n g s u f f i c i e n t e n e r g y t o t h e f i l m t o m e l t it entirely. I n t h e s i m p l e s t t y p e of laser p r o c e s s i n g , a cw argon laser i s o p e r a t e d i n t h e m u l t i l i n e mode w i t h a l a s e r beamwidth on t h e sample s u r f a c e of -40pm. The l a s e r s p o t i s scanned a c r o s s t h e sample a t a r a t e of 1-10 cm/sec. The sample
1. BEAM PROCESSING OF SILICON
47
FIGURE 1.32. Illustrating the laser recrystallization of thin polysilicon films. is usually held at 35OOC during this operation, with the laser power adjusted to produce complete melting. A typical value of laser power under these conditions is 10-12 watts. Surface melting is initiated at 5-6 watts, after which there is an increase in reflectivity as noted previously, and a consequent reduction of power absorbed in the film. The result is that only a thin surface layer is melted until the laser power is increased to 10-12 watts, where sufficient energy is coupled into the film to produce melting throughout. 1.4.2
Basic Recrystallization Process
The basic recrystallization process is suggested in Fig. 1.32. First the laser beam melts a small area of the film that is approximately the size of the laser spot. This molten region cools as the spot is scanned to the right. Crystallites having various orientations grow as the melt cools, each serving as a potential seed crystal for the material that is being melted in the adjacent laser-irradiated region. The shape of the liquidsolid interface is clearly of critical importance in determining the number of crystallites that are formed and the likely crystalline orientation of the crystallites that are grown (see Sec. 5.4 for further details). These recrystallized films are of course not normally single crystal, though single crystal islands can be grown by first etching the film into an array of islands [1.53], similar to the first step that would be taken in fabricating an SOS integrated circuit. The conditions for single crystal island growth are probably critical to the future of SOI, and a significant effort is now directed toward achieving this goal (cf Sec. 5.4). 1.4.3
Basic Device Results
Lee et al. [1.54] first showed that both enhancement and depletion mode MOSFETS could be made directly on the top surface of a laser-recrystallized polysilicon film with device properties that are surprisingly similar to devices made on single crystal substrates. Drain characteristics for both types of devices are shown in Fig. 1.33. Fabrication details are discussed in Sec. 5.3.
48
J. F. GIBBONS
FIGURE 1.33. (a) Source-drain I-V characteristics for a (b) Source-drain deep depletion-mode device (Vc = 0 to -12 V). I-V characteristics for an enhancement-mode device.
The films used by Lee et al. were recrystallized on Si3Nq/Si substrates, however, which was ultimately found to be less satisfactory than SiO2/Si substrates. Lam et al. [1.551 obtained improved characteristics by using a thermally grown SiO2-on-Si substrate rather than Si3N4, which led to less strain at the lower surface of the film; and they made 11 stage ring oscillators with a minimum propagation delay of 44 n s per stage compared to a propagation delay of 36 ns per stage for the same circuit made on a single crystal substrate.
1.4.4
Improvements in the Recrystallization Process
These first results, together with the knowledge that polysilicon islands (rather than continuous films) could be recrystallized into single crystal silicon, led to a rapid development of the field. Using a LOCOS technology, Lam and coworkers [1.56] fabricated islands that were surrounded by an oxide "moat" that helps to clamp the polysilicon island during its resolidification from the melt. This improves surface flatness and makes device
1. BEAM PROCESSING OF SILICON
49
fabrication significantly easier. Biegelsen et al. [1.57] showed that recrystallization could be improved by shaping the island prior to recrystallization into a chevron-like configuration so that the point of the chevron melts first as the beam is scanned by. This point then serves as a seed for the subsequent recrystallization of the entire island. This, together with careful attention to the heat flow problem, led to significant improvement in the crystallinity of the laser processed films and corresponding improvements in devices. Later Stultz and Gibbons [1.58] showed that recrystallization could be improved by shaping the laser beam into a tilted, half moon configuration, obtaining large crystallites that were almost entirely in their crystal orientation. An alternative to the use of lasers or arc sources for both annealing and recrystallization, the graphite strip heater, was pioneered by Fan, Geis and ther co-workers [1.59,1.60]. For recrystallization studies (Geis, et al. [1.601), a lower, large graphite strip serves as both the sample holder and a heater. It is heated to -120O0-13OO0C in typical recrystallization experiments with silicon films. A movable, linear strip heater is then mounted above the sample and scanned across it to produce the lateral thermal gradient necessary for recrystallization. The movable graphite strip is heated to a sufficient temperature to melt the silicon film as it is scanned over it.
A different approach was initiated when Kamins et al. [1.611 and Lam et al. [1.62] reported that single crystal films could be obtained by using a via through the Si02 film to permit the polysilicon to contact an underlying single crystal substrate, thus providing a single crystal seed for the subsequent laserinduced crystallization of the polysilicon that lies over the oxide (Fig. 1.34). Trimble et al. [1.63] then combined this seeded lateral eiitaxy concept with the beam shaping technique of Stultz and Gibbons to produce very promising SO1 films.
FIGURE 1.34. Geometry for laser-assisted seeded lateral epitaxy (after Trimble et al. [1.63]). @1982 Elsevier Science Publishing Co.
50
J. F. GIBBONS
Films r e c r y s t a l l i z e d by many o f t h e s e p r o c e s s e s now have s u r f a c e c a r r i e r m o b i l i t i e s t h a t are e s s e n t i a l l y e q u a l t o t h o s e of bulk s i l i c o n , and s u p e r i o r t o SOS. Furthermore t h e carrier m o b i l i t y i s found t o be uniform through t h e f i l m , i n s h a r p c o n t r a s t Arrays of r e l a t i v e l y l a r g e geometry MOSFETs have been t o SOS. made f o r use as l a r g e area LCD d i s p l a y d e c o d e r s , an a p p l i c a t i o n t h a t i s based p r i p l a r i l y on c o s t r e d u c t i o n and i s r e l a t i v e l y u n c r i t i c a l w i t h regard t o d e v i c e performance. For smaller geometry (VLSI) a p p l i c a t i o n s , t h e r e remains one b a s i c problem, however, r e l a t e d t o t h e f a c t , t h a t none of t h e r e c r y s t a l l i z a t i o n p r o c e s s e s h a s y e t succeeded i n f a b r i c a t i n g s i n g l e c r y s t a l f i l m s o r i s l a n d s every t i m e . I n s t e a d , g r a i n boundaries are u s u a l l y p r e s e n t , and dopant d i f f u s i o n a l o n g g r a i n boundaries d u r i n g c o n v e n t i o n a l s o u r c e and d r a i n f a b r i c a t i o n has been shown t o l e a d t o s o u r c e - d r a i n s h o r t s when t h e g a t e l e n g t h i s t o o small (<3pm) f1.641. This problem can be avoided by u s i n g a very s h o r t h e a t c y c l e t o a n n e a l t h e s o u r c e and d r a i n i m p l a n t s , such as i s provided by any of t h e a n n e a l i n g t e c h n i q u e s d e s c r i b e d above. I n d e e d , Ng e t a l . [1.651; see a l s o Sec. 5.4.4 of t h i s volume] have b u i l t d e v i c e s w i t h a g a t e l e n g t h of -1.5um and a 10 s t a g e r i n g o s c i l l a t o r w i t h a p r o p a g a t i o n d e l a y of 118 p s / s t a g e on non-seeded material t h a t was r e c f y s + t a l l i z e d d i r e c t l y w i t h a s c a n n i n g c i r c u l a r laser beam; i.e., t h e least f a v o r a b l e r e c r y s t a l l i z a t i o n c o n d i t i o n s . These r e s u l t s provide s i g n i f i c a n t m o t i v a t i o n f o r f u r t h e r work i n this field.
1.4.5
Three Dimensional I n t e g r a t i o n
F u r t h e r m o t i v a t i o n f o r t h i s f i e l d w a s provided by Gibbons and Lee [1.66], who demonstrated t h a t S O 1 can be used t o f a b r i c a t e v e r t i c a l l y i n t e g r a t e d (i.e., t h r e e dimensional) d e v i c e s . These a u t h o r s used t h e lower s u r f a c e of t h e r e c r y s t a l l i z e d f i l m t o make an n channel d e v i c e t h a t i s d r i v e n by t h e same g a t e t h a t c o n t r o l s a p channel device i n t h e u n d e r l y i n g s i n g l e c r y s t a l This substrate. The b a s i c s t r u c t u r e i s shown i n Fig. 1.35. s i n g l e g a t e CMOS i n v e r t - e r i s s i m i l a r i n p r i n c i p l e t o what one would o b t a i n by c o n c e p t u a l l y f o l d i n g t h e n channel d e v i c e of a c o n v e n t i o n a l CMOS i n v e r t e r over t h e p channel d e v i c e , t h u s o b t a i n i n g immediately a n e a r l y two-fold increase i n inverter packing d e n s i t y . I n a d d i t i o n , because t h e n and p channel d e v i c e s a r e on p h y s i c a l l y d i f f e r e n t s u b s t r a t e s , t h e r e i s no p o s s i b i l i t y f o r l a t c h up, a comrnon problem w i t h bulk CMOS circ u i t s . A l t e r n a t i v e p r o c e s s i n g sequences f o r s t a c k e d CMOS d e v i c e s have been developed by s e v e r a l a u t h o r s , [1.67, 1.681. More r e c e n t l y , Kawamura e t a l . [1.691 have f a b r i c a t e d t h r e e dimens i o n a l CMOS i n t e g r a t e d c i r c u i t s w i t h one type of d e v i c e f a b r i c a t e d d i r e c t l y above a t r a n s i s t o r of t h e o p p o s i t e type w i t h s e p a r a t e g a t e s and an i n s u l a t o r between t h e d e v i c e s . Seven
1.
BEAM PROCESSING OF SILICON
51
LASER-RECRYSTALLISED CVD POLYSILICON
(100)
\GATE OXIDE FOR BOTTOM LEVEL DEVICE
FIGURE 1.35. One-gate-wide CMOS inverter with devices built in both the underlying single crystal silicon substrate and a laser recrystallized polysilicon film. stage ring oscillators have been made with this form of threedimensional integration having a propagation delay of 8.2 ns per gate. 1.5
SILICIDE FORMATION WITH SCANNING BEAMS
The reaction of thin metal films with both single crystal and amorphous silicon to form silicide compounds has been successfully demonstrated using both scanned cw and pulsed beam processing. A detailed summary of the work that has been done with cw beams is given in Chapter 6 of this volume. We give an overview of some of the principal results in this section. Interest in the formation of silicide compounds arises in part from the fact that for given thickness, their sheet resistance is significantly lower than that of heavily doped polysilicon. This is of importance in silicon integrated circuit technology since the drive toward higher device packing densities and decreased dimensions will require low resistance, thin film interconnects with high temperature properties that are compatible with other silicon processing operations. Similar to the case for beam annealing of ion implanted silicon, both the basic mechanisms of silicide formation and the experimental results are found to depend on whether pulsed or scanned beams are used to promote the reactions. For the pulsed beam (Q-switched laser) case, melting and rapid solidification occurs in the metal-silicon films. The principal result is that the reacted film is frozen in a mixed phase compound which
J. F. GIBBONS
52
CW LASER BEAM
(100) Si SUBSTRATE
FIGURE 1.36. (a) Example of sample structure used for scanned-cw laser silicide formation. (b) Nomarsky micrograph of CW laser formed Pd2Si. 01981 Elsevier S c i e n c e P u b l i s h i n g Co.
normally includes several silicide compounds and silicon and metal precipitates. By contrast, the use of a scanned cw system t o react the metal-silicon films results in the formation of a large area, uniform, single phase silicide with essentially no silicon o r metal precipitation. 1.5.1
Sample Preparation
A schematic illustration of the samples that are used for c h i s work is shown in Fig. 1 . 3 6 . The metal film is electron-beam
evaporated in a high vacuum, oil-free system. Sputter deposition of the metal film is generally avoided because of residual argon impurity incorporation and its affect on both growth mechanisms a n d film properties. For cw laser processing, a thin overlayer of silicon is also electron beam evaporated in the same pump down
1.
BEAM PROCESSING OF SILICON
53
as that used for depositing the metal film. This overlayer provides an anti-reflection coating that enables efficient coupling of the laser into the film. If an electron beam is used to promote the reaction, this antireflection coating is not necessary. We also show in Fig. 1.36 a Nomarski micrograph of a silicide formed by cw laser processing. The surface of this material is essentially featureless at the magnification indicated, suggesting a very uniform reaction. 1.5.2
Growth Mechanisms
Measurements of the kinetics of silicide formation via scanned laser processing have been carried out for several metal/silicon systems. In Fig. 1.37 we show Rutherford backscattering data, taken from Shibata, et al. [1.70] obtained on cw laser reacted Si/Nb/Si structures. The substrate temperature during laser reaction for this example was 35OoC. The advancement of the interface between the silicon and niobium materials, indicative of silicide formation, occurs for constant laser power as a function of the number of laser scan frames. It can be seen that the thickness of the silicide film (NbSi2) increases with the number of scan frames. The average composition of this reacted layer was found to be NbSI2, as determined from backscattering and verified by glancing angle X-ray diffraction analysis. I
I
I
I
I
I
I
ENERGY (MeV)
FIGURE 1.37. Rutherford backscattering spectra taken on laser reacted Nb/Si. The formation of NbSip as a function of number of laser scan frames is shown. O1981 E l s e v i e r Science Publishing CO.
J. E GIBBONS
54 lo4,
F
,
I
I
NORMALIZED LASER POWER. p: 0 9
NUMBER OF SCANS
FIGURE 1.38. P l o t of t h e i n c r e a s e of NbSi2 t h i c k n e s s v e r s u s t h e number of l a s e r scan frames. 01981 E l s e v i e r Science Publishing Co. Reprinted by permission of the publisher, the Electrochemical S o c i e t y , Ioc.
Using t h e n e a r s u r f a c e S f a c t o r , [Sp]:tsi2 = 83.7 e V / i , the f i l m t h i c k n e s s v e r s u s number of scan rames was determined and plotted. This r e s u l t i s shown i n Fig. 1.38. One can s e e t h a t a p s r a b o l i c growth behavior e x i s t s f o r t h e NbSi2 f i l m s i n c e t h e number of s c a n s i s p r o p o r t i o n a l t o t h e e f f e c t i v e a n n e a l i n g t i m e . This r e s u l t i s i n agreement w i t h t h e g e n e r a l behavior of t h e r e f r a c t o r y metals which form d i s i l i c i d e s by f u r n a c e annealing. These f a c t s t h u s s u g g e s t t h a t t h e laser induced r e a c t i o n f o r t h e NbSi2 f o r m a t i o n p r o c e s s i s v i a a s o l i d phase d i f f u s i o n mechanism. T h e o r e t i c a l models p r e s e n t e d i n Chapter 2 of t h i s volume can be used t o compute t h e e f f e c t i v e s u r f a c e temperature and e f f e c t i v e a n n e a l i n g t i m e f o r given scanning c o n d i t i o n s (laser power, scan v e l o c i t y , beam r a d i u s and sample r e f l e c t i v i t y ) . These d a t a can be used t o g e t h e r w i t h measured f i l m t h i c k n e s s growth t o s t u d y t h e growth mechanism and o b t a i n k i n e t i c d a t a t o c h a r a c t e r i z e t h e process. We show such a development i n Fig. 1.39. On t h e r i g h t hand s i d e of t h i s f i g u r e t h e v a r i o u s normalized l a s e r powers used t o o b t a i n a c e r t a i n t h i c k n e s s of growth of NbSi2 a r e used t o c a l c u l a t e t h e maximum s u r f a c e temperature f o r a s u b s t r a t e t e m p e r a t u r e of 350°C. The r e c i p r o c a l of t h i s temperature i s t h e n p l o t t e d v s t h e s q u a r e of t h e r e a c t e d l a y e r t h i c k n e s s t o determine t h e growth kine t i c s of t h e material. T h i s p l o t is shown on t h e l e f t s i d e of F i g . 1.39. It can be seen t h a t t h e d a t a are w e l l f i t by a s t r a i g h t l i n e i n d i c a t i n g t h a t a f i r s t o r d e r p r o c e s s is o c c u r r i n g . A least s q u a r e s f i t of t h e d a t a r e s u l t s i n an a c t i v a t i o n energy f o r t h e
1. BEAM PROCESSING OF SILICON
i
-
20 LASER SCANS teff 7 msec Ea = 1.47eV 'Qb.0 7.5 8.0 tO"/T
55
a:
L 10
(OK-')
FIGURE 1.39 Comparison of theoretical calculations with experimental measurements for Pd2Si formation by scanned cw laser. process. By using an iterative procedure that takes into account that the effective time is dependent on both the activation energy and maximum surface temperature, Shibata, et al. were able to determine a value o f activation energy of 1.47 eV. The rate constant for the process was found to be A = 1.08 x 1014 i2/sec. The formation and analysis of MoSi2 and WSi2, silicides of special importance for integrated circuit applications, are discussed in detail in Chapter 6 of this volume, along with attempts that have been made to form superconducting silicides by carrying out metal/silicon reactions at temperatures that are difficult to achieve without a laser system. Important results on silicide oxidation mechanisms and kinetics are also presented in that chapter. 1.6
RAPID THERMAL PROCESSING WITH STATIONARY INCOHERENT SOURCES
To conclude this chapter, we describe a simple form of cw beam processing in which annealing of an entire wafer may be achieved in seconds using a uniform source of incoherent radiation. As with the scanning Hg arc and graphite strip heater systems discussed previously, the high throughput capabilities of these systems make them attractive for application in practical cal semiconductor processing lines.
J. F. GIBBONS
56
T h e s e systems a r e u s u a l l y c a l l e d " r a p i d t h e r m a l p r o c e s s i n g systems" because they are r a p i d compared t o a c o n v e n t i o n a l f u r n a c e processing. The b a s i c p h y s i c s of t h e p r o c e s s i s v e r y s i m i l a r t o f u r n a c e a n n e a l i n g , however.
Rapid thermal p r o c e s s i n g systems employ r a p i d l y cycled banks of tungsten-halogen lamps, a r c s o u r c e s o r e l e c t r i c a l l y heated g r a p h i t e s t r i p s t o raise t h e t e m p e r a t u r e of a t h e r m a l l y i s o l a t e d This t e m p e r a t u r e i s wafer t o 100O0-12OO0C i n a few seconds. maintained f o r a few seconds, a f t e r which t h e wafer i s allowed C a l c u l a t i o n s of t h e t o c o o l by r a d i a t i o n o r f o r c e d a i r c o o l i n g . temperature-time p r o f i l e s f o r such systems are g i v e n i n Chapter 2. The f i r s t system of t h i s t y p e used f o r a n n e a l i n g i o n implanted s i l i c o n was r e p o r t e d by Nishiyama e t a l . [1.71], who used A schematic i l l u s t r a t i o n of a somewhat quartz-halogen lamps. l a r g e r system, used by Lischner and Celler [1.83] f o r b o t h anneali n g s e v e r a l w a f e r s a t a t i m e and r e c r y s t a l l i z i n g t h i c k f i l m s of s i l i c o n on i n s u l a t i n g s u b s t r a t e s , i s shown i n Fig. 1.40, along w i t h a p l o t of lamp power and wafer temperature. The r a d i a t i o n f i e l d i n t h i s system i s found t o be uniform over an area of -500cm'. Wafers may be h e a t e d t o 100O0-12OO0C i n a few seconds and cooled e q u a l l y r a p i d l y . The w a f e r peak t e m p e r a t u r e and a l s o t h e v a r i o u s time i n t e r v a l s and ramps may be c o n t r o l l e d w i t h g r e a t p r e c i s i o n u s i n g microprocessor-based c o n t r o l systems. The a p p a r a t u s i s a l s o energy e f f i c i e n t . For s i l i c o n i n t e g r a t e d c i r c u i t technology, t h e movement of t h e f i e l d toward VLSI appears t o p r o v i d e an o p p o r t u n i t y f o r a p p l i c a t i o n s of r a p i d thermal p r o c e s s i n g t h a t would probably n o t e x i s t
I
REFLECTOR ~
I
TUNGSTEN HALOGEN LAMPS
_nlR,00000000000000alR, I W ART 2 PLATE I I WARTZ PLATE I
I
OVEN WAFER
FIGURE 1.40 Schematic c r o s s - s e c t i o n of t h e RTA a p p a r a t u s . ( a f t e r Lischner e t a l . [1.831). el982 E l s e v i e r S c i e n c e P u b l i s h i n g
co
.
1. BEAM PROCESSING OF SILICON
57
o t h e r w i s e . This o p p o r t u n i t y can be f a i r l y judged a g a i n s t s p e c i f i c t e c h n i c a l c r i t e r i a such as t h e u n i f o r m i t y of a c t i v a t i o n of i m p l a n t e d dopants, dopant r e d i s t r i b u t i o n , wafer warpage, minimizat i o n of r e s i d u a l d e f e c t s , energy consumed d u r i n g p r o c e s s i n g , and W e t h e throughput of l a r g e diameter (100 t o 125 mm) wafers. w i l l review t h e s e p o i n t s b r i e f l y below. F u r t h e r d e t a i l s on s p e c i f i c a s p e c t s of t h e process w i l l be found i n succeeding c h a p t e r s of t h e book. 1.6.1
Annealing Requirements f o r VLSI A p p l i c a t i o n s on Large D i a meter S i l i c o n Wafers
Doping u n i f o r m i t y over a wafer and t i g h t c o n t r o l of v a r i a t i o n s i n r o u t i n e p r o d u c t i o n are key advantages of t h e i m p l a n t a t i o n p r o c e s s i n modern s i l i c o n technology. Current production i o n i m p l a n t e r s are capable of a t h r o u hput of a t l e a s t 100 w a f e r s p e r hour a t doses up t o about 2x10f5/cm2, w i t h doping u n i f o r m i t y b e t t e r t h a n 0.5% f o r b o t h 100 and 125 mm wafers. Furnace anneali n g i s capable of r e t a i n i n g t h i s u n i f o r m i t y and, e q u a l l y import a n t , keepingwafers f l a t w i t h i n t h e t o l e r a n c e s necessary f o r f i n e scale p h o t o l i t h o g r a p h i c s t e p s . However t h e t y p i c a l a n n e a l i n g c o n d i t i o n s (600OC f o r 30 minutes followed by 1000°C - l l O O ° C f o r 20 minutes) produce dopant r e d i s t r i b u t i o n t h a t makes i t d i f f i c u l t t o f a b r i c a t e t h e shallow j u n c t i o n s and small device s t r u c t u r e s r e q u i r e d f o r VLSI. F i g u r e 1.41 has been prepared t o provide a comprehensive Here w e show v i s u a l i z a t i o n of a n n e a l i n g p r o c e s s e s and r e s u l t s . a s e t of time-temperature c o o r d i n a t e s on which l i n e s (---As--1 have been drawn g i v i n g t h e time-temperature l o c i f o r given amounts Several experimental of d i f f u s i v e r e d i s t r i b u t i o n of A s (3dDt). r e s u l t s t h a t are h e l p f u l f o r j u d g i n g t h e r e l e v a n c e of r a p i d t h e r m a l p r o c e s s i n g f o r VLSI a p p l i c a t i o n s are a l s o given. It should be mentioned t h a t t h e d i f f u s i o n c o e f f i c i e n t used f o r t h i s construct i o n i s o b t a i n e d from Tsai e t a l . [1.72] and is c o n c e n t r a t i o n We have assumed an A s c o n c e n t r a t i o n of 5 ~ 1 0 ~ ~ /f oc rm ~ dependent. t h e construction.
1.6.1.1
Uniformity
The u n i f o r m i t y of a c t i v a t i o n of implanted dopants over l a r g e w a f e r s i s of paramount importance f o r any p r o c e s s t h a t i s a candid a t e f o r a p p l i c a t i o n i n s i l i c o n VLSI technology. Current and Yee [1.73], q u o t i n g a p r i v a t e communction from P e r l o f f , s t a t e t h a t f u l l a c t i v a t i o n of moderate dose (10l5/cm2) i m p l a n t s of B , P and A s c a r r i e d out at low e n e r g i e s (-50 keV) i n t o l a r g e wafers (100-125mm) can be s u c c e s s f u l l y o b t a i n e d a t temperatures of 1000°C i n t e n seconds u s i n g a halogen lamp system (open c i r c l e i n Fig.
J . F. GIBBONS
58
PROCESSING TEMPERATURE ("C) 14 0
lo4
1200
1100
1000
1
1
I
9(
lo2
iA
loo
c
VBaumgart. e t a l . (1016/cm2 B 5 0 k e ~ ; 1016/cm2 A s , 150 keV) OHodgson. et a l . ( 5 ~ 1 As, 0 ~ 50 ~ keV) O(hrrent and Yee ( 1 .Current
I
I
1
7
.
5
I015/cm2 B , P , As, 50 keV)
and Yee ( 6 ~ 1 O ~ ~ /B c , m P ,~ A s , 50 keV) i
I
8
1 0 ~( K1) ~
FIGURE 1.41 Time-temperature coordinates with experimental annealing results and estimated damge annealing locus for heavy implants. a1984 Elsevier Science Publishing Co.
1.41). This confirms the earlier work of Burton, et al. [1.74] and Baumgart, et al. [1.75], who achieve high activity with little diffusive redistribution and who also show (on smaller samples) that residual defect densities as judged by TEX can be reduced to levels that are at or below those obtaioned in furnace annealing. Figure 1.41 suggests that about 300A of diffusive redistribution should be expected to take place under these conditions. However, at higher doses, Current and Yee find that uniformity of activation is unacceptable until the annealing time is increased to 60 seconds at 1000°C (filled circle on Fig. 1.411, citing sensitivity of activation to temperature gradients across the wafer as the probable cause of this behavior. The consequent redistribution is then estimated to be near 0.1 wn, which is approaching the limit of acceptability for MOS-VLSI and is already too large for really high speed bipolar devices.
1. 1.6.1.2
BEAM PROCESSING OF SILICON
59
Wafer Flatness
Wafer flatness is an absolute necessity for both mask registration and for high yield, since warped wafers are easily broken in subsequent processing operations. Current and Yee [1.73] have studied wafer flatness after implant and rapid anneal on both 100 and 125 mm wafers. Their results show that implantation improves wafer flatness, probably by relieving surface stress; and that halogen lamp annealed wafers have flatness that is equivalent to that of furnace annealing. Table 1.7 gives a partial summary of their data. TABLE 1.7
Wafer flatness changes with implantation and anneal.
Wafer Diameter
Peak to Valley in um Before After
Mount
Anneal
100 100 100 100 100
Unclamped Unclamped Unclamped Clamped Clamped
Furnace Rapid None Furnace Rapid
5.5 4.1 4.4 4.6
100
Clamped
None
4.1
4.6
125 125 125
Unclamped Unclamped Unclamped
Furnace Rapid None
6.5 7.2 7.4
4.8 5.3 4.9
125 125 125
Clamped Clamped Clamped
Furnace Rapid None
6 .O 4.9 4.9
3.6 3.1 1.9
(mm) ~~~~
~~~
4.1
4.9 5.1 3.6 4.9
4.6
Reprinted with permission of Solid State Technology, published by Technical Publishing, a company of Dun and Bradstreet Corporation.
1.6.1.3
Residual Damage (TEM)
A potentially troublesome feature of rapid thermal annealing not discussed by Current and Yee concerns the annealing time/temperature combination necessary to achieve defect free material, as judged by TEM, after high dose implants. et al. [1.75] have shown that, for B implants of Baum%art5 101 /cm at 50 keV and As implants of 10l6/cm2 at 150 keV, annealing at 1200°C for 30 seconds is required to reduce the. damage to an acceptable level. Annealing at llOO°C for 30 seconds for these implant conditions leaves a high density of irregularly shaped dislocation loops that would almost certainly produce unacceptable device properties.
60
J. F. GIBBONS
Hodgson et al. [1.76], using an ultrahigh power, vortex cooled ar on arc lamp obtain similar data for As implants of 5 ~ 1 0 ~ ~ /atc m40~ - 50 keV. To achieve no observable defects in TEM, temperatures in the vicinity of 120OoC are required with anneal times of a few seconds. Defects are observed when the anneal temperature is reduced to 1150"C, for -3 second annealing time. To call attention to these points, we have drawn a line in Fig. 1.41 (?"""D"""~ . ) representing an estimate of the time-temperature locus that may be necessary to achieve high quality annealing for high dose implants. Hill (1.771 has suggested, based on a plot similar to that shown in Fig. 1.41, that in the far future device dimensions might be reduced to the point where only 0.01 Um of impurity redistribution would be allowed, and that annealing times of 10 milliseconds and temperatures of 1200°C would then be required, suggesting the use of a scanning cw laser. However, if the residual damage line has the shape suggested in Fig. 1 for high dose implants, higher temperatures and longer times than those suggested by Hill might be required for successful annealing. In fact, it may be difficult to achieve a simple implant1 anneal process with lOOA accuracy that is totally satisfactory. New device design procedures may simplify processing requirements, however.
1.6.1.4
Enhanced Redistribution
One of the advantages often cited for rapid cw processing is that little diffusive redistribution occurs if times and temperatures are managed properly. This conclusion has been based on work carried out primarily in the 1000°C temperature range. However, Hodgson et al. 11.761 mention that enhanced diffusion may be occurring under the experimental conditions they employ; and the recent data of Narayan et al. [1.78] suggest that the diffusion coefficient for As at 1050°C is about two orders of magnitude greater than the literature value. To illustrate this, we show in Fig. 1.42 A s concentration profiles measured by Narayan et al. before and after a 1O5O0C, 10 sec annealing cycle. The implant was carried out at 100 keV to a dose of 1016/cm2. The likelihood that the enhanced diffusion is a damageenhanced phenomena f s suggested by the work of Seidel [1.79], who implanted 2x101 /cm2 BF2 into a preamorphized Si target and found full electrical activity following a rapid thermal anneal at 77OoC for 10 seconds. SIMS profiles showed no motion of the annealed profile at temperatures up to 1000°C for 10 seconds. However, TEM studies show residual damage that may require higher temperature anneals for its removal. In any case, the possibility that enhanced diffusion is occurring during the rapid annealing
1. lo5
BEAM PROCESSING OF SILICON
61
100 keVAs IMPLANTED Si(i00) I
I
I
l
I
.... ..
I
[
I
b
g. 0.
TOTAL-ANNEALED 4050 'C FOR 5 sec SUBSTITUTIONAL TOTAL-AS IMPLANTED
.
bm
TOTAL-ANNEALED 4050 'C FOR 40 5cc 0 SUBSTITUTIONAL TOTAL-AS IMPLANTED
40'
-030
0
oto
0.20
Li
a30
DEPTH (microns)
FIGURE 1.42 Arsenic concentration profiles before and after rapid thermal processing: (a) 1O5O0C/5 sec; (b) 105OoC/10 sec.
of very shallow ion implanted layers seems to be worth attention. The future of rapid thermal annealing for very shallow structures and hence i t s appropriateness for VLSI applications will depend partly on the the answer to this question. 1.6.1.5
DLTS
A further point of interest regarding defect minimization arise from the work of Pens1 et al. r1.801. These authors find DLTS defect spectra and densities (Fig. 1.43) that appear to depend on both cooling rate and thermal gradients. If the cooling rate is indeed responsible for the excess defect density obtained, then more careful attention will have to be given to this feature of the annealing process.
There is also the possibility that, for high dose implants at least, the percolation of oxygen into the implanted layer prior to annealing (from direct exposure to the laboratory environment) may be responsible for some of the residual damage. These points may suggest that very well controlled annealing in situ (i.e., in
J. E GIBBONS
62
Epi- Silicon (100) p-Type (Boron) p = 7 flcm
2-
SEk,/'
?-
-5+
,
TPA: -
X'
Z
t = 2 0 sec fast ramp down
z 0
(5O-15O0C/~)
AG 21?,--0,
t
\
11:
/'
'. .
VARIAN '\ GRAPHITE i, HEATER I
a Q Lz
t-
L
do FIGURE 1.43.
I. 1
\
' d o '
annealing.
.
'1 1
lo&
' l&
b ' 12bo '
'
OC
'
ANNEALING TEMPERATURE T,
DLTS spectra for various types of rapid thermal
Olga6 Elsevier Science Publishing co.
the ion implantation chamber) may be required to achieve the ultimate limits of the technology. 1.6.1.6
Summary
To summarize, for doses of B , P and As below 1015/cm2, any of several forms of rapid thermal annealing (quartz-halogen lamps, arc lamps, and hot graphite strips) appear to be satisfactory for application in silicon IC technology, measured against reasonably exacting criteria. The adequacy of the technique f o r very shallow structures and higher dose cases is less clear if redistribution is to be strictly limited, though it is adequate f o r the near future. In addition there are device structures and alternative fabrication techniques to be discussed later that may also create a role for rapid annealing.
1.6.2
Applications of Rapid Annealing for Si Microwave Bipolar Transistors
It is clear from recent progress in the microwave semiconductor industry that very shallow bipolar structures can be made without putting the entire burden of profile control on low
1 . BEAM PROCESSING OF SILICON
63
/"
pO'y-si
FIGURE 1.44. Super self-aligned bipolar transistor (Sakai et al. [1.81]. 01984 Elsevier Science Publishing C o . voltage ion implantation. A s an example we show in Fig. 1.44 a bipolar transistor that uses a "double poly" process to form both low resistivity base contacts and a polysilicon "extended emitter" [1.81]. The active base region is formed by direct ion implantation and furnace annealing. Rapid thermal annealing could perhaps be used to some advantage here, though base widths much less than 0.1 pm are not as useful as might appear at first sight because of the undesirably large base resistance (rb) that is then obtained. The advantage of rapid thermal processing would be primarily one of convenience. The emitter in this structure is prepared by i o n implantation of A s into the polysilicon "extended emitter", followed by furnace annealing to drive the As into the underlying single crystal silicon. The final device has an fT of -12GHZ and has been used to make ECL-type counting circuits that operate at 5 G W clock rates and dissipate -0.5 mW per gate. These are among the fastest IC building blocks ever made. Attempts to use rapid annealing to fabricate a similar bipolar structure have been made by Natsuaki, et al. r1.821, who have concentrated on using the process both to activate a 1 . 5 ~ 1 0 ~ ~ As implant into the polysilicon extended emitter and to control A s penetration into the underlying silicon. The "rapid anneal" was accomplished by simply mounting wafers on a quartz boat of low thermal mass, inserting them rapidly into a furnace, and then removing them quickly. Grain growth in the polysilicon and As diffusion into the underlying crystalline siliconwere measured as a function of time at an annealing temperature of 115OOC. Figure 1.45 shows the cut off frequency fT obtained as a function of collector current. The data is similar to that of Sakai et al. [1.81] described earlier, who employ furnace annealing. We also show fT vs collector current for a state-of-the-art double-implanted, furnace-annealed
J. F. GIBBONS
64 12 /
state-of-the-art
/
10 anneal
-r
aGeneral purpose
N
2
6-
c
furnace anneal
'c
I
0
I
1
I
I
30 40 COLLECTOR CURRENT (mA) 10
20
I
50
*
FIGURE 1.45. Dependence of fT on collector current for microwave bipolar silicon transistors. O1984 Elsevier Science Publishing Co.
bipolar transistor, and a similar curve for a general purpose microwave bipolar transistor, also double-implanted but with less attention given to ultimate frequency performance. The rapid anneal looks promising but there is need for improvement before it replaces existing technology. It should also be mentioned that SIPOS polysilicon emitter structures are very likely to become the dominant technology for emitter fabrication in bipolar transistors intended for very high speed, high gain operation. The structure is shown in Fig. 1.46 and has an As-doped polysilicon extended emitter with a thin oxygen doped layer at the base-emitter junction. The absence of significant As penetration through this oxide layer gives very low emitter side wall capacitance. In addition, the asymmetry in electron and hole injection through this oxygendoped layer results in very large emitter efficiency, leading to gain enhancements of more than an order of magnitude. A s pointed out by Hill [1.77], this structure is such a step forward from conventional bipolar transistors that it will undoubtedly be the natural successor of the current microwave bipolar transistor. From the processing viewpoint, rapid thermal annealing may be useful in redistributing the implanted A s in the polysilicon layer and simultaneously activating the B base implant, though this structure will be more forgiving than a standard bipolar
1 . BEAM PROCESSING OF SILICON
65
E
g 20-
0
c
4
‘190 I-
4
(L
t
z w
18-
z
0
0 17 -
I2
\
\B
Bipolar transistor with SIPOS emitter (after Hill [1.77]). 01983 Elsevier Science Publishing Co.
i
U
9160
I: \
V
FIGURE 1.46
I
I I
0
150-4 OEPTH (micron)
transistor of longer thermal cycles, so rapid annealing would be of importance primarily for convenience and possibly energy savings, both of which may be of considerable appeal. 1.6.3
Rapid Thermal Processing of Compound Semiconductors
Several rapid thermal processing techniques have been applied to GaAs with encouraging initial results and also important limitations. This subject has been reviewed by Williams [1.83] with the general conclusion that rapid thermal processing has not yet realized its initial promise of overcoming long standing problems with conventional furnace annealing for activation of implanted dopants in GaAs. Work done since the review of Williams has produced a somewhat more hopeful outlook, though problems still remain. This work will be reviewed in Chapter 7. Apart from its use for standard annealing experiments, however, there is also now an increasing interest in more novel applications in which rapid thermal processing is used to study and control impurity diffusion in compound semiconductors and also to control surface properties without using ion implantation. We also review these applications in Chapter 7.
J. F. GIBBONS
66 REFERENCES
1.1 1.2 1.3
1.4 1.5
1.6 1.7 1.8
1.9 1.10
1.11
1.12 1.13
1.14 1.15 1.16 1.17 1.18 1.19 1.20
of High Power Laser Radi t ion. " Ready, J. F., "Effect Academic P r e s s , 1971. K h a i b u l l i n , I. B., Shtyrkov, E. I., Zaripov, M. M., Galya n t d i n o v , M. F., and Zakirov. G. G., Sov. Phys. Semicond u c t o r s 11, 190 (1977). B v a r e c h e n s k i i , A., SmirAntonenko, A., Gerasimenko, N.N., nov, L. A., and T s e i t l i n , G. M., Sov. Phys. Semiconductors 10, 81 (1976). Nadaev, E. V., Khodyachikh, A. V., and Kachurin, G. A., Kovaleva, L. A., Sov. Phys. Semiconductors 10, 1128 (1976). Baeri, P., Campisano, S. U., F o t i , G., and Rimini, E., Appl. Phys. L e t t . 33, 2, 137 (1978). M u l l e r , J. C., Grob, A., Stuck, R., and S i f f e r t , P., Appl. Phys. Lett. 33, 2 , 137 (1978). and A f f o l t e r , K., Appl. Phys. von Allmen, M., Luthy, W., Lett. 33, 9 , 824 (1978). Celler, G. K., P o a t e , M. J., and Kirnerling, L. C., Appl. Phys. Lett. 32, 8, 464 (1978). White, C. W., C h r i s t i e , W. H., Appleton, B. R., Wilson, S. R., Pronko, P. P., and Magee, C. W., Appl. Phys. L e t t . 33, 7 , 662 (1978). Venkatesan, T. N. C., S l u s h e r , Auston, D. H., Surko, C.M., R. E., and Golovchenko, J . A , , Appl. Phys. L e t t . 33, 5, 437 (1978). K i r k p a t r i c k , A. R., Minnucci, J . A., and Greenwald, A. C., I E E E Trans. on Elec. Dev. ED-24, 429 (1977). G a t , A , , and Gibbons, J. F., Appl. Phys. Lett. 32, 3 , 142 (1978). Gat, A., Gibbons, J . F., Magee, T. J., Peng, J., W i l l i a m s , P., D e l i n e , V., and Evans, C. A. Jr., Appl. Phys. L e t t . 33, 5 , 389 (1978). Gibbons, J. F., L i e t o i l a , A., Sigmon, R e g o l i n i , J. L., T. W. Magee, T. J., Peng, J., Hong, J. D., K a t z , W., and Evans, C. A., Jr., Appl. Phys. Lett. 50, 6, 4388 (1979). Gat, A., Gerzberg, L., Gibbons, J. F., Magee, T. J., Peng, J., and Hong, J. D., Appl. Phys. L e t t . , 33, 8, 775 (1978). Lee, K. F., Gibbons, J. F., Saraswat, S., Kamins, T. I., Appl. Phys. Lett. 35, 2 , 173 (1979). Sheng, T. T., and Celler, G. P o a t e , J . M., Leamy, H. J., K., Appl. Phys. L e t t . 33, 11, 918 (1978). Appl. Phys. L e t t . 34, 3 , Von Allmen, M. and W i t t m e r , M., 221 (1979). Tsaur, B. Y., and Magee, J. W., Appl. Phys. Liau, 2. L., L e t t . 34, 3, 221 (1979) S h i b a t a , T., Gibbons, J.F., and Sigmon, T. W., Appl. Phys. L e t t . 36, 7, 566 (1980).
1. 1.21
1.22 1.23 1.24 1.25 1.26 1.27 1.28 1.29 1.30
1.31 1.32 1.33 1.34
1.35 1.36 1.37
1.38 1.39 1.40
1.41 1.42
BEAM PROCESSING OF SILICON
67
Auston, D. H., Galovchenko, J. A., Simons, A. L., S l u s h e r , R. E., Smith, P. R. and Surko, C. M., "Laser-Solid I n t e r a c t i o n s and Laser P r o c e s s i n g - 1978," (S. D. F e r r i s , H. J. Leamy, and J. M. Poate, eds.) pp. 11-16. AIP (1979). Bloembergen, N., op c i t , p. 1-9. "Thermophysical P r o p e r t i e s of Matter," Vol. 4, p. 204. IFI/Plenum, New York (1970). Ho, C. Y., Powell, R. W. and L i l e y , P. E., J. Phys. Chem. Ref. Data 3, 1, 1-588 (1974). P h i l l i p , H. R. and T a f t , E. A., Phys. Rev. 120, 37 (1960). Brodsky, M. H., T i t l e , R.S., Weiser, K. and P e t t i t , G. D., Phys. Rev. 1, 6, 2632 (1970). Dash, W.C. and Newman, R., Phys. Rev. 99, 4, 1151 (1951). L e i t o i l a , A. and Gibbons, J. F., J. Appl. Phys. 53, 4 (1982). and Gibbons, 3. F., Appl. Phys. Lett. 40, 7 L e i t o i l a , A. ( 1982). Van Vechten, J. A., Tsu, R. S a r i s , F. W. and Hoonbout, D., Phys. L e t t . 74A, 417 (1979). Kurz, H. , and Bloembergen, N. , Appl. Liu, J. M., Yen, R., Phys. L e t t . 39, 755 (1981). Baeri, P., Campisano, S. V., F o t i , G. and Rimini, E., J. Appl. Phys. 50, 2, 788 (1979). Wood, R. F. and Pronko, P. P., Appl. Phys. Wang, J. C., L e t t . 33, 5, 455 (1978). White, C. W., Clark, G. J., Narayan, J., Young, R. T., C h r i s t i e , W. H., Murakami, M., King, P. W. and K r a m e r , S. D., Appl. Phys. L e t t . 32, 3, 139 (1978). Kodera, H., Jpn. J. Appl. Phys. 2, 212, (1965). Wilson, S. R., Appleton, B. R. and Young, White, C. W., F. W., Jr., J. Appl. Phys. 51, 1, 738 (1980). "Laser and E l e c t r o n Beam Narayan, J. and White, C. W., P r o c e s s i n g of Materials", p. 65-70, (C. W. White and P. S. Peercy, eds.). Academic P r e s s (1980). See a l s o Narayan, J., Young, R. T. and White, C. W., J. Appl. Phys. 49, 3912, ( 1978). Lang, D. V., J. Appl. Phys. 45, 3023 (1974). Johnson, N. M., B a r t e l i n k , D. J., Gold, R. B., and Gibbons, J. F., J . Appl. Phys. 50, 7, 4828 (1979). Burton, J. L., Kimerling, L. C., Miller, G. L., Robinson, D. A. H. and Celler, G. K., "Laser-Solid I n t e r a c t i o n s and Laser P r o c e s s i n g - 1978," (S. C. F e r r i s , H. J . Leamy, and J. M. P o a t e , eds.) p. 543. AIP (1979). K i r k p a t r i c k , A. R., Minnucci, J. A., and Greenwald, A. C., IEEE Trans. on Elec. Dev. ED-4, 429 (1977). Kennedy, E. F., Lau, S. S., Golecki, I., Mayer, J. W., Tseng, W., Minnucci, J . A., and K i r k p a t r i c k , A. R., "LaserS o l i d I n t e r a c t i o n s and Laser P r o c e s s i n g - 1978," (S. C. F e r r i s , H. J. Leamy, and J. M. P o a t e , eds.) p. 470. AIP (1979).
68 1.43 1.44
1.45
1.46 1.47 1.48
1.49
1.50 1.51
1.52
1.53 1.54 1.55 1.56 1.57
1.58 1.59
J . F. GIBBONS
Cespregi, L., Mayer, J. W. and Sigmon, T. W., Phys. Lett. 54A, 157 (1975). Bicknell, R. W. and Allen, R. M., Proceedings of the First International Conference on Ion Implantation in Semiconductors, Thousand Oaks, Calif. Gordon and Breach, New York (1970). Chu, A. and Gibbons, J. F., Proceedings 1976 International Conference on Ion Implantation in Semiconductors, (F. Churno, J. Borders, and D. K. Price, eds.) p. 711. Plenum Press, New York (1977). Gibbons, J. F., "Ion Implantation in Semiconductors - Part 11: Damage Production and Annealing," Proc. IEEE 60, 9, 1062 (1972). Stultz, T., Ph.D. Dissertation, Stanford University, Stanford, CA (1983). Nissim, Y. I., Lietoila, A., Gold, R. B. and Gibbons, J. F., J. Appl. Phys. 51, 1 (1980). Roth, J. A., Olson, G. L., Kokorowski, S . A. and Hess, L. D., "Laser and Electron-Beam Solid Interactions and Materials Processing," (J. F. Gibbons, L. D. Hess, and T. W. Sigmon, eds.) p. 413. North Holland (1981). Gibbons, J. F., Johnson, W. S. and Mylroie, S. W., "Projected Range Statistics in Semiconductors," distributed by John Wiley & Sons. Dowden, Hutchinson and Ross (1975). Johnson, N . M., Bartelink, D. J., Moyer, M. D., Gibbons, J. F., Lietoila, A . , Taknakumar, K. N., Regolini, J. L., "Laser and Electron Beam Processing of Materials," (C. W. White and P. S. Peercy, eds.), p. 423. Academic Press (1980). Stultz, T. J., Sturm, J. and Gibbons, J. F., "Laser-Solid Interactions and Transient Thermal Processing of Materials," (J. Narayan, W. L. Brown, and R. A. Lemons, eds.) p. 463. North Holland (1983). Gibbins, J. F., Lee, K. F., Magee, T. J., Peng, J. and Ormond, R., Appl. Phys. Lett. 34, 831 (1979). Lee, K. F., Gibbons, J. F., Saraswat, K. C . and Kamins, T. I., Appl. Phys. Lett. 35, 173 (1979). Lam, H. W., Tasch, A. F., Jr., Holoway, T. C., Lee, K. F. and Gibbons, J. F., IEEE Elec. Dev. Lett. EDL-1, 6 (1980). Lam, H. W., Tasch, A. F., Jr. and Holloway, T. C., IEEE Elec. Dev. Lett. EDL-1, 10 (1980). Biegelsen, D. K., Johnson, N. M., Bartelink, D. J. and Moyer, M. D., "Laser and Electron-Beam Solid Interactions and Materials Processing," (J. F. Gibbons, L. D. Hess, and T. W. Sigmon, eds.) p.487. North-Holland (1981). Stultz, T. J. and Gibbons, J. F., Appl. Phys. Lett. 39, 6 (1981). Fan, J. C. C . , Geis, M. W. and Tsaur, B.-Y., Appl. Phys. Lett. 38, 365 (1981).
1. BEAM PROCESSING OF SILICON 1.60
1.61 1.62 1.63
1.64
1.65
1.66 1.67
1.68 1.69 1.70 1.71 1.72 1.73 1.74
1.75
1.76
69
Geis, M. W., Smith, H. I., Tsaur, B.-Y., Fan, J. C. C., S i l v e r s m i t h , D. J., Mountain, R. W. and Chapman, R. L., "Laser-Solid I n t e r a c t i o n s and Transient Thermal Processi n g of Materials ( J . Narayan, W. L. Brown, and R. A. Lemons, eds.) p. 477. North Holland (1983). Karnins, T. I., Cass, T. R., Dell'Oca, C. J., Lee, K. F., Pease, R. F. W. and Gibbons, J. F., J. Electrochem. Soc. 128, 5 (1981). Lam, H. W., P i n i z z o t t o , R. F. and Tasch, A. F., Jr., J. Electrochem. Soc. 128, 9 (1981). Trimble, L. E., C e l l e r , G. K., Ng, K. K., Baumgart, H. and Leamy, H. J., "Laser and E l e c t r o n B e a m I n t e r a c t i o n s w i t h S o l i d s " , p. 505. North Holland (1982). Ng, K. K., Celler, G. K., P o v i l o n i s , E. J., F r y e , R. C., Leamy, H. J. and Sze, S . M., IEEE Trans. on Dev. L e t t . EDL-2, 12 (1981). Ng, K. K., Celler, G. K., P o v i l o n i s , E. I., Trimble, L. E. and Sze, S . M., "Laser-Solid I n t e r a c t i o n s and T r a n s i e n t Thermal P r o c e s s i n g of Materials," (J. C. C. Fan and N. M. Johnson, eds.). North Holland, 1983, i n p r e s s . Gibbons, J . F. and Lee, K. F., IEEE Elec. Dev. L e t t . EDL-1, 6 (1980). Goeloe, G. T., Maby, D. J., S i l v e r s m i t h , D. J., Mountain, R.W. and Antoniadis, D. A., IEDM Conference A b s t r a c t s , p. 554 (1981). Colinge, J. P. and Demoulin, E., IEDM Conference A b s t r a c t s , p. 557 (1981). Kawamura, S . , S a s a k i , N., I w a i , T., Kakano, M. and Takagi, M., IEEE Elec. Dev. L e t t . EDL-4, 1 0 , 366 (1983). S h i b a t a , T., Gibbons, J. F. and Sigmon, T. W., Appl. Phys. L e t t . 36, 566 (1980). Nishiyama, K., Arai, M. and Watanabe, N., Jap. J. Appl. Phys. 19, 10 (1980). T s a i , M. Y., Morehead, F. F., Baglin, J. E. E. and Michel, A. E., J. Appl. Phys. 51, 3230 (1981). S o l i d S t a t e Tech., 197 (1983). C u r r e n t , M. and Yee, A., Burton, J . L., Celler, G. K., Jacobsen, D. C., Kemerling, L. C., Lischner, D. J., Miller, G. L. and Robinson, McD., "Laser and E l e c t r o n Beam I n t e r a c t i o n s w i t h S o l i d s , " (B. R. Appleton and G. K. Celler, eds.) p. 765. North Holland (1983). Lischner, D. J., Robinson, Baumgart, H., Celler, G. K., McD. and Sheng, T. T., "Laser-Solid I n t e r a c t i o n s and Trans i e n t Thermal P r o c e s s i n g of Materials," ( J . Narayan, W. L. Brown, and R. A. Lemons, eds.) p. 349. North Holland (1983). Hodgson, R. T., Baglin, J . E. E., Michel, A. E., Mader, S . , Gelpy, J. C., "Laser-Solid I n t e r a c t i o n s and T r a n s i e n t Thermal P r o c e s s i n g of Materials," (J. Narayan, W. L. Brown, and R. A. Lemons, eds.) p. 355. North Holland (1983).
70 1.77
1.78 1.79 1.80
1.81 1.82
1.83
J. F. GIBBONS H i l l , C., "Laser-Solid I n t e r a c t i o n s and T r a n s i e n t Thermal P r o c e s s i n g of Materials," ( J . Narayan, W. L. Brown, and R. A. Lemons, eds.) p 381. North Holland (1983). Narayan, J . , Holland, 0. W., and Eby, R. E., Appl. Phys. Lett. 43, 957 (1983). IEEE Elec. Dev. Lett. EDL-4 (1983). S e i d e l , T. E., P e n s l , G., Schulz, M., S t o l z , P., Johnson, N. M., Gibbons, J . F. and Hoyt, J., "Laser-Solid I n t e r a c t i o n s and Trans i e n t Thermal P r o c e s s i n g , " ( J . C. C. Fan and N. M. Johnson, eds.) North Holland (1984) i n p r e s s . S a k a i , T., Konaka, S., Kabayashi, Y., Suzuki, M., and K a w a i , Y., Elec. L e t t . , 284 (1982). Natsuaki, N., Tamura, M., Miyazaki, T. and Yanagi, Y., Extended A b s t r a c t s of 1 5 t h Conf. on S o l i d S t a t e Devices and M a t e r i a l s , Japanese S o c i e t y f o r Appl. Phys., Tokyo (1983) p . 47-50. Lischner, D. J. and Celler, G. K., "Laser-Solid I n t e r a c t i o n s w i t h S o l i d s , " (B. R. Appleton and G. K. C e l l e r , eds.) p. 759. North Holland (1982).
CHAPTER 2
Temperature Distributions and Solid Phase Reaction Rates Produced by Scanning CW Beams Arto Lietoila*, Richard B. Goldf,James F. Gibbons, and Lee A. ChristeM STANFORD EI.E€lRONICS LABORATORIES STANFORD UNIVERSITY STANFORD, CALIFORNIA
2.5
INTRODUCTION . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . THE KIRCHHOFFTRANSFORM ......................... TEMPERATURE CALCULATIONS FOR LASERS . . . . . . . . . TEMPERATURE CALCULATIONS FOR WIDE AREA OPTICAL SOURCES: THE ARC LAMP . . . . . . . . . . . . . . . . . . . . TEMPERATURE CALCULATIONS FOR SCANNING
2.6
CW ELECTRON BEAMS . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . REACTIONRATE CALCULATIONS ........................
2.1 2.2 2.3 2.4
cw
REFERENCES
2 1
.......................................
71 72 74 89 93 10 1 105
INTRODUCTION
Quantitative studies of the effects of cw beam processing in various applications require accurate knowledge of the temperature produced by the scanning beam. By comparison to conventional furnace processing, a direct measurement of the temperature is difficult since at any one time only a small fraction of the sample is heated. As a result, theoretical calculations of the temperature produced by a scanning beam are important for choosing processing parameters which achieve a set of required experimental conditions. These temperature calculations are the subject of this chapter. The temperature is calculated by solving the heat flow equation with appropriate source terms and boundary conditions. Since the thermal conductivity of semiconductors is a strong function of temperature, this equation is strongly nonlinear, and cannot be solved analytically in the general case. However, under steady state conditions, analytic solutions to the heat flow equation can be obtained using the Kirchhoff transform. We will review this method in Section 2 . 2 below. *Present address: Micronas, Inc., 00101 Helsinki. Finland. tPresent address: Adams-Russell Company, Burlington. Massachusetts 01803 $Resent address: SEX4 Solar Corporation. Santa Clam, California 95054.
SEMICONDUCTORS AND SEMIMETALS, VOL. 17
71
Copyright Q 1984 by Academic Press. Inc. All rights of reproduction in any form reserved. ISBN 0-12-752117-8
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ART0 LIETOILA ET AL.
With regard to the geometry of the calculations, several distinct regimes are important for practical applications. When lasers or circular electron beams are used, the dimensions of the focused beam are normally small compared to the semiconductor sample thickness. In this case, three-dimensional solutions for the heat equation need to be sought. At the other extreme, typical beam "diameters" in arc lamp and some ribbon e-beam annealing systems are at least a few millimeters. Since semiconductor samples are typically only a fraction of a millimeter thick, the lateral heat flow then becomes negligible and the problem can be treated as one-dimensional. In this case one also has a choice between two different boundary conditions at the sample backsurface. 'Ihe sample may be left thermally isolated or it may be heat sunk to a fixed temperature. For very narrow ribbon beams, an intermediate regime obtains and the problem becomes two dimensional. Each of the above cases will be considered in more detail below. In addition to the calculation of the temperature, it is also important to know how scanning speed, or more precisely the beam dwell time, affects various solid state reactions promoted by the beam. In principle this problem can be solved by computing the temperature vs time for a given point of the sample during the scan. However, the lateral intensity distribution of a laser or e-beam can often be well approximated by a Gaussian distribution, and the temperature profiles are similar in shape. This makes it feasible to model the effect of any scan by imagining it to produce at each point a constant temperature, equal to the maximum temperature achieved, with a reduced dwell time. This model, to be developed in Section 2 . 6 , provides a quantitative framework for analyzing various beam-induced reactions such as solid phase epitaxial regrowth and silicide formation. 2.2
THE KIRCHHOFF TRANSFORM
The equation to be solved for the heat flow produced by the beam is the diffusion equation:
where P is the density, Cp the specific heat and K the thermal conductivity of the semiconductor; and Q is the rate of energy deposition. Equation (2.1) is nonlinear, because all three coefficients are, in general, temperature dependent. The temperature dependences of p and Cp are usually weak, however, and will be neglected.
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
73
Equation (2.1) can be simplified by introducing the Kirchhoff transform [2.1],
where 8 will be called the linear temperature and T is the true temperature. Using Eq. (2.2), the equation for 0 becomes (2.3) The nonlinearity is retained for the time-dependent case, but in the steady state, Eq. (2.3) further reduces to the linear heat flow equation
In most semiconductors, the thermal conductivity can be accurately approximated by the following expression 12.2-2.61:
Values of the parameters A and TK for different semiconductors are given in Table 2.1. Using Eqs. (2.2) and (2.5), the Kirchhoff transform can now be expressed as
The inverse of the Kirchhoff transform is T = T, + (To-TK> exp
8
The steady state temperature distributions can now be readily obtained by first solving Eq. (2.4) €or 8 , and then applying the inverse Kirchhoff transform Eq. (2.7) to obtain the true temperature. The normal boundary Conditions for T are
ART0 LIETOILA ET AL..
74 T (z=L)
=
To
(2.9)
where z i s the depth coordinate, and L is the sample thickness. TABLE 2.1. The Fitting Parameters, A and TK, for the Thermal Conductivity of Various Semiconductors [ K = A/(T-T,].
Material
A(W/cm)
si
299
GaAs InP
sic
91 115.5 430
TK("K>
99 90.9 130.5 320
Temperature Range
> > > >
300°K 250°K 300°K 500°K
Refs. 2.2,2.4,2.6 2.2,2.3,2.6 2.3, 2 . 6 2.5, 2.6
Condition (2.8) follows from the fact that the heat losses through the sample front surface (radiation, convection) are assumed to be negligible. Condition (2.9) means that the sample is in perfect thermal contact with a heat sink having a temperature Those conditions imply for 8 : To. (2.10) 0 (z=L) =
0
(2.11)
In the case of a thermally isolated sample, radiation losses dominate and these boundary conditions become invalid. This case will be discussed in Section 2.4. In Section 2.3 we develop solutions to Eq. ( 2 . 4 ) using the boundary conditions Eqs. (2.10) and (2.11). 2.3
TEMPERATURE CALCULATIONS FOR CW LASERS
The majority of scanning beam annealing experiments to date have been performed using cw lasers. We therefore begin by presenting, in this section, temperature calculations for cw laser processing. As was mentioned in Section 2.1, cw laser (and circular electron) beams are, in annealing applications, typically so tightly focussed that the beam dimensions are clearly much less than the thickness of the semiconductor wafers. Therefore, both the vertical and lateral heat flows must be considered, and the temperature calculations have to be three-dimensional. On the other hand the actual layers which are subject to the beam
2. TEMPERATURE DISTRIBUTIONS AND REACTION RATES
75
processing (e.g., implanted layers or deposited thin films) are usually less than a micron thick, and are therefore thin compared to the typical beam diameters (20-100 pm). Therefore, the vertical temperature drop across the layer to be processed is negligible, and the surface temperature can be assumed to prevail across the entire layer. (An exception to this is the use of a modified SEM for beam annealing, in which case the beam dimensions can be comparable to the thickness of the layer.) We will begin by showing how the maximum surface temperatures are obtained, and we will demonstrate that in practical applications the temperature can be assumed to be uniform in depth throughout the the layers to be processed. The effects of the wafer thickness, absorption coefficient, beam shape (elliptical, circular), and scanning speed are then studied. Finally, we will present a calibration procedure to facilitate the determination of the processing temperatures in practical annealing experiments. The Maximum Temperature Obtained with a Scanning CW Laser
2.3.1
The maximum temperature is achieved at the sample surface at the center of the beam under the following conditions:
1.
The beam is circular.
2.
The beam is stationary.
3.
The wafer is semi-infinite.
4. The absorption depth can be assumed to be zero. In most annealing experiments, these conditions are practically met because: spherical lasers are used, the scanning speeds are slow, the wafer thickness is usually several times the beam diameter, and e.g., for the Ar+ laser, the absorption depth is only about 0.7 pm. In what follows we present the temperature calculations assuming that the above conditions 1-4 prevail. The effects of deviations from these conditions are studied in later sections. The high power cw lasers used for beam annealing are usually designed to operate in the TEMoo mode, in which case the lateral beam intensity dietribution can adequately be approximated by a Gaussian distribution:
I = I~ exp (-r2/2)
(2.12)
76
ART0 LIETOILA ET AL.
where r i s t h e d i s t a n c e from t h e beam c e n t e r and w i s d e f i n e d as t h e beam r a d i u s . We w i l l f o l l o w t h i s d e f i n i t i o n of w ( l / e - p o i n t s i n i n t e n s i t y ) throughout t h i s volume. However, o t h e r d e f i n i t i o n s have a l s o been used i n t h e l i t e r a t u r e . C l i n e and Anthony [2.7], and N i s s i m e t a l . (2.21 d e f i n e t h e beam r a d i u s by t h e l/&-points i n i n t e n s i t y , which g i v e s a r a d i u s a f a c t o r of fi smaller t h a n o u r s . On t h e o t h e r hand, t h o s e doing r e s e a r c h on lasers p e r se, sometimes d e f i n e t h e beam r a d i u s by t h e l / e p o i n t i n t h e e l e c t r i c f i e l d s t r e n g t h , which corresponds t o t h e l / e 2 - p o i n t s i n i n t e n larger s i t y . This g i v e s a beam r a d i u s which i s a f a c t o r of than t h a t determined a c c o r d i n g t o t h e d e f i n i t i o n i n Eq. (2.12). The d i f f e r e n t d e f i n i t i o n s are compared i n Table 2.2. TABLE 2.2. Different Definitions Gaussian Laser Beam (TEMoo Mode)
Defining v a l u e of i n t e n s i t y
Symbols used
of
the
Relative value f o r a g i v e n beam
I/&
R, r
1/n
1/e l/e2
W
1
W
fi
for a
Beam, Radius
Refs.
2.2, 2.8, 2.17
2.6, 2.9
2.7
Assuming a Gaussian beam shape and c o n d i t i o n s 1-4, above, Lax [2.8] h a s shown t h a t t h e maximum l i n e a r s u r f a c e t e m p e r a t u r e a t t h e beam c e n t e r i s g i v e n by (2.13)
W e define the quantity within brackets absorbed power p e r u n i t r a d i u s , Pa:
Pa =
P (1-R) w
i n Eq.
(2.13)
as
the
(2.14)
.
The maximum t r u e t e m p e r a t u r e v s Pa f o r d i f f e r e n t semiconthe inverse d u c t o r s can now be o b t a i n e d u s i n g E q s . (2.13-2.14), Kirchhoff t r a n s f o r m , and t h e e x p r e s s i o n s f o r t h e t h e r m a l conducThe r e s u l t s are g i v e n i n F i g . 2.1 f o r S i , Sic, t i v i t y , Eq. ( 2 . 5 ) . GaAs and I n P u s i n g d i f f e r e n t s u b s t r a t e backsurface t e m p e r a t u r e s T o . That f i g u r e shows- t h a t u s e of h i g h e r v a l u e s of To r e s u l t s i n a s u b s t a n t i a l r e d u c t i o n of laser power f o r a g i v e n s u r f a c e t e m p e r a t u r e . This i s due t o t h e f a c t t h a t t h e thermal c o n d u c t i v i t y f o r a l l of t h e s e materials d e c r e a s e s w i t h i n c r e a s i n g t e m p e r a t u r e . It i s f o r t h e same reason t h a t t h e T v s . P curves are n o n l i n e a r and concave up.
2500
-
0,
2000
-
5
1500
-
u
a Si C
I
500
1000 Pa --
1400 -
To =550°C 35OoC
1500
2000
2500
500
.
I
I
.
1500 (1-R)P
(W/cm)
15OoC
.
1000
paZw
I
2000
.
~
2500
.
1
.
3000
(W/crn)
25OC
C InP
100
200
300
400
500
600
700
800
FIGURE 2.1. Maximum surface temperature a t the beam center as a f u n c t i o n of t h e absorbed power p e r u n i t radius f o r ( a ) s i l i c o n , (b) s i l i c o n carbide, ( c ) gallium arsenide and ( d ) indium phosphide.
78
ART0 LIETOILA ET AL.
2.3.2
Vertical and Lateral Temperature Distributions
I f the maximum temperature vs. depth must be calculated, the following formula, derived by Lax [2.8], can be used for the linear temperature at the beam center:
where Cmx(Z=O) is given by Eq. (2.13) and Z = z/w. We plot in Fig. 2.2 the ratio G(z)/Gmx, which we call the normalized linear temperature nz, as a function of the dimensionless depth coordinate Z = z/w. From Fig. 2.2(a), it is evident that in most cases the surface temperature can be used throughout the layers of interest. For example, using a typical beam radius of 40 pm, the temperature within the first 1 pm of the sample does not vary by more than 4 % .
b
x
2
-? N
Y
0
t
0.92 0901
0
"
0.01
'
1
0.02
'
z
' 003
'
1
.
1 I
0.04 0.05
0
1
2
z
3
4
5
FIGURE 2.2. Normalized linear temperature G(Z)/Gmx as a function of normalized depth Z = z/w at the beam center (x=y=O) for (a) small and (b) large values of 2. The linear temperature at the sample surface at a distance r from the beam center can be expressed L2.81 as a function of the dimensionless distance R = r/w: O(R) , =,O
(R=O) exp
(- F)I, (5)
(2.16)
where here I, is the modified Bessel function of order zero. We plot in Fig. 2.3 the ratio O(R)/Q,,(R=O) as a function of R and compare this to the normalized laser intensity distribution I(R)/Io. That figure or Eq. (2.16) together with Eqs. (2.7) and (2.13), can be used to obtain the true surface temperature as a function of distance from the beam center.
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
79
R = r/w
FIGURE 2.3. Normalized linear temperature O(R)/Omax as a function of normalized lateral distance R = r/w from the beam center at the sample surface (PO; r2=x2+y2). For comparison, the dashed line shows the profile of laser intensity. 2.3.3
Effect of Sample Thickness
The formulas given above for 0, Eqs. (2.13-2.15), were derived assuming that the sample is infinitely thick compared to the beam dimensions. This is not always the case: use of very powerful lasers and long focal length lenses may result in beam diameters in excess of 100 pm and the sample thickness must then be taken into account. A quick check of the validity of the infinite sample assumption can be performed by using Eq. (2.15) to calculate what the linear temperature rise would be at a depth equal to the actual sample thickness in an infinitely thick wafer. If the result is close to zero, no correction is required. Otherwise, the formulas given above, and Fig. 2.1, give temperatures that are too high, and an appropriate correction must be carried out. We will now present this correction for the linear temperature 8. Equation (2.7) can then be applied to obtain the true surface temperatures in wafers of finite thickness.
It has been shown using the method of images f2.91 that the linear surface temperature, k, in a sample whose thickness is L can be expressed as the following series: m
4,=
Omax [l + C (-l)k 2exp(4k2L2/w2) erfc (2kL/w)J k=l
(2.17)
ART0 LIETOILA ET AL.
80
Figure 2.4 shows the ratio of OL/&~ = r& as a function of We can see that for w/L < 0.3, the infinite sample approximation is good. On the other hand, for w/L > 2 , the temperature approaches the values obtained from one-dimensional solutions indicated by the dashed line in Fig. 2.4.
w/L.
1.oc
'-ONE \
DIMENSIONAL APPROXI MATlON
-
0" Q- P/w w 3 1/e -points in intensity
0.1L BEAM RADIUS/SAMPLE THICKNESS, w / L
FIGURE 2.4. Dependence of the linear temperature 0 on the ratio w/L of the beam radius to the sample thickness. The value of 0 is normalized to the maximum value, , O, obtained for a semiinfinite sample. The dashed line is valid for the one-dimensional approximation. (After Ref. 2 . 9 ) 01982 Elsevier Science Publishing Co.
For intermediate wafer thickness, the correct temperatures can be obtained in two ways. If the deviation of the infinite sample approximation is significant (i.e., w/L = l), one has to OL is then obtained using first calculate ,,Q using E q . ( 2 . 1 3 ) . Eq. (2.17) or Fig. 2.4, and the true temperature is calculated using Eq. ( 2 . 7 ) . However, if the change in 0 due to the wafer thickness is small, a much simpler procedure can be used. Differentiation of E q s . (2.6,2.7) leads to the following expression for the change in the real temperature: (2.18) is the maximum real temperature achieved for infinite where,,T samples (see Fig. 2.1). [Note that temperatures in Eq. (2.18) The relative change in G, have to be expressed in Kelvin.] AO/Omax, is obtained from Fig. 2.4.
2. 2.3.4
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
81
Effect of Absorption Coefficient
Lax [2.81 has shown that for a finite absorption coefficient ,,the maximum linear temperature is reduced. To account for this, , depends on we calculate the linear normalized temperature ‘Iwhich the dimensionless parameter W = aw: (2.19) where Omax (W=m) is given by Eq. (2.13). na can be expressed as 2
Lax i2.81 has shown that
(2.20)
Figure 2.5 gives the normalized linear temperature na as a function of W. The figure shows that the assumption of an infinite absorption coefficient is enerally good. For example, for the Ar+ laser, a is 1 . 2 ~ 1 0 4 cm-5 at 300°K for Si, and a typical value for w would be 20 um. This gives W = 24, and na = 0.97. However, if a Kr laser is used, the absorption coefficient is substantially lower, namely about 4000 cm-l at 300’K. The lower power output of that laser may necessitate the use of a smaller spot to obtain high enough temperatures, and the effect of the absorption coefficient may then have to be taken into account. The situation will be further complicated by the fact that the absorption coefficient also depends on temperature.
0.8
10
20
30
40
50
W:aw
FIGURE 2.5. Normalized linear surface temperature ‘I, = OmX(W)/ O,(W=m) at the beam center (x=y=z=O) as a function of the product, W, of the absorption coefficient, a, and the beam radius, w. [After Ref. 2.9.1
82
ART0 LIETOILA E 7 AL.
2.3.5
E f f e c t of Scanning Speed
C l i n e and Anthony [2.7] have shown t h a t t h e e f f e c t of t h e scanning speed on t h e l i n e a r temperature depends on t h e dimensionless parameter V:
v =
pcpwv 2
(2.21)
K(T~)
where v is t h e scanning speed. Again, t h e maximum l i n e a r temperaN i s s i m e t a l . 12.21 g i v e v a l u e s t u r e i s reduced by a f a c t o r , nv. of nv f o r a wide range of scanning speeds and f o r v a r i o u s a s p e c t r a t i o s 8 f o r e l l i p t i c a l beams ( s e e S e c t i o n 2.3.7). For low scanning speeds, nv can be approximated by t h e f o l l o w i n g expression:
nv
= 1.0
-
Figure 2.6 g i v e s
(2.22)
0.72 V
nv f o r large v a l u e s of V.
a W
a
I 06
a a w
2
5 0.4
0 W
1 '!
;; 0 2 z
a
0
z
"
1
3
2
4
v FIGURE 2.6. Normalized l i n e a r s u r f a c e temperature n = 6 ( V ) / Omax a t t h e beam c e n t e r (x=y=z=O) as a f u n c t i o n of t h e normalized scan speed V [ s e e Eq. (2.21)] f o r a c i r c u l a r beam (8=1) and e l l i p t i c a l beams of v a r i o u s a s p e c t r a t i o s 8 (see S e c t i o n 2.3.7). It should be emphasized t h a t normally, V << 1, i n which case Eq. 2.22 can be used f o r n. ( A f t e r Ref. 2.2.)
I n c o n t r a s t t o c o r r e c t i o n s due t o t h e wafer t h i c k n e s s and a b s o r p t i o n c o e f f i c i e n t , t h e i n v e r s e Kirchhoff t r a n s f o r m cannot be a p p l i e d t o t h e c o r r e c t e d l i n e a r temperature t o o b t a i n t r u e temperatures f o r t h e scanning beam s i n c e t h e Ktrchhoff t r a n s f o r m i s v a l i d only i n t h e s t e a d y state. Upper and lower limits f o r
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
83
the true temperature can, however, be obtained using the following procedure. Values for the parameter V in Eq. (2.21) are calculated using both the substrate backsurface temperature and the temperature which would be obtained if the beam were steady. Corresponding correction factors 'I, for 0 are then determined. If nv is close to one, Eq. (2.18) can be used to obtain the lower and higher limits for AT due to the scanning. For large deviations of nv from unity, corrected values of 0 should be calculated and the inverse Kirchhoff transform should be applied to obtain the upper and lower limits for the true temperature.
2.3.6
Calibration Procedure for Temperature Calculations
The determination of the annealing temperature in a practical experiment is greatly simplified if the system is calibrated in terms of the melting point of the semiconductor to be processed. This can be achieved by noting the reading af the laser power meter at which melting of the material just occurs at the beam center (allother conditions, To, etc., should be the same as will be used for the actual annealing). This power level will be called Pmelt. The annealing power is then expressed as a fraction, P/Pmelt, of the melting power, and Fig. 2.7 can be used to determine the annealing temperature.
1400 -
u ~,lOOO
-
w
=
0.2
O
0.4
0.6 0
0.8
J1 .o
FIGURE 2 . 7 . Maximum true surface temperature at the beam center in silicon as a function of normalized power P/Pmelt for various values of substrate backsuface temperature To. (After Ref. 2.9.)
ART0 LIETOILA ET AL
84
This procedure is especially useful for Si, because no decomposition of the material occurs at submelting temperatures and once melting takes place, it leaves clearly visible rough traces on the otherwise mirror-like sample surface. The observation of melting is made easy if an interference contrast microscope is used. However, compound semiconductors, such as GaAs and InP, may decompose before melting occurs, and that will make the use of the calibration procedure more difficult. The calibration procedure described above will completely remove the uncertainty in temperature calculations due to the inaccuracy of the laser power meter (providing it still is linear), the reflectivity of the material and the spot size of the laser beam. If the same wafer thickness is used for calibration and annealing, possible errors due to finite sample thickness are also eliminated. However, if the product of the absorption coefficient and beam radius is so low that it affects the maximum temperature, a slight error will result due to the fact that the absorption coefficient depends on temperature. Finally, this calibration procedure will give the most accurate estimate for the temperature for scanning beams. However, some error will still occur, since the effect of the the scanning speed depends on the sample temperature as described above. Calculated temperatures which this procedure yields will be lower than the real temperatures, because the thermal conductivity in semiconductors decreases when temperature increases and the lowering of temperature due to scanning is more profound for low conThus, the effect ductivity material (see Section 2 . 3 . 5 above). of the scanning speed is more significant at Pmelt than at the power used for annealing. 2.3.7
The Elliptical Beam
In certain applications it is desirable to use an elliptical laser beam rather than a circular one. For example, if wide scan lines are desired, one could use an elliptic beam with the long axis perpendicular to the scan direction. On the other hand, scanning in the direction parallel to the long axis would yield a long dwell time without reducing the scanning speed. Elliptic beam shapes are readily obtained by using cylindrical lenses, either alone or in combination with spherical lenses. The incident intensity distribution of an elliptical beam can be described as
I
=
1, exp
-
(2 +$)
(2.23)
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
where the intensity .I follows : P I0 = r wx w Y
85
is related to the total incident power as
(2.24)
Nissim et al. 12.21, who have treated the case of elliptical beams in great detail, use the convention that wx is the short radius of the elliptic beam. The aspect ratio is then B = !Y>1 w
(2.25)
X
Figure 2.8 shows the effect of increasing f? for the linear temperature 0. In that figure, we plot the normalized linear tempera* circ ure, 0 , / as a function of position at the sample surface for different values of 8 . We can see that the increase of f3 results in a substantial decrease of the maximum temperature while the temperature distribution spreads out in the direction of the larger radius. [2.2] have also shown that the normalized circ linear temperature, rl = O/Gmax, at the center of the beam at the sample surface can be expressed as:
Nissim et al.
(2.26)
where K is the complete elliptic integral of the first kind. We plot Eq. (2.26) as a function of f3 in Fig. 2.9. This figure, together with Eqs. (2.7) and (2.13), can be used to calculate the maximum annealing temperature for any elliptical beam.
We would like to emphasize that Eq. (2.26), and Figs. 2.8 and 2.9, have been obtained by assuming that the sample is infinite. This assumption is not valid, if even the larger radius of the elliptical beam is comparable to or larger than the wafer thickness. In that case, the best way to determine the actual annealing temperature is to use the calibration procedure described in Section 2.3.6 above.
rt
circ
0 max is the maximum linear temperature achieved with a circular beam of radius w = wx, [see Eq. (2.13)].
ART0 LIETOILA ET AL
86
I Z
-
20 40
(r
0
O
1
I
I
2
3
I
I
I
4
5
6
Y = ylw,
FIGURE 2.8.
Normalized
linear surface
temperature 11 = G /
circ
emax
for various aspect ratios 6 = wy/wx as a function of normalized distance from the beam center along (a) the long radius (Y = y/wx) and (b) the short radius (X = x/wx> of the elliptical beam. V = 0 for all cases.
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
0 0.01 1
2
.
3 4 5 6 78910
0
20
87
2
~
60 80100
40
P = wy / wx circ FIGURE 2.9. Normalized l i n e a r s u r f a c e temperature O/Qma, a t t h e beam c e n t e r a s a function of t h e aspect r a t i o B of t h e e l l i p t i c a l beam. Thin Layers on t h e Semiconductor Sample: Reflectivity
2.3.8
E f f e c t s of
In many beam processing a p p l i c a t i o n s t h e semiconductor sample i s covered by t h i n metal o r d i e l e c t r i c layers. These l a y e r s have F i r s t , the reflectwo e f f e c t s on t h e temperature c a l c u l a t i o n s . t i v i t y of t h e layered s t r u c t u r e can be s i g n i f i c a n t l y d i f f e r e n t from t h a t of t h e bare semiconductor. Secondly, t h e thermal prope r t i e s of the s t r u c t u r e may a l s o change somewhat. Fortunately t h e l a t t e r e f f e c t i s n e g l i g i b l e i n most p r a c t i c a l beam processing a p p l i c a t i o n s , s i n c e the l a y e r thickness ( t y p i c a l l y l e s s than 0.5 pm) i s very small compared t o both the beam diameter and t h e thickness of the underlying s u b s t r a t e . For t h i s reason w e consider i n t h i s s e c t i o n t h e e f f e c t s of a change i n t h e r e f l e c t i v i t y only. I f t h e sample i s covered by a t h i n nonabsorbing d i e l e c t r i c l a y e r (e.g., Si02 o r Si3N4), t h a t l a y e r w i l l a c t as an a n t i r e f l e c t i o n coating. The r e f l e c t i v i t y of such a s t r u c t u r e i s given by [2.101:
__
nf2(1-nSl2 cos 2
+
+
(ns-nf2I2 s i n 2
n f 2 (1+nsl2 cos2+
+
(ns
R =
+ nf2I2
+
s i n2
+
(2.27)
88
ART0 LIETOILA ET AL.
where nf and n, are the indexes of refraction of the film and the semiconductor substrate, respectively. The phase difference Q. in the film is (2.28) where d is the film thickness and X o is the wavelength of the laser radiation in free space. The above formula is valid €or thick substrates. We plot in Fig. 2.10 the reflectivity of Si coated with SiOz or Si3N4 for the Ar+ laser (mean wavelength = 5000 a ) . Note that the maximum reflectivity is that of an uncoated semiconductor.
LAYER THICKNESS
(A)
FIGURE 2.10. Reflectivity of Si covered by an AR coating of Si02 or Si3N4 for the Ar ion laser (mean wavelength = 5000 a ) as a function of coating thickness.
If the coating film is of strongly absorbing material, e.g., metal, the entire radiation is absorbed within the layer and the reflectivity of the structure is simply the reflectivity of the coating material.
2. 2.4
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
89
TEMPERATURE CALCULATIONS FOR WIDE AREA OPTICAL SOURCES: THE ARC LAMP
One of the primary reasons for using arc radiation sources for annealing purposes, instead of lasers or tightly focused electron beams, is the need to achieve large spot sizes and thus large throughput for industrial applications. Large spot size of course implies that lateral heat flow in the sample is negligible, so the temperature calculations then become one-dimensional. Arc lamp annealing has been performed in two ways. When the isothermal 12.111 mode is utilized, the sample is thermally isolated, and will be uniformly heated. In the heat sink [2.12] mode, the backsurface is kept at a constant temperature, as in the case of laser annealing described in Section 2.3 above. We will first consider the isothermal mode and then present calculations for the heat sink mode using a xenon arc lamp as the radiation source. 2.4.1
The Isothermal Mode
When this mode is used, the wafer to be processed is thermally floating, and the whole wafer is illuminated at the same time. 'Ihe temperature of the wafer will be practically uniform (hence the term isothermal), and the temperature achieved is governed mainly by radiation of the sample. This method has the advantage of requiring very low power densities. However, the time constant for heating the sample is of the order of 10s [2.11], making very rapid anneals impossible. Also, if the heat is to be confined to a thin surface layer of the sample, this method is not practical. A convenient geometry for this form of annealing is obtained by imagining the arc source to be located near the pole of a spherical volume having walls of reflectivity Rs 1 with the sample mounted on thermal insulators in the center of this volume. Under these conditions, the wafer will receive radiation on both surfaces from all angles and will be cooled by radiation only. The total power Q (Watts) of the source (as well as radiation from the cavity walls) is coupled to the wafer with an efficiency n which is given by (2.29) where A is area (cm2) and R the reflection coefficient. The subscripts w and s refer to the wafer and sphere respectively. Since the radiation losses from the wafer are similarly recoupled back
90
ART0 LIETOILA ET AL.
to the wafer, these losses are reduced by the factor 1-n as compared with the losses which would occur in vacuum. Assuming that the walls of the sphere remain at temperature To (by external cooling or high thermal mass), the equation describing the temperature rise of the wafer is given by (2.30) where CJ is the Stefan-Boltzmann constant (5.67.10-12W~m-2~K-4) and the emissivity of silicon is taken to be sW = 0.5. (Note we assume E = 1-R for both wafer and sphere.) The steady state temperature Tf is
(2.31)
Q, Eq. (2.30) can be written
As an experimental guide it is useful to note that, because of the dependence, (dT/dt) is within 90% of its initial value until T = 0.6 Tf. Hence a first order estimate of the time required for T to reach its final value is obtained by extrapolating the initial slope to the final temperature. This gives
(2.33)
For a cavity of radius 15 cm and reflectivity of 99% containing a 300 um thick 4 " diameter wafer we obtain a coupling efficiency n of 0.741. For Tf = lOOO"C, such a wafer reaches steady state in a time [from Eq. (2.34)] of At = 16.3 sec.
Although Eq. (2.33) predicts that the time to reach a given steady state temperature increases as n increases, it should be noted that for constant power Q, At will decrease rapidly as n approaches unity. Powell et al. [2.11] have calculated the temperature evolution of a Si wafer subjected to xenon arc radiation of 22 W/cm2 in the absence of a reflecting cavity. The results are shown in Fig. 2.11. In this case, one needs only to equate the blackbody radiation of the sample with the incident intensity to obtain the steady state temperature.
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
f-- ARC
1000
91
LAMP OFF
800
2
I
600
c 400
200 0
I
0
10
I
20
I
1
30
TIME (SEC)
Experimental temperature v s . t i m e of a 2 i n c h FIGURE 2.11. diameter S i wafer exposed t o c o n s t a n t xenon a r c lamp r a d i a t i o n of 22 W/cm2 i n t h e i s o t h e r m a l mode. ( A f t e r Ref. 2.11)
2.4.2
The Heat Sink Mode
T h i s mode, where t h e backsurface of t h e sample i s k e p t a t a c o n s t a n t temperature, was f i r s t d e s c r i b e d by Gibbons [2.12]. The power d e n s i t i e s r e q u i r e d are s u b s t a n t i a l l y h i g h e r than in t h e i s o t h e r m a l mode, s i n c e t h e h i g h thermal c o n d u c t i v i t y of s i l i c o n causes a l a r g e h e a t flow through t h e sample. On t h e o t h e r hand, t h e t i m e c o n s t a n t d e s c r i b i n g t h e approach t o e q u i l i b r i u m i s s u b s t a n t i a l l y s h o r t e r (because of t h e h i g h i n t e n s i t y u s e d ) , and very r a p i d a n n e a l s are p o s s i b l e . Also, t h i s method i s p r e f e r r e d f o r a p p l i c a t i o n s where melting of t h e sample s u r f a c e i s d e s i r e d ; t h e m e l t depth i s e a s i e r t o c o n t r o l because t h e r e i s a temperature gradient across t h e wafer. The temperature c a l c u l a t i o n s i n t h i s mode are complicated by t h e broad wavelength spectrum of arc r a d i a t i o n s o u r c e s , which makes a s i n g l e a b s o r p t i o n c o e f f i c i e n t i n a p p l i c a b l e . This problem can be overcome by u t i l i z i n g t h e s u p e r p o s a b i l i t y of t h e l i n e a r temperature 0. To do t h i s , t h e wavelength spectrum of t h e r a d i a t i o n source is divided i n t o s m a l l i n t e r v a l s , each of which i s assumed t o have a constant a b s o r p t i o n c o e f f i c i e n t . The l i n e a r temperatures produced by t h e s e i n t e r v a l s a r e c a l c u l a t e d and summed. The i n v e r s e Kirchhoff t r a n s f o r m is then a p p l i e d t o t h e sum t o o b t a i n t h e real temperature. L i e t o i l a e t a l . , 12.131 have a p p l i e d t h i s procedure t o calc u l a t e t h e temperatures achieved i n s i l i c o n i r r a d i a t e d by a xenon a r c l a m p . It was shown t h a t s t e a d y s t a t e s o l u t i o n s t o t h e
92
ART0 LIETOILA ET AL.
temperature can be used, if the dwell-time of the annealing spot is about 30 ms or more. The results are given in Fig. 2.12 for two different backsurface temperatures and wafer thicknesses. We can see that use of a thicker sample and higher backsurface temperature results in a saving of arc lamp power. However, even in the best case shown in Fig. 2.12, the absorbed intensity required to reach useful annealing temperatures is about 5 kW/cm2, which may be difficult to achieve in practice.
t I
XENON ARC L A M P
SI S U B S T R A T E I
1
I
I
I
1
I
I
2
4
6
8
10
12
14
16
3
ABSORBED POWER D E N S I T Y ( k W / c r n Z )
FIGURE 2.12. Calculated steady state surface temperature of Si in the heat sink mode as a function of absorbed xenon arc lamp intensity for different values of heat sink temperature To and sample thickness L. (After Ref. 2.9) Olga2 Elsevier Science Publishing Co. An easy way to achieve the temperature needed for annealing, while still maintaining a well-controlled temperature gradient across the wafer, is to place a poorly conducting layer between the semiconductor sample and the heat sink. A wafer of quartz serves well as such a layer. Lietoila et al., I2.131 have calculated temperatures achieved under the same conditions as in the optimum case of Fig. 2.12 (375 Dm thick wafer, To = 5 O O 0 C ) , using quartz wafers of different thicknesses. The temperature dependent thermal conductivity of quartz given in Ref. 12.141 was used in these calculations. The results are given in Fig. 2.13, and show that the required intensities are in a range much more easily achievable than if the wafer were placed directly on the heat sink.
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
I
I
I
I
I
I
I
I
I
I
I
1
I
93
1
S i : L . 3 7 5 prn# T o = 5 0 0 0 C I
0.2
I
0.4
I
1
1.0 1.2 1.4 ABSORBED INTENSITY ( kW/crn2)
0.6
0.8
I
1.6
I
1.8
2.0
FIGURE 2.13. Calculated effect of a quartz layer placed between the heat sink and the Si sample being annealed with a xenon arc lamp. The figure shows the surface temperature for the optimum case of Fig. 2.12 (To = 500°C, L = 375 pm) when different quartz layer thicknesses are used. 01982 E l s e v i e r Science Publishing co. Finally, we would like to emphasize that the above calculations have been carried out assuming that the thermal contact between the wafer and the heat sink is perfect. This is not necessarily the case in practice, unless a good thermal conductor, e.g., liquid tin or gallium, is used to improve the contacts. Therefore, if accurate information about the annealing temperature is needed in a particular experiment, it would be advisable t o perform the melt temperature calibration described in Section 2.3.6 above. 2.5
TEMPERATURE CALCULATIONS FOR SCANNING CW ELECTRON BEAMS
Processing of semiconductors by scanning cw electron beams is broadly analogous to cw laser processing. There are, however, a few differences which have to be considered in making temperature calculations. The principal dissimilarities are as follows. a. b.
The reflectivity is essentially zero f o r any material. The lateral intensity distribution may not be Gaussian.
94
ARTO LIETOILA ET AL.
c.
d.
The energy deposition profile is not exponential in depth, but approximately Gaussian, with the peak normally inside the material; this "absorption" profile does not depend on temperature. The beam will spread out when entering the material; this effect may be very significant if narrow beams are used.
The lack of reflectivity is a simplifying factor, since dielectric layers will not change the absorbed power. This feature is especially useful for processing of device structures, which have different layer thicknesses. The other three points above, b-d, make the formulas presented in Sections 2.3.1-2.3.4 above inapplicable for e-beam annealing. However, the temperature calibration procedure given in Section 2.3.6 is valid for (stationary) e-beams, no matter what the intensity profile is laterally or in depth. Moreover, since the energy deposition profile vs. depth does not change with temperature as does the optical absorption coefficient, the procedure is accurate in this respect for the e-beam. However, if the beam is scanning fast, the determination of the temperature will become inaccurate in the same fashion as for scanning cw lasers. Temperature Calculations for Finely Focused Electron Beams
2.5.1
Some applications, e.g., the direct writing of patterns during lithography, require an electron beam with extremely small lateral dimensions, on the order of 1 um. In such cases the beam dimensions are not only small compared with the sample thickness (necessitating a three-dimensional analysis) but are in addition small compared with the range of the electrons themselves. In analogy with the absorption depth for laser radiation, the relevant parameter in this case is called the Grun range [2.18] and is denoted Rg. In Fig. 2.14 the Grun range Rg is plotted as a function of electron energy Eb for electrons in silicon. More generally, Rg is given approximately by Rg
=
4.57 Ei*75 pg/cm2
5 keV
<
E,,
<
25 keV
(2.34)
where Eb is expressed in keV. To calculate temperatures in this case, one must first obtain a realistic representation of the heat source. This has been accomplished using Monte-Carlo techniques to determine a universal energy dissipation function I2.19, 2.201. Analytic techniques based on Green's functions were then used to calculate the temperature from the known source function.
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
8,
I
I
I
95
I
FIGURE 2.14. Griin range Rg as a f u n c t i o n of energy Eb f o r electrons i n silicon.
I n Fig. 2.15, we p l o t 0, = @/Om,, t h e normalized l i n e a r temperature a t t h e c e n t e r of a c i r c u l a r Gaussian e-beam a s a funct i o n of t h e dimensionless parameter Rg/w where w i s t h e beam /w small, t h e r e s u l t s a g r e e w i t h t h e cw r a d i u s . Note t h a t f o r laser c a s e i n w h i c h t h e a s o r p t i o n c o e f f i c i e n t a = 0. An a n a l y t i c approximation t o t h e curve i n Fig. 2.15 i s given by
%
= 0.981
-
0.359
(%)+
0.0628
(>)
3
- 3.96'10-3
(k)
(2.35)
The t r u e maximum temperature is once a g a i n obtained v i a Eqs. (2.13) and (2.7) w i t h t h e power P given by t h e beam voltagec u r r e n t product and t h e r e f l e c t i v i t y R = 0. S i m i l a r c a l c u l a t i o n s have been performed f o r a "ribbon" a beam very long i n t h e y d i r e c t i o n but e l e c t r o n beam, i.e., narrow i n t h e x d i r e c t i o n . I n t h i s case, i n s t e a d of Eq. (2.13), t h e l i n e a r temperature a t t h e c e n t e r of t h e beam i s given by t h e modified one-dimensional e x p r e s s i o n 0 -
PYL
vrii w
K(T~)
On
(2.36)
where Py is now t h e power p e r u n i t l e n g t h of t h e beam. On i s p l o t t e d i n Fig. 2.16 a s a f u n c t i o n of w/L where h e r e w i s t h e h a l f - w i d t h of t h e Gaussian ribbon beam and L i s a g a l n t h e sample
96
ART0 LIETOILA ET iQL.
thickness. The results have been found to be insensitive to the value of Rg. Note that for w/L greater than about 10, the onedimensional solution applies.
0.8
e W a
z w
c
0.6
[r
a W
5
0.4
-I
0 W N
5 0.2 a E [r
0 Z
FIGURE 2.15. Normalized linear surface temperature at the center of a circular Gaussian e-beam as a function of the dimensionless parameter Rg/w. The effect of scan speed on the lateral temperature distribution for a ribbon e-beam of Gaussian half-width 1.4 lim and energy 20 keV incident on a 250 lim thick silicon wafer is shown in Fig. 2.17. It can be seen from the figure that, for this case, distortion in the lateral temperature profile is significant only for velocities greater than about 20 cm/sec. 2.5.2
Temperature Calculations f o r Wide Area Ribbon Electron Beams
We now wish to present temperature calculations for the surface of a thin sample which is irradiated with a ribbon electron beam which has a width of order 1 mm. For these calculations it is assumed that the electrons deposit their energy in a very narrow region near the surface. For typical electron energies of 5 keV, this is a good approximation since the range of such an electron is only about 0.5 pm. We also assume a constant thermal conductivity K and apply the Kirchhoff transform to obtain the true temperature when such a transformation is valid.
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
97
1 .o
co= W
IL
2 k
0.8
IL W
a
5
0.6
W
N 0.2 J
a
E
0 2
w/L
FIGURE 2.16. Normalized linear surface temperature at the center of a Gaussian ribbon e-beam as a function of the ratio of beam half-width w to the sample thickness L.
1 .o
0
Gaussian ribbon e-beam
%
0.8-
U Q
z w
0.6-
-v.0
- .-
E 20keV w:2.8pm L: 250pm
v: 2 0 c m l s
-
Q 3
w t
2a nk
0.4-
-
p I-
0.2
-
-
11 IYw
I
I
I
I
---+--
FIGURE 2.17. Normalized linear surface temperature as a function of lateral position for a Guassian ribbon e-beam showing the asymmetry induced at high scan speeds.
ART0 LIETOILA ET AL.
98
Consider a beam which deposits energy at the rate Q per unit area per unit time in a region of width W = 2b at the surface of a semi-infinite solid. Furthermore, assume that this beam is scanned with velocity V in the -x direction. It has been shown [ 2 . 2 1 ] that the linear temperature rise due to such a beam in the steady state is given by X+B O,(X,Z)
=
npc
eu%
v
(4u2 + ) du Z2
(2.37)
X-B where X = nx, 2 = nz and B = nb, n = Pcpv/2K and K, and K1 are modified Bessel functions of the second kind which satisfy (2.38)
For a sample of finite thickness L and a back surface which is either thermally isolated or thermally sunk, we once again utilize the method of images to obtain the surface temperature (2.39)
where the upper sign corresponds to the case of thermal isolation and the lower sign to the case of a heat sunk back surface. Equations (2.37) and (2.39) have been evaluated by numerical integration with the following results.
2.5.2.1
Thermally Isolated Wafer
When the scan speed is less than about 1 cm/sec, the front and back surfaces are very nearly at the same temperature for a given lateral position x and a first estimate of the temperature rise is simply given by the total energy deposited per unit volume divided by the specific heat. The first approximation is independent of the thermal conductivity and scan speed (for slow enough scan speeds V). For very slow scan speeds, however, one must account for the cooling effect of radiation from both surfaces. In our calculations the effects of radiative cooling are accounted for by correcting the temperature rise at each lateral position (starting from the leading edge of the beam) by an amount equal to the cooling which occurs during a time T=dx/V where dx is the distance between lateral points considered. This is just the dwell time corresponding to each lateral point.
2. TEMPERATURE DISTRIBUTIONS AND REACTION RATES
99
When the scan speed exceeds about 1 cm/sec, the sample has little time to cool so radiative losses can be neglected. However, in this case another effect becomes important; i.e., the thermal conductivity limits the rate at which the back surface can come to equilibrium with the front surface. Thus the maximum surface temperature is higher than would be obtained assuming uniform heating. Furthermore, this maximum is dependent on the thermal conductivity. Since this maximum is also scan-speeddependent, the Hrchhoff transform cannot be used to remove the temperature dependence of the thermal conductivity. For these cases we have performed calculations using both the thermal conductivity at the initial temperature (OOC), and the thermal conductivity at the final temperature. The true temperature must therefore lie between these two values. The results of these calculations are presented in Fig. 2.18. The abscissa is now the deposited energy density and thus has units of JIcm3. The dashed line gives the resulting temperature assuming uniform heating of the wafer with no losses due to radiation.
-
I200
-
ISOLATED BACK SURFACE
L* 400pm
W*lmm
0
L loOO-
_ _ _ UNIFORM HEATING WITHOUT RADIATIVE COOLING
E
r
800-
Y Y
3a
6W-
3
400
-
200
-
I
z
2
' 0
300
600
900
1200
1500
NORMALIZED W W E R
le00
2wx)
2400
2700
3000
( J /cm3)
FIGURE 2.18. Maximum silicon surface temperature as a function of normalized power for a scanning ribbon e-beam with rectangular intensity distribution of width W. The sample is assumed to be thermally isolated. O1982 Elsevier Science Publishing C O .
For slower scan speeds, radiative losses result in a lower maximum temperature while for higher scan speeds, the thermal gradient across the wafer thickness results in a higher maximum surface temperature. The resulting temperatures obtained for a scan speed
100
ART0 LIETOILA ET AL
of 5 mmlsec show that in this case radiative cooling is significant f o r temperatures above about 800°C. When the scan speed is 1 cm/sec, both cooling effects and thermal gradient effects are evident. The range of power values obtained using the maximum and minimum thermal conductivities is indicated in Fig. 2.18 by the width of the shaded region. For a scan speed of 10 cm/sec, radiative cooling is negligible, but the maximum surface temperature is strongly dependent on the thermal conductivity. 2.5.2.2
Heat Sunk Back Surface
When the back surface of the sample is heat sunk, much more power is necessary to reach a given temperature. A s a result, losses due to radiation are negligible in all cases. In addition, for scan speeds less than about 5 cm/sec, the results are independent of scan speed. Thus in these cases the Kirchhoff transform is valid and the temperature dependence of the thermal conductivity can be properly accounted for. In Fig. 2.19 we show the results for back surface temperatures of 0 and 350°C for a lmm wide beam. Also shown are calculations at a scan speed of 1 cmlsec for beams of width 100 um and 500 um. For these narrow beams it is evident that lateral diffusion is severe enough to cause considerable reduction in the
1200
-
HEAT SUNR BACK SURFACE
LiLwpm
r
I
E
1ooo-
Q
" w
zoo
0
2
4
6
8
10
POWER DENSITY
12
I4
16
18
20
I rw/crn'~
FIGURE 2.19. Maximum silicon surface temperature as a function of power density for a scanning ribbon e-beam with rectangular intensity distribution of width W incident on a sample which is heat sunk. el982 Elsevier Science Publishing C o .
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
101
maximum s u r f a c e temperature achieved, even a t c o n s t a n t power dens i t y . It has been found, however, t h a t t h e r e i s l i t t l e dependence on t h e beam width f o r widths g r e a t e r than about l m m . Thus i n such cases t h e h e a t problem i s e s s e n t i a l l y one-dimensional. 2.6
REACTION RATE CALCULATIONS
A v a r i e t y of p h y s i c a l p r o c e s s e s induced by scanning cw energy beams have been i n t e r p r e t e d i n terms of solid-phase r e a c t i o n mechanisms. These processes i n c l u d e t h e e p i t a x i a l regrowth of i o n implanted s i l i c o n a s w e l l as t h e formation of c e r t a i n metals i l i c i d e s , and are d i s c u s s e d i n d e t a i l i n subsequent c h a p t e r s of t h i s book. The behavior of t h e s e systems during low-temperature f u r n a c e p r o c e s s i n g can be d e s c r i b e d by a s i n g l e a c t i v a t i o n energy, and t h e r e a c t i n g i n t e r f a c e can be considered t o be a t a c o n s t a n t temperature throughout t h e a n n e a l . The e f f e c t of cw beam i r r a d i a t i o n i s t o produce a high temperature and an a c c e l e r a t e d r e a c t i o n r a t e , but i n t h i s case t h e temperature is a r a p i d l y changing W e discuss here the a n a l y t i c a l formulation f u n c t i o n of t i m e . f o r solid-phase r e a c t i o n k i n e t i c s induced by a scanning cw beam, i n c l u d i n g t h e dependence of t h e temperature on t i m e . The r e s u l t s of such c a l c u l a t i o n s permit a comparison of observed beam processi n g r e s u l t s w i t h e x t r a p o l a t e d furnace-processing d a t a .
For a r a t e - l i m i t e d process such a s t h e regrowth of amorphous implanted S i , t h e reacted-layer ( o r regrown-layer) thickness furnace anneal can be r e s u l t i n g from a constant-temperature expressed as d = R o t exp(-Ea/kT)
(2.40)
where t and T a r e t h e anneal t i m e and temperature, r e s p e c t i v e l y , and Ea i s t h e a c t i v a t i o n energy. For d i f f u s i o n - c o n t r o l l e d react i o n s , such as t h e formation of c e r t a i n m e t a l - s i l i c i d e s , t h e d i n Eq. (2.40) i s replaced by d2, but a l l o t h e r f e a t u r e s of t h e following d i s c u s s i o n are s t i l l a p p l i c a b l e . The t r a n s i e n t n a t u r e of t h e temperature rise produced by a scanning cw beam means t h a t t h e r e a c t e d - l a y e r t h i c k n e s s should be expressed by 00
d = Ro
exp[-E,/kT(t)]dt
(2.41)
4
If t h e dependence of temperature on t i m e i s known, Eq. (2.41) can be reduced t o Eq. (2.40) by determination of an " e f f e c t i v e a n n e a l i n g t i m e , " t e f f . The f i r s t s t e p is t h e c a l c u l a t i o n of t h e
102
ART0 LIETOILA ET AL.
s p a t i a l temperature p r o f i l e . Then, u s i n g knowledge of e x p e r i mental scan c o n d i t i o n s , t h i s s p a t i a l p r o f i l e i s c o n v e r t e d t o a temporal one. We c o n s i d e r f i r s t t h e case of a s i n g l e s c a n through t h e p o i n t of i n t e r e s t ( x , y ) = ( 0 , O ) . The beam i s assumed t o have a Gaussian i n t e n s i t d i s t r i b u t i o n given by Eq. (2.12), w i t h an a b s o r p t i o n depth much less t h a n t h e beam r a d i u s W. We s t i p u l a t e f u r t h e r t h a t t h e beam i s moving w i t h a v e l o c i t y v i n t h e x d i r e c t i o n , and t h a t i t s c e n t e r p a s s e s over ( x , y ) = ( 0 , O ) a t t = 0. Using t h e r e s u l t s of Eq. (2.16) t h e l i n e a r temperature r i s e O ( t ) a t t h e s u r f a c e can be shown t o be
ct-l
O(t) =
~ , ( v2t 2 /2w 2 )
omax exp(-v2t2/2w2)
(2.42)
where O ,, = ' ( 0 ) i s g i v e n by Eq. (2.131, and I, i s t h e modified Bessel f u n c t i o n of o r d e r 0. A s p o i n t e d o u t e a r l i e r , w e assume t h a t t h e r e a c t i n g - l a y e r t h i c k n e s s i s much less than t h e beam r a d i u s and t h a t , consequently, t h e s o l u t i o n [Eq. (2.42)] i s a l s o Since t h e growth r a t e i s an valid at the reacting interface. extremely s t r o n g f u n c t i o n of temperature, v e r y l i t t l e accuracy i s l o s t (2.151 by a f i r s t - o r d e r expansion of O ( t ) n e a r Omax. Hence w e e x p r e s s O ( t ) as O(t)
2
omax -
'max
v2,2
(2.43)
2w2
The t r u e t e m p e r a t u r e T ( t ) i s t h e n approximately
T(t)
'
Tmax
-
%ax
v 2 t 2-dT -
2w2
(2.44)
dO 0X ,
Equation (2.6)
T(t)
Z
T ,,
[
1-
can be used t o e l i m i n a t e 0, w i t h t h e r e s u l t
2 2 v t
(To-T,)ln[(T,,-T,)/(To-TK>
2w2
TmaX
1
(Tmx-TK)
(To
-
1
TK>
(2.45)
where Tmax = T, + (To-TK> exp
P( 1-R) 2 71'/ 2wA
(2.46)
can be d e r i v e d u s i n g Eqs. (2.7) and (2.13), and A i s t h e parameter The e x p r e s s i o n f o r T ( t ) i s now i n s e r t e d i n l i s t e d i n Table 2.1. Eq. (2.41) t o c a l c u l a t e t h e t o t a l r e a c t i o n produced by a scan through t h e p o i n t of i n t e r e s t . This i n t e g r a t i o n y i e l d s t h e r e s u l t d
3
R o t e f f exP(-Ea/kTeff)
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
103
where
and
(2.47) The result has been expressed in terms of an "effective temperature," Teff, and an "effective time," teff. The effective temperature is the peak temperature, while the effective time is equal to the beam dwell-time, 2w/v, multiplied by a "dwell-time reduction factor," f. The calculated dependence of f on T, and Tmax is shown in Fig. (2.20). It can be seen that the effective reaction time for a thermally-activated process is typically onethird of the actual beam dwell-time. Higher activation energies and/or lower substrate temperatures lead to smaller values of teff for a given TmX, as would be expected intuitively.
0 0
100
200
300
400
SUBSTRATE TEMPERATURE, To ( O C )
FIGURE 2.20. Calculated dependence on To and Tmax of the dwell time reduction factor f. ' h o values (1.5 and 2.35 eV) were used for the activation energy of the reaction process.
104
ART0 LIETOILA ET AL.
This single-scan calculation can be extended to the case of multiply-scanned, uniformly-reacted samples, such as those which would be processed for large-area device structures. The reacted layer thickness produced at a given point is now the result of many different beam scans, not all of which pass directly through that point. The solution can be written as [2.16]
where
and (2.48) The variables t and A represent the total scan time and total scan area, respectively, and the dwell-time reduction factor f is identical to that defined in Eq. (2.48). Examples of the use of these formulas will be given in Chapters 3 and 6, where beam-induced solid phase epitaxial regrowth of implantation-amorphized silicon and silicide formation are described, respectively.
2.
TEMPERATURE DISTRIBUTIONS AND REACTION RATES
105
REFERENCES 2.1 2.2 2.3 2.4 2.5 2.6 2.7 2.8 2.9 2.10 2.11 2.12
"Conduction of Heat Carslaw, H. S., and Jaeger, J. C., in Solids", 2nd ed., p. 11. Clarendon Press, 1959. Nissim, Y. I., Lietoila, A., Gold, R. B., and Gibbons, J. F., J. Appl. Phys. 51, 274 (1980). Maycock, P. D., Solid St. Electron. 10, 161 (1967). Ho, C. Y., Powell, R. W., and Liley, P.E., J. Phys. Chem. Ref. Data 3, Suppl. 1, 1-588 (1974). Slack, G. A., J. Phys. Chem. Solids 34, 321 (1973). Nissim, Y. I., Ph.D. Thesis, Stanford University, Department of Electrical Engineering (1981). Cline, H. E., and Anthony, T. R., J. Appl. Phys. 48, 3895 (1977). Lax, M., J. Appl. Phys. 48, 3919 (1977). Lietoila, A., Ph.D. Thesis, Stanford University, Department of Applied Physics (1981). Hecht, E., and Zajac, A., "Optics", p. 314. Addison-Wesley (1974). Powell, R. A., Yep, T. O . , and Fulks, R. T., Appl. Phys. Lett. 39, 150 (1981). Gibbons, J. F., in "Laser and Electron-Beam Solid Interactions and Materials Processing" (J. F. Gibbons, L. D. Hess and T. W. Sigmon, eds.), p. 449. North-Holland, (1981).
2.13 2.14 2.15 2.16 2.17 2.18 2.19 2.20 2.21
Lietoila, A . , Gold, R. B., and Gibbons, J. F., to be published. Touloukian, Y. S., Powell, R. W., Ho, C. Y., and Klemens, D. G., "Thermophysical Properties of Matter," Vol. 2, p. 182. IFI/Plenum (1970). Gold, R. B., and Gibbons, J. F., J. Appl. Phys. 51, 1256 (1980). Gold, R. B., and Gibbons, J. F., in "Laser and Electron Beam Processing of Materials" (C. W. White and P. S. Peercy, eds.) p. 77. Academic Press (1980). Siegman, A. E., "An Introduction to Lasers and Masers," p. 306. McGraw-Hill (1971). Grun, A. E., 2. Naturfosch. 12a, 89 (1957). Hawryluk, R. J., Hawryluk, A. M., and Smith, H. I., J. Appl. Phys. 45, 2551 (1974). Iranmanesh, A. A., SEL Technical Report, Stanford University (1980). Ref. 2.1, Chapter X.
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CHAPTER 3
Applications of CW Beam Processing to Ion Implanted Crystalline Silicon Arto Lietoila* and James F. Gibbons STANFORD ELECTRONICS LABORATORIES STANFORD UNIVERSITY STANFORD. CALIFORNIA
3.1 INTRODUCTION . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2 EXPERIMENTAL METHODS............................ 3.3 cw LASERINDUCED CRYSTALLIZATION OF AMORPHOUS SILICON............................... 3.4 ANNEALING OF 75AS IMPLANTED SI: C < C,, . . . . . . . . . . . . OF 75AS IMPLANTED SI: c > c,, . . . . . . . . . . . . 3.5 ANNEALING 3.6 HIGHRESOLUTIONSELECTIVE E-BEAM ANNEALING OF A S IMPLANTED SILICON . . . . . . . . . . . . . . . . 3.7 ANNEALING OF B IMPLANTED SI ...................... REFERENCES. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3.1
107 108 122 137 147 165 166 173
INTRODUCTION
Historically, cw laser processing was first applied to the annealing of As-implanted crystalline Si [3.1]. The dose and energy chosen for these experiments was sufficient to amorphize the silicon to a depth of about 0.1 pm. It was found that with appropriate laser processing, the amorphized silicon could be completely recrystallized with 100% electrical activity and no residual defects as judged by "EM. Furthermore, the recrystallization was observed to produce no redistribution of the As from the as-implanted profile. 'Ihe lack of redistribution itself strongly suggests that the recrystallization process is closely akin if not identical to solid phase epitaxy. Subsequent to these first results, it was shown directly by several workers [3.2-3.41 that the annealing process for amorphized samples is a solid phase epitaxial regrowth identical to that obtained in conventional furnace annealing. It was also shown [3.5] that B-implanted silicon can be satisfactorily annealed (though with less ideal results than for amorphized samples); and it was further demonstrated [3.6-3.81 that 100% electrical activity can be obtained in As-implanted silicon even for impurity concentrations that exceed the equilibrium solubility. Metastable concentrations are successfully produced *Present address: Micronas, Inc., 00101 Helsinki, Finland.
SEMICONDUCTORS AND SEMIMETALS, VOL. 17
107
Copyright 0 1984 by Academic Press, Inc. All rights of reproduction in any form reserved. ISBN 0-12-752117-8
108
ART0 LIETOILA AND JAMES E GIBBONS
also by pulsed laser annealing, but since this process inevitably involves melting f 3 . 9 1 , the as-implanted profile cannot be preserved. The crystalline quality, as judged by TEM, is good in cw laser annealed silicon layers. However, DLTS measurements have revealed production of high densities of point defects well beyond the range of the original implantation, a topic which is discussed in detail in Chapter 4 . Results obtained using a scanning e-beam, both in an electron beam welder and in a scanning electron microscope, are largely similar to those for cw laser scanning. ‘Ihe most significant difference between these methods and the use of a laser occurs if the silicon sample to be annealed is covered by an oxide or other dielectric layer. If a laser is used, the dielectric layer will act as an anti-reflective coating, and the laser power must then be adjusted accordingly (see Fig. 2.10 for reflectivity values). This may cause problems in device structures, where the oxide thickness is not uniform throughout the scanned area. For the e-beam, however, this problem does not exist, since no reflection takes place from either the silicon or from silicon dioxide. On the other hand, charging of the oxide may be important for certain choices of the beam current and voltage. In this chapter, we first discuss the analytical techniques used in evaluation of cw laser and e-beam annealing of ion implanted single crystal silicon. We then give a brief discussion of solid phase epitaxy, after which we explain in detail the results of cw beam processing in amorphous and non-amorphous Si. 3.2
EXPERIMENTAL METHODS
When the results of cw beam annealing of ion implanted silicon are evaluated the experimentalist is usually most concerned about the behavior of the impurity atoms. It is important to know to what extent the doping atoms are electrically active and how the distribution of them has changed during the annealing cycle. These questions can be answered by Secondary Ion Mass Spectroscopy (SIMS) and electrical Van der Pauw (VdP) measurements. The former gives the distribution of all the impurity atoms, without making a distinction between electrically active and inactive species, whereas the Van der Pauw measurements yield the distribution of the electrically active atoms. The VdP measurements give also the Hall mobilities of charge carriers. If the mass of the impurity atoms is larger than that of Si, the total number of impurity atoms (per unit area) can be accurately determined by Rutherford Backscattering (RBS) techniques; this provides a valuable calibration meth0.d for the implanted dose. The distribution of the impurity atoms can also be determined, though the accuracy is not very good unless a glancing angle method is used. When RBS measurements are performed in the channeling mode, information is obtained about the lattice location (interstitial/substitutional) of impurity atoms. A usual
3.
ION IMPLANTED CRYSTALLINE SILICON
109
aesumption is that shallow impurities which look substitutional by channeling are electrically active. However, we will show in Section 3.5.3 that this is not always true. Since the practical goal of annealing of implanted Si is application in device processes, the crystalline quality of Si is also of major importance. RBS channeling techniques provide one powerful method for judging the crystalline quality of the annealed layers. Quantitative information about the structure of large defects (> 5 nm) can be obtained by Transmission Electron Microscopy (TEM). The depth distribution and electronic properties of point defects are detected by Deep Level Transient Spectroscopy (DLTS)
.
In this section, very brief reviews of VdP, SIMS and RBS measurement techniques are given. The DLTS technique and its applications are treated in Chapter 4 . With regard to TEM, we refer to standard textbooks; e.g., Ref. [3.10]. 3.2.1
Van der Pauw Measurements
Van der Pauw measurements of a doped layer can be performed in two ways. The sheet measurements give the true sheet resistivity of the layer and an estimate, Ns, for the active sheet carrier concentration; a weighed average mobility v, can then be calculated. Differential Van der Pauw measurements, in turn, give the carrier concentration and mobility as a function of depth. In the following, the two types of measurements are explained. The details of the anodic sectioning for stripping measurements are then given. Some of the error sources are also discussed. 3.2.1.1
Sheet Van der Pauw Measurements
.
The geometrical configuration for the measurements has been presented by Van der Pauw [3.11] These measurements give directly the sheet resistivity Rs and an estimated sheet carrier concentration Ns of the measured layer. The following formulas are valid for Rs and N, [3.12] :
0
ART0 LIETOILA AND JAMES F. GIBBONS
110
where d is the layer thickness, p is carrier concentration, and u mobility. The Hall scattering factor rH is defined as
where T is the carrier momentum relaxation time. electrons, rH = 1 [3.13].
For degenerate
The weighed average mobility, II, can be calculated from Rs and Ns: 1 (3.4)
!J=
4 Rs Ns
Use of formulas (3.1) and (3.2) yields
If the mobility were constant throughout the doped layer, the sheet measurements would give correct values for the active sheet concentration and mobility. However, the carrier mobility is a monotonically decreasing function of carrier concentration. In this case, the measured Ns will be lower than the real active dose. 'Ihe difference can be up to 30% for a heavy n-type dose, since p varies rapidly when the electron concentration is greater than 1019 cm-3. If the distribution of impurities is known, the above formulas (3.1, 3.2, 3.5) can be used to calculate the expected values for R,, u , and Ns. The measured values are then compared with the calculations, and if the agreement is good, the annealing is considered to be successful. This method is especially suitable for cw beam annealing, since the impurity redistribution is negligible during annealing, and the doping profiles can be accurately predicted. 3.2.1.2
Differential Van der Pauw Measurements
If information is needed about the carrier concentration and mobility vs. depth in the material, differential Van der Pauw measurements must be performed. In this method, thin sections of material are removed and the sheet measurements are performed after removal of each successive layer. This then allows the calculation of the sheet resistivity and sheet carrier concentration
3.
ION IMPLANTED CRYSTALLINE SILICON
111
(per unit area) in each layer. In other words, the limits of integrals in Eqs. (3.1, 3.2) are replaced by the limits of each layer. The weighed average mobility in each layer can also be calculated: the expression for it is Eq. (3.5) with appropriately modified integration limits. Providing that the thickness of each layer is known, the mean carrier concentration in it can be calculated, and concentration and mobility can be plotted as a function of depth. 3.2.1.3
Anodic Sectioning
The most widely used method to perform sectioning for differential Van der Pauw measurements employs anodic oxidation and oxide stripping. This is an elaborate procedure, but it is far superior to etching, since the thickness of the grown oxide can be measured after each step, allowing accurate determination of of the layer thicknesses. The standard anodization solution is prepared from [3.14] 1 gal. ethylene glycol, 15.4 g KNO3 (0.04 N solution), 90 ml H20 (2.5 vol - X ) , 6 g Al(N03)3*9H20. It is customary to illuminate the sample through the front surface during the oxidation. This is not essential for n-type layers on p-substrate. However, for p on n, the light is necessary to get enough current through the reverse-biased pn-junction (as the name of the method suggests, the Si substrate is positively biased in anodic oxidation). The thickness of the anodic oxide layer depends on the voltage prevailing across the oxide when this process is stopped, and calibration curves exist in literature [3.141. However, the main reason for using the anodization method is a desire to obtain accurate values for layer thicknesses. Therefore, the oxide thicknesses should be measured by ellipsometry or some other suitable method. Once the oxide thickness is known, it then has to be converted to the thickness of removed silicon. For thick oxide layers (> 1000 a), the ratio of silicon thickness to oxide thickness is expected to be about the same as for thick thermal oxide, namely 0.44. However, this ratio is larger when the oxide gets thinner. The ratio for about 45 nm thick oxide has been measured to be 0.56 f 0.01 [3.151. 3.2.1.4
Discussion of Error Sources
As was pointed out in Section 3.2.1.1., sheet Van der Pauw measurements give values that are systematically too low for the
112
ART0 LIETOILA AND JAMES F. GIBBONS
sheet carrier concentration, unless the mobility is constant throughout the whole layer. This is true also for differential Van der Pauw measurements: this measurement yields the sheet concentration in each layer and the total active dose is obtained by summing up the sheet concentrations of each layer. However, since the individual layers are thin, the variation of mobility is minute within each of them. Therefore, the differential measurements yield a much more accurate value for the active dose. The inaccuracy in the determination of the thickness of the removed layers does not cause any error in the value of the total active dose, since the differential measurements give the total number of carriers in each layer. However, when the volume concentration versus depth is calculated, a possible error in the thickness measurement causes errors in the shape of the carrier profile. The Van der Pauw method gives correct results only if the four contacts are point contacts and are at the edge of the sample. To achieve this, various etched "Van der Pauw structures" have been proposed [3.16]. Such structures are unnecessary, however, if sufficiently large square samples can be used. The contacts are formed by placing tungsten probes directly onto the sample surface. l'hese probes can be regarded as point contacts, but of course they cannot be placed exactly at the sample edge. Probably the probes can be placed within 0.2 mm of the sample edge. Van der Pauw [3.11] has calculated that the relative error caused by this in sheet carrier concentration is about that fraction of the sample dimension by which the probes are separated from the edge. Thus if 5x5 mm samples are used, the error should be less than -5%. Additional uncertainty in differential Van der Pauw measurements results if only a small fraction ((5%) of the active carriers is removed with each step. [In this case, two numbers which are very close to each other will be subtracted.] This problem usually occurs for the first point of measurement because the surface concentration of typical implants is quite low unless a very low energy is used. However, from the second layer on, a sufficient fraction of the remaining dopants is usually removed at each step to avoid this type of error.
3.2.2.
S IMS Technique
The popularity of secondary ion mass spectrometry (SIMS) as a technique for detecting impurities in semiconductors stems from the fact that, under appropriate conditions, it has the sensitivity to measure depth profiles of impurities down to concentrations on the order of 1 x 1015 cm-3 I3.17-3.191. In this section we shall give a brief overview of the general technique and will
3.
ION IMPLANTED CRYSTALLINE SILICON
113
show how SIMS may be applied to measure depth profiles of common impurities in semiconductors. 3.2.2.1
General Description
The traditional method of secondary ion mass spectrometry is schematically illustrated in Fig. 3.1 i3.201. A primary source of low energy ions is generated, which collide with the atoms of the surface of a specimen. This induces a beam of sputtered secondary ions which may then be mass analyzed. Monitoring the concentration of constituents in this secondary ion beam as a function of time gives a depth profile of each of the host and impurity atoms.
ION MICROPROBE 'ON
TO MASS SPECTROMETER/
FIGURE 3.1. Artist's conception of the secondary ion production process (after Ref. 3.20). The efficiency of ionization of the secondary ions determines the sensitivity of SIMS to these elements. There are two processes by which these atoms become ionized, (a) kinetic ionization and (b) chemical ionization. In the first of these, the incident ion must impart enough kinetic energy through electronic stopping losses to actually ionize the secondary sputtered ion. Chemical ionization is the process whereby the primary ions chemically ionize the surface atoms and is statistically the most favorable. It is therefore desirable to use a chemically reactive ion such as Oz, 0-, or Cs+ as the primary beam. Other elements such as Ar+ are used, but they tend to produce lower yields of either positive or negative secondary ions because they induce only kinetic ionization.
114 3.2.2.2
ART0 LIETOILA AND JAMES F. GIBBONS
Depth P r o f i l i n g of I m p u r i t i e s
Ion implanted i m p u r i t i e s i n S i a r e u s u a l l y d i s t r i b u t e d a t depths l e s s than 300 nm, r e q u i r i n g a depth p r o f i l i n g technique w i t h a depth r e s o l u t i o n 5 nm. One method of a c h i e v i n g t h i s The primary i o n high depth r e s o l u t i o n i s i l l u s t r a t e d i n Fig. 3.2. beam i s r a s t e r scanned i n o r d e r t o c r e a t e a wide flat-bottomed c r a t e r . An e l e c t r o n i c g a t i n g technique i s a p p l i e d such t h a t t h e d e t e c t o r only a c c e p t s i o n s w h i l e t h e beam i s scanning i n t h e In t h i s way, only t h e f l a t c e n t e r r e g i o n of t h e c r a t e r bottom. p o r t i o n of t h e c r a t e r c o n t r i b u t e s t o t h e s i g n a l r e c e i v e d . N e u t r a l primary i o n s , however, can n o t be r a s t e r scanned and a r e t h e r e f o r e a b l e t o s p u t t e r any p o r t i o n of t h e specimen a t any t i m e . Elect r o n i c a p e r t u r i n g i s of course unable t o reduce t h e n o i s e produced by n e u t r a l i o n s s t r i k i n g t h e c r a t e r s i d e w a l l s and t h e sample s u r f a c e . This l i m i t s t h e r e s o l u t i o n and s e n s i t i v i t y of SIMS, p a r t i c u l a r l y f o r t h e c a s e of an i o n implanted p r o f i l e where t h e s i d e w a l l impurity c o n c e n t r a t i o n may be s e v e r a l o r d e r s of magnitude h i g h e r than t h a t on t h e c r a t e r f l o o r .
-i
BEAM
1-26v
i SAMPLE
I,
SAMPLE
I
I
CRATER PRODUCED BY RASTERED PRIMARY BEAM
FIGURE 3.2. I l l u s t r a t i o n of craters produced by a r a s t e r e d and u n r a s t e r e d beam ( a f t e r Ref. 3.60). F o r c e r t a i n atoms such a s s u l f u r , t h e r e i s a n o t h e r l i m i t a t i o n t o t h e u l t i m a t e system d e t e c t a b i l i t y [3.211. Diatomic oxygen h a s t h e same atomic mass number a s monatomic s u l f u r and i s t h e r e f o r e A high background of r e s i d u a l oxygen analyzed simultaneously. w i l l t h e r e f o r e produce a f a l s e b a s e l i n e on t h e SIMS output. E f f e c t s such a s t h i s due t o molecular "feedthrough" can u s u a l l y
3.
ION IMPLANTED CRYSTALLINE SILICON
115
be minimized by improving the ultimate vacuum in the analysis chamber. An example of a SIMS output is given in Fig. 3.3. for a Te implant in GaAs. The horizontal axis represents the number of Tecounts in the detector during a detecting time-interval, and the horizontal axis is the sputtering time. This data is converted to impurity concentration vs. depth through suitable calibration procedures.
lx10’4 Te/cm2 120 keV
-0 o 1 SECONDS
FIGURE 3 . 3 . SIMS output from a Te implanted GaAs sample (after Ref. 3.60). The depth calibration is conveniently performed by measuring the depth of the sputtered crater by a step profiler. Since the total sputtering time is known, it is then easy to convert the sputtering time to depth. This procedure of course assumes that the sputtering rate is constant, which is usually the case for silicon. When converting counts/interval to impurity concentration, the background count rate is first subtracted. That rate is obtained from the end of the SIMS output. For example, in Fig. 3 . 3 , the background rate (for time > 200s) is about 0.5 countsfs. Final calibration is then performed by requiring that the integrated impurity dose be the same as the implanted dose. Another way to calibrate concentrations is to use known calibration samples, which should be analyzed in the same run together with the unknown samples.
116
ART0 LIETOILA AND JAMES F. GIBBONS
Thus t h e SIMS technique cannot provide a b s o l u t e q u a n t i t a t i v e i n f o r m a t i o n about t h e number of i m p u r i t i e s i n t h e sample i n quest i o n . If t h a t i n f o r m a t i o n i s needed (e.g., i f t h e implanted dose i s n o t a c c u r a t e l y known), RBS a n a l y s i s must be performed. This method i s e x p l a i n e d i n t h e f o l l o w i n g s e c t i o n .
3.2.3
RBS and MeV Ion Channeling Analysis
A s w a s i n d i c a t e d above, t h e Rutherford b a c k s c a t t e r i n g and channeling techniques can be used t o determine a b s o l u t e numbers of i m p u r i t i e s i n semiconductors, t h e i r l a t t i c e l o c a t i o n s Unfortunately t h e and c r y s t a l l i n e q u a l i t y of t h e h o s t l a t t i c e . s t u d y of i m p u r i t i e s must, i n p r a c t i c e , be l i m i t e d t o atoms which In a r e h e a v i e r t h a n t h e h o s t atoms, a s w i l l be e x p l a i n e d below. t h e f o l l o w i n g w e w i l l g i v e a s h o r t d e s c r i p t i o n of t h i s technique i n g e n e r a l , and t h u s show how i t can be used f o r t h e t h r e e abovementioned a p p l i c a t i o n s .
This s e c t i o n i s aimed t o h e l p a r e a d e r u n f a m i l i a r w i t h t h i s method t o b e t t e r understand t h e r e s u l t s p r e s e n t e d l a t e r in t h i s c h a p t e r . Those wishing a n in-depth d e s c r i p t i o n of t h e backscatt e r i n g technique should c o n s u l t Ref. 3 . 2 2 .
3.2.3.1
General D e s c r i p t i o n
A schematic diagram of a t y p i c a l b a c k s c a t t e r i n g s p e c t r o A Van de Graaff a c c e l e r a t o r metry system i s g i v e n i n Fig. 3.4. Momentum g e n e r a t e s a m n o e n e r g e t i c MeV i o n beam of He i o n s . a n a l y s i s i s made by a magnet t o o b t a i n a beam of 4He n u c l e i a t a w e l l d e f i n e d energy. This beam i s then c o l l i m a t e d by a s e t of s l i t s , e n t e r s t h e t a r g e t chamber and impinges upon t h e sample t o be analyzed. The i n c i d e n t p a r t i c l e s c o l l i d e w i t h t h e l a t t i c e and impurity atoms of t h e sample and some of them are b a c k s c a t t e r e d from t h e near s u r f a c e r e g i o n through a s i n g l e e l a s t i c c o l l i s i o n . A s o l i d s t a t e d e t e c t o r i s used t o c o l l e c t and a n a l y z e t h e energy d i s t r i b u t i o n of b a c k s c a t t e r e d p a r t i c l e s f o r a given d i r e c t i o n . For a t a r g e t s u r f a c e p e r p e n d i c u l a r t o t h e i n c i d e n t beam, any d i r e c t i o n l o c a t e d a t an angle g r e a t e r than 90" from t h e incoming beam can be s e l e c t e d t o c o l l e c t t h e b a c k s c a t t e r e d 4He p a r t i c l e s . A common choice i s t o p l a c e t h e d e t e c t o r a t 170' where t h e backs c a t t e r e d 4He p a r t i c l e s can be considered as having passed through t h e same d e p t h of analyzed m a t e r i a l on t h e i r way i n and o u t of t h e sample. D e t e c t o r s may a l s o be placed a t g r a z i n g a n g l e s when b e t t e r depth r e s o l u t i o n i s needed i 3 . 2 3 1 . The d e t e c t o r i s t y p i c a l l y d e v i c e which i s n o t s e n s i t i v e t o p a r t i c l e . The e l e c t r i c a l s i g n a l detector is proportional t o the
a solid s t a t e surface barrier t h e charge s t a t e of t h e c o l l e c t e d generated i n t h e s u r f a c e b a r r i e r i n c i d e n t p a r t i c l e energy. This
IMASS SELECTION VAN DE GRAAFF ACCELERATOR (MODEL K )
3 MeV MFLECTION
MULTI- CHANNEL ANALYSER
U
BY'\'. 9
'+' I
GRAZING DETECTOR ( ~ - 8 3 0 )
/ /
SCHEMATIC DIAGRAM OF THE BACKSCATTERING SPECTROMETRY SYSTEM
11s
A R T 0 LIETOILA AND JAMES F. GIBBONS
signal is amplified at several stages and then stored in a multichannel analyzer. The final form of the data shows the number of backscattered events per channel as a function of energy. This spectrum can then be displayed on a plotter or printer. The sample to be analyzed is held on a precision goniometer (typically 0.01" accuracy). Several constructions exist, with the typical system providing the following three degrees of freedom: a 360" rotation perpendicular to the incident beam, -+ 90' sample tilt from its normal to the beam direction, and vertical translation. 3.2.3.2
Basic Physical Process: Shape of a Spectrum
A s indicated above, the output of anRBS system is the number of backscattering events vs energy of the backscattered particles. The spectrum is digitized in the multi-channel analyzer to give the number of particles in each channel (typically 3-5 keV wide).
There are three factors governing the shape of an RBS spectrum for a given incident particle energy and flux. 'he kinematic factor, €$,, determines the fraction of energy the projectile retains in the elastic collision. The probability of collision yielding backscattering to the detector angle is proportional to the scattering cross section, u. Finally, the projectile loses energy when it travels inside the target material; this energy loss is described by the stopping cross section, E. The kinematic factor depends on the backscattering angle and the mass ratio of the projectile and the target atom. For a constant projectile mass, the kinematic factor increases with increasing target mass. This is why trace impurities in Si can only be detected if their mass is larger than that of Si; in this case the signal from the impurities will be well separated from the silicon signal. The scattering cross section for 4He particles is proportional to the square of the target atom number. This is another reason why heavier impurities are more easily detected. The stopping cross section, which determines how much energy the projectile loses while traveling in the target material, increases when the projectile energy decreases. This causes nonlinearity in depth vs energy calibration. (The concentration of implanted impurities in Si is always so low that the energy loss of the projectile is determined by Si material only.) An example of the backscattering spectrum of 2.0 MeV 4He+ The sample analyzed was crystalions is given in Fig. 3.5 L3.241. line Si implanted with a nominal dose of 1.2 x 1015 As atomsfcm2
3.
ION IMPLANTED CRYSTALLINE SILICON
119
at 250 keV. The plot gives, point by point, the number of counts in each channel (5.0 keV channel width). The silicon signal is a step with leading edge at 1.13 MeV. The particles scattered from inside the sample have a lower energy since energy is lost in going through the material. The As signal (plotted on amplified scale) appears at higher energies (1.5-1.6 MeV), since the kinematic factor is higher for As (mass = 75 amu) than for Si (mass = 28 amu). The As-signal is relatively higher than the Si-signal because the scattering cross-section is about a factor of 6 higher for As.
Energy ( M e V )
FIGURE 3.5. Ener spectrum of 2 MeV 4He+ ions backscattered from Si implanted to aY5As dose of 1 . 2 ~ 1 0c 2~ ~ at 250 keV. The ' m vertical arrows indicate the energies of particles scattered from surface atoms of 28Si and 75As (from Ref. 3.24). 3.2.3.3
Total Number of Impurities
The total number of impurity atoms in the Si sample (i.e., the implantation dose) can be determined directly by RBS techniques. The impurity dose, Ni, is given by (Ref. 3.22, p.139):
where Ai is the total number of counts (area) of the impurity signal and HSi is the height of the Si-signal at the edge; &is the width of a channel in the multi-channel analyzer; osi and oi are
120
ART0 LIETOILA AND JAMES E GIBBONS
the scattering cross sections for Si and the impurity atoms, respectively; and [ E ~ I S is ~ the stopping cross section factor for the He particles in Si. HSi is a machine parameter, and us$, ui, and [ E ~ ] Sare ~ tabulated in Ref. 3.22. The RBS technique thus yields an absolute value for Ni. Unfortunately, this method is applicable only to impurities of higher mass than Si. 3.2.3.4
Channeling Technique: Thickness of Amorphous Layer, Crystalline Quality and Lattice Location of Impurity At oms
The high resolution goniometers of a backscattering system can be used to align the principal crystalline orientation of the sample with the direction of the incident 4He+ beam. If the sample has good crystalline quality, the yield will drop to 3-42 of the random value, since most of the host atoms lie in rows that terminate in a surface atom, and they are therefore hidden from the beam by it. This feature can be used to study the crystalline quality of the sample. An example of a series of channeling spectra is given in Fig. 3 . 6 . The Si sample was first amorphized with implantation and then subjected to low temperature thermal annealing for increasing times. The annealing caused the amorphous layer to regrow via solid phase epitaxy (SPE; see Section 3.3.2 below). The spectrum for the as-implanted sample has a high yield at energies from 1.13 MeV down to about 0.93 MeV, corresponding to particles scattered from the amorphous layer. Particles which go through the amorphous layer to the underlying crystalline material will partly channel, and the yield is markedly lowered. However, low angle scattering in the amorphous layer causes the beam to lose coherence, and the yield is higher than for a completely crystalline sample. The energy scale of the particles scattered from silicon can be translated to a depth scale (Ref. [ 3 . 2 2 1 , p. 59), as is shown in Fig. 3.6. Thus the thickness of the amorphous layer can be determined. Figure 3.6 shows the trailing edge of the amorphous Si signal moving to higher energies with increasing annealing time, corresponding to movement of the crystalline-amorphous interface closer to sample surface. Simultaneously, the yield from the underlying crystalline material is decreased because a thinner amorphous layer causes less incoherence in the helium beam. Figure 3.6 shows that the crystalline structure of the sample was completely restored after 60 minutes of annealing. The quality of the crystal can now be judged by comparing the
3.
ION IMPLANTED CRYSTALLINE SILICON
121
yield of the crystalline sample (taken just behind the surface peak) to the amorphous yield. This ratio is called the minimum yield, xmin. For a perfect Si sample of orientation, )cdn = 3.4%. The sample illustrated in Fig. 3.6 has a minimum yield of 6%; this indicates that some volume damage exists in the sample. However, a part of the high yield is also due to the high surface peak characteristic of a damaged surface layer in regrown samples. 4000 I
Omin.'
2000
I
I
I
depth
1
I5minx 30mina 45min" 0
.-......
x x
60min.
(A)
-
%
.&
*-wo 09
I
I
10
1.1
1.2
Energy (MeV)
FIGURE 3.6. Aligned spectra of 2 MeV 4He ions for Si. The sample was self implanted to a total dose of 8x1015 cm-2 at multiple energies ranging from 50 to 250 keV, and then annealed at 5OO0C for different times (0-45 min) (from Ref. 3.25). Finally, if the sample were implanted with impurities, such as As, their lattice location can be determined by the channeling technique. If the impurities are substitutional in the crystal, they are "hidden" behind host atoms, and little or no signal i s then obtained from the impurities. It is customary to calculate the ratio of the total impurity yield (signal area) to the implanted dose. If all impurities are substitutional, that ratio has the value of x d n (Ref. 3.22, p.269); anything above that indicates presence of interstitial impurities.
122 3.3 3.3.1
ART0 LIETOILA AND JAMES F. GIBBONS
CW LASER INDUCED CRYSTALLIZATION OF AMORPHOUS SILICON Introduction
It has been well established that when implantation amorphized silicon is furnace annealed at low temperatures (4006OO0C), the annealing mechanism is solid phase epitaxy (SPE) which proceeds from the amorphous crystalline interface towards the sample surface f3.251. However, simultaneously with the epitaxial regrowth, nucleation and growth of polycrystallites can take place in the amorphous layer [3.26]. When the polysilicon formation has proceeded far enough, the epitaxial regrowth will stop. Thus there is a limit on the thickness of the amorphous layer that can be regrown to single crystal silicon [3.27,3.28]. Since the early days of laser processing, the experiments have suggested that the mechanism of cw laser annealing of amorphous layers is SPE. However, due to the extreme rapidity of the phenomenon (annealing typically takes place in milliseconds or less), a clear proof for this has only recently been developed. Olson et al. [3.29] were able to observe in situ the cw laser induced regrowth of amorphous silicon. Their measurement technique utilizes interference effects within the amorphous layer, which has different optical properties than the underlying crystalline material. The regrowth can be observed by monitoring the reflectivity of the sample which is being annealed with a stationary beam. In addition to SPE, the nucleation and growth of polycrystalline silicon has also been observed in laser processed amorphous layers. Gat [3.27] has shown that an attempt to anneal a deep 4000 a thick) self implantation amorphized layer with a Kr+ cw laser leads always to formation of a polycrystalline layer (of unknown thickness) on top of the sample. Hess et al. i3.281 have found that while thin ( - 1200 a thick) UW-evaporated amorphous layers could be epitaxially regrown (in solid phase), 5000 a thick layers showed development of polycrystalline material. Little further characterization of the polycrystalline layer was performed in either of the above experiments. (13
Studies on the cw laser induced regrowth of amorphous silicon are the subject of this section. First we will give a short review of the kinetics of SPE as observed in low temperature furnace processing. m e n we report on the observation of SPE and measurement of regrowth rates under typical laser annealing conditions. Finally, we give results on regrowth and nucleation of polycrystalline silicon in deep amorphous layers.
3. 3.3.2
ION IMPLANTED CRYSTALLINE SILICON
123
Low Temperature Regrowth Kinetics
The kinetics of SPE in the temperature range of 450-600°C have been extensively studied by Csepregi and coworkers [3.25, 3.30, 3.311. The most immportant features are as follows. The regrowth proceeds at a constant rate from the amorphouscrystalline interface towards the sample surface for all crystalline orientations except <111>. The latter exhibits two distinct growth velocities, the first 1000 A growing considerably more slowly than the rest of the layer t3.311. The growth rate vg is thermally activated, obeying the formula vg
=
vgo exp (-Ea/kT)
(3.7)
The activation energy, Ea, was measured to be, for all crystal orientations, 13.311: Ea
=
2.35
f
0.1 eV
(3-8)
The pre-exponential terms for the most interesting crystal orientations are 13.311:
f
= 1.4
1014 A i s
a10>
=
0.2
1014 a i s
a i i > , slow
=
0.35x 1014
ais
a i l > , fast
(3.9)
A recent measurement on Si gave the values of 2.85 0.1 eV for Ea and 3.68~10~’A/s for vgo.
The growth rate can be greatly enhanced by doping the amorphous layer with electrically active impurities, such as As, B, or P [3.30]. The n-type dopants do not change the activation energy, whereas B, while enhancing the regrowth still more than P and A s , appears to decrease the activation energy. Table 3.1 gives the enhancement of the regrowth caused by P-doping at different concentrations. 3.3.3
Observation of cw Laser Induced SPE
Williams et al. [3.2] were the first to find evidence that the mechanism of cw beam induced crystallization of amorphous Si is SPE. These authors used glancing angle ion channeling measurements to measure amorphous layer thickness in 75As-implanted and cw laser annealed silicon samples. They found that under suitable
ART0 LIETOILA AND JAMES F. GIBBONS
124
c o n d i t i o n s , t h e amorphous-crystalline i n t e r f a c e had moved a s h o r t d i s t a n c e towards t h e sample s u r f a c e , w h i l e a h i g h e r l a s e r power r e s u l t e d i n complete regrowth of t h e amorphous l a y e r ; t h i s beh a v i o r i s ' i l l u s t r a t e d i n Fig. 3.7. TABLE 3.1. The E f f e c t of Phosphorus Doping C o n c e n t r a t i o n on t h e Rate of S o l i d Phase E p i t a x i a l Regrowth of S i l i c o n a t 475°C ( a f t e r Ref. 3.30).
Enhancement factor
Regrowth rate (Urnin)
Doping concentration ( 1020 0 - 3
-
3.25
0.93
8.3
1.10 1.60 2.20 2.80
9.1 12.5 1 7 .O 19.3
j
30keV As-SI
1
DOSE 3 r 1 0 ' 5 ~ sCm-*
(100)
1 9 MeV He'
CWARGON LASER ANNEALED
i w 2
2.6 2.8 3.8 5.2 5.9
.
('gp)
+!z-
.
SOLID STATE DETECTOR
( A ! St REGROWTt
280
300
320
340
360
380
CHANNEL NO ( 4 . 2 keV/CHANNEL!
FIGURE 3.7. Glancing a n g l e RBS random and a l i g n e d s p e c t r a f o r s i l i c o n which w a s implanted w i t h A s at 30 keV and annealed w i t h a s c a n n i n g A r + laser. The a l i g n e d s p e c t r a a r e f o r ( a ) asimplanted, (b) p a r t i a l l y a n n e a l e d , and ( c ) f u l l y annealed samples (from Ref. 3 . 2 ) .
3.
ION IMPLANTED CRYSTALLINE SILICON
125
While Williams e t a l . 13.21 found only one p o i n t of i n t e r mediate s t a g e i n t h e regrowth p r o c e s s , t h e technique introduced by Olson e t a l . [3.4] allowed i n s i t u o b s e r v a t i o n of t h e growth of t h e e n t i r e l a y e r . T h e i r technique i s based on t h e i n t e r f e r e n c e e f f e c t s w i t h i n t h e amorphous l a y e r which has d i f f e r e n t o p t i c a l p r o p e r t i e s t h a n t h e underlying c r y s t a l l i n e material. That i n t e r f e r e n c e makes t h e sample r e f l e c t i v i t y depend p e r i o d i c a l l y on t h e t h i c k n e s s of t h e amorphous l a y e r , as i l l u s t r a t e d i n F i g . 3.8. Thus, by monitoring t h e r e f l e c t i v i t y of a sample being annealed w i t h a s t a t i o n a r y laser beam, i t i s p o s s i b l e t o observe t h e regrowth. Figure 3.9 shows an example of a r e f l e c t i v i t y vs t i m e p l o t measured by Olson e t a l . 13.41, which c l e a r l y demonstrates t h a t t h e regrowth is proceeding from t h e o r i g i n a l i n t e r f a c e towards t h e sample s u r f a c e .
0.490 0.470 I>
-5
0.450
Iu
0.430 w
a 0.410 0.390 0.370 L 4QM)
1
3000
2000
loo0
0
AMORPHOUS FILM THICKNESS, A
FIGURE 3.8. Calculated r e f l e c t i v i t y ( a t 0.6328 um) of amorphous Si f i l m on c r y s t a l l i n e S i s u b s t r a t e a s a f u n c t i o n of f i l m t h i c k n e s s . (Note: Thickness scale is reversed t o s i m p l i f y comparison w i t h experimental d a t a , F i g . 3.9 .) The r e f r a c t i v e f o r amorphous S i i n d i c e s used i n c a l c u l a t i o n s are 4.85-0.6121 01981 Elsevier and 4.16-0.0181 f o r c r y s t a l l i n e S i (from Ref. 3 A ) . Science Publishing Co.
3.3.4
Rate of CW Laser Induced Regrowth
L i e t o i l a e t a l . 13.321 have measured t h e e p i t a x i a l regrowth rates under a c t u a l cw laser a n n e a l i n g c o n d i t i o n s , u s i n g a scanning beam and l a s e r powers t h a t y i e l d a maximum s u r f a c e temperature c l o s e t o 1000°C. This was accomplished by determining t h e s c a n speed of a s i n g l e l a s e r s c a n r e q u i r e d t o regrow a known amorphous l a y e r a t t h e c e n t e r of t h e beam. The regrowth r a t e s can be d e t e r mined by u s i n g t h e formalism p r e s e n t e d i n Chapter 2 which g i v e s t h e e f f e c t i v e dwell t i m e f o r a scanning Gaussian beam. S p e c i a l
ART0 LIETOILA AND JAMES F. GIBBONS
126
attention was paid to the determination of the annealing temperature: one of the precautions was to leave the laser power unchanged after melting power calibration and regulate the annealing temperature by changing the sample holder temperature. This experiment is described below. 0.49
7-rc '
0.47
- 0.45 1 $ w
0.43
3
U
0.41 0.39 0.37
o,35i
I
i
0
_10
i . . 5
~
-.._
15
-
20
d i
A 25
TIME, sec
FIGURE 3.9. The reflectivity of a HeNe probe vs time for a Si sample which is being annealed with a stationary Ar' laser turned on at t = 0). The sample is Si implanted to a j 5 A s dose of 5x1014 at 158keV. The laser power is 7.5W and substrate temperature To = 460°C (from Ref. 3.4). 9981 Elsevier Science Publishing Co.
3.3.4.1
Sample Preparation
The samples were of polished Si. The amorphous layer was formed by self implantation to avoid impurity effects. The implantation energy and dose were 40 keV and l ~ l O ~ ~ c r nrespec-~, tively. The samples were held at LN2 temperature during implantation to prevent any self annealing. The thickness of the as implanted amorphous layer was measured by using 2.2 MeV4He+ channeling techniques (see Section 3 . 2 ) , and was found to be 960 8 . Prior t o laser annealing,the samples were cleaned in methanol and D.I. water, the purpose of which is mainly to remove major dust particles. Annealing was performed in the laboratory atmosphere.
3. 3.3.4.2
ION IMPLANTED CRYSTALLINE SILICON
127
Measurement of the Laser Beam Diameter
The cw laser scanning system described in Chapter I was used with a 136 mm focusing lens. The focussed beam radius was measured by determining the power required to reach surface melting of Si with the substrate temperature To = 332°C. A scanning speed of 6 cm/s was used, which is slow enough to allow the steady state temperature calculations (Chapter 11) to be used. The effect of scanning speed on temperature will be discussed later. The incident power (discounting the losses in the beam shaping devices) required to cause surface melting of crystalline Si was found to be 5.0 W. The parameter pa [see Eqs. (2.13,2.14)] has to be 1210 W/cm to reach the maximum temperature of 1412OC for To = 332°C [Eqs. (2.7,2.13)]. 'Ihe reflectivity of crystalline Si was measured to be 38% for the Ar+ laser radiation; thus w = 26 pm is obtained. We would like to stress that the beam radius determined above is used only to compute the dwell time of the spot. The calculation of the annealing temperature is based solely on the calibration against the melting power. 3.3.4.3
Procedure to Measure the Regrowth Rates
After the laser power required to reach melting was determined and the beam radius was calculated, the laser was kept at that power. The annealing temperature was then controlled by changing the substrate temperature, To. This was done because the beam radius i s known to vary with laser power and a constant beam radius is essential for accurate determination of the annealing temperature. For each substrate temperature To, two scanning speeds, v1 and v2, were determined so that the sample region underneath the center of the beam was just regrown at vl, whereas v2 was too fast to allow full regrowth. The results of the scan were judged optically: the regrown area appears dark against the lighter amorphous background [3.3]. The values of To as well as v1 and v2 are given in the first two columns of Table 3.2. 3.3.4.4
Analysis of Results
The peak annealing temperatures for each substrate temperature, To, were obtained using the calculation presented in Chapter 11. The scan speeds were quite low suggesting the use of Eqs. (2.7,2.13 ) derived for a stationary beam. However, first order corrections to the temperature due to scan speed were made by
128
ART0 LIETOILA AND JAMES F. GIBBONS
c a l c u l a t i n g t h e l i n e a r temperature 0 f o r t h e scanning beam [Eqs. (2.21,2.12)] and applying t h e i n v e r s e Kirchhoff t r a n s f o r m [Eq. (2.7)1 * The Kirchhoff transform i s not q u i t e e x a c t l y v a l i d i f V ( a parameter r e l a t e d t o t h e scanning v e l o c i t y : see Chapter 11) d e v i a t e s from z e r o . However, V g e t s q u i t e small i n t h i s case: t h e change i n t h e temperature due t o scanning w i l l be a t most about 20°C, s o t h i s i s not a c r i t i c a l p o i n t . The peak a n n e a l i n g temperatures are now determined a s f o l lows. Since t h e melt temperature c a l i b r a t i o n w a s performed on c r y s t a l l i n e s i l i c o n , w h i l e t h e a c t u a l samples were amorphous, t h e m e l t i n g v a l u e , 1210 W/cm, f o r p has t o be m u l t i p l i e d by t h e r a t i o of a b s o r p t i v i t i e s (1-R) of amorphous and c r y s t a l l i n e s i l i con. The r e f l e c t i v i t y of t h e samples used h e r e w a s measured t o be 4 5 % , s o t h e e f f e c t i v e p i s 1073 W/cm f o r t h e s e samples. For each s u b s t r a t e temperature and scanning speed, t h e parameter V i s c a l c u l a t e d u s i n g Eq (2.21). Equation (2.22) g i v e s t h e c o e f f i c i e n t s n , which determine how much t h e l i n e a r maximum temper,, i s reduced from t h e value f o r a s t a t i o n a r y b e a m . ature @ Equation (2.13), with p = 1073 W/cm, is used t o c a l c u l a t e t h e Om, stat i s t h e n d i v i d e d by t h e v a l u e of s t a t i o n a r y OmX,Stat. n W l t f o r t h e m e l t i n g c a l i b r a t i o n scan (.986, see Table 3 J), and m u l t i p l i e d by t h e corresponding value of n f o r each s c a n . The r e s u l t i n g v a l u e of 0 i s f i n a l l y converted t o r e a l temperature The r e s u l t s are g i v e n i n Table 3.2. using Eq. ( 2 . 7 ) .
.
Next t h e e f f e c t i v e dwell times are determined u s i n g t h e f o r m a l i s m p r e s e n t e d i n S e c t i o n 2.5. For any thermally a c t i v a t e d process t h e e f f e c t of a s i n g l e laser s c a n a t t h e c e n t e r of t h e scan l i n e can be d e s c r i b e d u s i n g a c o n s t a n t temperature, equiv a l e n t t o t h e peak temperature, Tmx, and a reduced dwell t i m e , teff : ( 3 .lo) The dwell time r e d u c t i o n f a c t o r f w a s shown i n S e c t i o n 2.5 t o depend mainly, but s t i l l n o t s t r o n g l y , on t h e s u b s t r a t e temperat u r e and t h e a c t i v a t i o n energy of t h e p r o c e s s . Values of f a r e given i n F i g . 2.14. I n t h i s case, where t h e s u b s t r a t e temperaW e used t h e curve f o r E, = t u r e s a r e around 200°C, f = 0.26. 2.35 e V i n F i g . 2.14; t h i s c h o i c e is n o t c r i t i c a l , however, s i n c e The e f f e c t i v e dwell-times c a l c u l a t e d f depends only weakly on Ea. f o r each s c a n a r e given i n Table 3.2. The l a s t s t e p i n t h e d a t a r e d u c t i o n i s t h e c a l c u l a t i o n of t h e regrowth r a t e s u s i n g t h e known t h i c k n e s s of t h e amorphous l a y e r and t h e e f f e c t i v e dwell t i m e s . For each s u b s t r a t e temperat u r e we o b t a i n a lower and a n upper l i m i t f o r t h e regrowth rate,
TABLE 3.2.
a)
Results of the Study of the CW Laser Induced SPE Regrowth of Si
174
1.5 2.3
0.86
0.003 0.005
0.998 0.996
357.0 356.3
797 795
0.88 0.57
1.1 1.7
184
1.5 3.0
0.84
0.003 0.007
0.998 0.995
367.3 366.2
825 822
0.88 0.44
1.1 2.2
193
3.0 4.5
0.81
0.007 0.010
0.995 0.993
375.3 374.6
847 845
0.44 0.29
2.2 3.3
202
3.0 6.0
0.80
0.007 0.014
0.995 0.990
384.6 382.6
872 866
0.44 0.22
2.2 4.4
211
6.0 9.0
0.78
0.014 0.021
0.990 0.985
391.8 389.6
89 1 886
0.22 0.15
4.4 6.4
332
6.0
0.59
0.019
0.986
578.1
1412
---
---
v1 is the velocity at which complete regrowth was observed at the center of the beam; for v2, the crystalline interface did not reach the sample surface.
130
ART0 LIETOILA AND JAMES F. GIBBONS
corresponding to scan velocities v1 and v2, respectively. The lower and upper limits for the regrowth rate vs temperature are given in Table 3.2 and in Fig. 3.10. In the figure a bar is drawn to connect the limits obtained at each substrate temperature. For comparison, Fig. 310 also shows the low temperature data given by Csepregi et al. (3.25, 3.30, 3.311 and in Ref. 3.59. Figure 3.10 shows that the regrowth rates induced by laser annealing are one to two orders of magnitude higher than the rates extrapolated from furnace annealing data. This difference is too large to be explained by the uncertainties in the observed regrowth velocities (a factor of 1.5 - 2 ) or in the calculated temperatures (-- 2 20°C). (The observed data are, however, too scattered to make any firm conclusions regarding the activation energy.) Thus, these measurements would indicate that the epitaxial regrowth is enhanced by the laser radiation. It should be emphasized however, that the extrapolation of growth rates from the low temperature data (see Fig. 3.10) is subject to considerable uncertainty, since the measurements performed in Refs. 3.25, 3.30, 3.31, and 3.59 were limited to a narrow temperature range. 3.3.5
Obstructed Regrowth of Thick Amorphous Layers
Hess et al. [3.28] have observed that under their experimental conditions, UHV deposited 0.5 um thick amorphous Si layers could not be regrown with a scanning cw laser. 'Ihis was shown to be a problem stemming from the layer thickness rather than layer preparation, since 0.1 um thick, identically deposited layers grew fully with laser scanning. Gat also observed 13.271 that implantation amorphized, 0.4 pm thick Si layers did not fully regrow but turned at least partly polycrystalline. Hess et al. [3.28] postulated that spontaneous nucleation and growth of polycrystalline silicon would stop the epitaxial regrowth in thick layers. They suggested that doping of the amorphous material would make it possible to regrow thicker layers, since doping is known to enhance the epitaxial regrowth rate [3.30]. The latter hypothesis is very speculative: the growth of polycrystalline silicon may also be enhanced by doping. Also, it is not known whether the nucleation of polycrystallites originates at the sample surface or in the bulk of the amorphous material. In the former case, the ratio of the epitaxially regrown and polycrystalline layer thicknesses should be about the same if the thickness of the original amorphous layer is varied. But if the nucleation takes place in the bulk, the regrown thickness should stay constant, and the thickness of the
3
ION IMPLANTED CRYSTALLINE SILICON
TEMPERATURE ("C) 700 600
1000 900 800 I
I
I
llkT
131
500
(d')
FIGURE 3.10. The rate of cw laser induced solid phase epitaxial regrowth in (100) Si (after Ref. 3.32). For comparison, the low temperature growth data given in Refs. 3.25 and 3.59 are also given.
remaining polycrystalline layer should vary with the original amorphous layer thickness. Lietoila 13-33] has conducted a study which sheds some light on the above questions. Different amorphous layer thicknesses, doped and undoped, were used to determine where the spontaneous nucleation originates and the effect of doping. Results of this study are reported in the following.
3.3.5.1
Sample Preparation, Experimental Techniques
Silicon wafers of orientation were amorphized by ion implantation. Three different implantation schedules were carried
ART0 LIETOILA AND JAMES F. GIBBONS
132
out to form two different thicknesses of undoped amorphous silicon and one P-doped layer. The implantation schedules are given in Table 3.3. The wafers were held at LN2 temperature during implantation; and all implants were performed in the same implanter. The laser annealing was performed using the system described in Chapter I with a 160 mm focusing lens. 'he substrate temperature was held constant, To = 332OC, and the processing temperature was controlled by changing the laser power. The beam radius for the different laser powers was determined by measuring the power required to melt the sample surface for elevated substrate temperatures (up to To = 530°C). This procedure is described in Section 3.3.4.2. and it was carried out to make the determination of the annealing temperature more accurate. The scan rate was kept constant, v = 2.7 cm/s, and the distance between adjacent scan lines was 10 um in all cases.
"he thicknesses of the amorphous layers were measured by 2.2 MeV 4He+ channeling techniques as described in Section 3.2. For calculation of the layer thicknesses, the symmetrical mean energy approximation was used. 3.3.5.2
Undoped Amorphous Layers
The effect of laser power on the regrowth thickness was studied using samples subjected to the implantation schedule I (see Table 3.3). nose implants result in a continuous amorphous layer of 3740 8, as measured by channeling; the channeling spectrum is given in Fig. 3.11. TABLE 3.3. Implantation Schedules for the Study of Laser Induced Epitaxial Regrowth. All implants were performed at LN2 temperature (after Ref. 3.32). ~
Schedule
~
Ion
Energy (keV)
Dose (cm2)
I
si
25 50 90 170
3.0~10~~ 6.0~1014 1.1~1015 2.8~1015
I1
si
100
l.OXlOl6
25
3.0~1014 6.0~1014 1 .iX1015 2.8~1015
I11
P
50
90 170
3.
ION IMPLANTED CRYSTALLINE SILICON
133
The laser powers used, the resulting maximum annealing temperatures, and results of regrowth are given in Table 3.6. The channeling spectra for a few cases are also given in Fig. 3.11. The fact that the channdJing yield behind the random peak drops clearly as the amorphous layer gets thinner means that the regrown layer has good crystalline quality. DEPTH ( p m )
0.8
0.4
1
0
~
SCHEDULE
1
~
1
I Tmax
990OC 117OoC 1320O C
0.8
1.0
1.2
1.4
6
ENERGY (MeV)
FIGURE 3.11. Channeling spectra for Si samples implanted according to Schedule I (see Table 3.3) and laser annealed at powers resulting in different maximum temperatures TmX. The solid line is the as-implanted spectrum (from Ref. 3.33).
In Table 3.4 are also included the calculated maximum thickness of regrowth, obtained using the extrapolated low temperature regrowth kinetics measured by Csepregi et al. [3.31] The thickness was calculated at the beam center, including the contributions of the two adjacent scans; the same dwell time, 0.6 m s , was assumed for each. This is not quite correct, since the contributions of the adjacent scans come from slightly off the beam center. However, since the overlapping is tight (1. 83%), the error is not significant.
.
Table 3.4 shows that throughout the power range used, the thickness of regrowth is almost constant. The thickness of the remaining polysilicon layer decreases slightly with increasing laser power, but even for the highest power, more than 800 a of Si has turned polycrystalline. That power resulted in a heavy network of sliplines on the sample surface, which is why no higher power was employed.
ART0 LIETOILA AND JAMES E GIBBONS
134
TABLE 3.4. Results of Laser Annealing of the Deep Siimplanted Amorphous Layer (Schedule I) (after Ref. 3.33).
Amorphous/ PolySi Layer(A)
-
-
0.65 0.71 0.78 0.84 0.91
990°C 1070°C 1170°C 1240°C 1320°C
Regrown Layer(A)
-
3740 1440 1300 1230 1120 830
2300 2440 2510 2620 2910
Calculated Regrowth(A) Cspregi et al. [3.25,3.31] Ref. 13.391
-
180
680 2890 6430 15700
1810 9000 5 .2x104 1.4~105 4.1~105
The question as to whether the formation of polycrystalline silicon originates at the sample surface or in the bulk of the amorphous material was studied by laser scanning wafers implanted according to Schedule I1 (see Table 3.3). This implant results in a continuous amorphous layer which is only 2570 a thick or about the same as the regrowth for the above samples of Schedule 1. If the spontaneous crystallization originated at the sample surface, one would expect the ratio of the random and regrown layers to be the same as for the above thicker amorphous layers. But, if the nucleation started homogeneously in the bulk of the amorphized material, the regrowth of the thinner amorphous layers (Schedule 11) should be complete. Table 3.5, and Fig. 3.12, show that the regrowth of the samples, which originally had 2570 a of amorphous material, is almost complete. The remaining layer is too thin to be measured by channeling, since the FWHM of the surface peak is about the same as the system resolution (7 channels or 25 keV). However, we can say that it is at most 200 A , most likely less. TABLE 3.5. Channeling Results of As-implanted and Laser Annealed Wafers with the Shallower Si-implanted Amorphous Layer Schedule (11). Results for a virgin sample are included for comparison (after Ref. 3.33). P/Pme1t
TmaX
- Virgin Sample 0.78 0.84
1170°C 1240°C
Amorphous/ Polysilicon Layer(A)
-
2570 <200 <200
Minimum Yield
X,X
Surface Peak XO
4.4%
8.8%
7.6% 5.8%
34% 28%
-
-
3.
ION IMPLANTED CRYSTALLINE SILICON
135
DEPTH ( p m )
0.8
0.4 JI
SCHEDULE
lO0Ot
1 1 I U
0
n
Tmax
-.----
117OoC 1240 O C
ENERGY (MeV)
FIGURE 3.12. Channeling spectra for Si samples implanted The solid line is the asaccording to Schedule I1 (Table 3.3). implanted spectrum (from Ref. 3.33). The above results would thus suggest that the spontaneous nucleation of crystallites is a bulk, rather than surface phenomenon, This would mean that amorphous layers below a critical thickness can be completely regrown, whereas attempts to regrow more will result in only the critical thickness being crystalline: the rest will be polycrystalline material. The critical thickness will vary with laser power and other conditions, which is why the quantitative results given here should be used as guidelines only. 3.3.5.3
Phosphorous Doped Amorphous Layers
The possible effects of doping on the thickness of regrowth, were studied by annealing samples amorphized with a multipleenergy phosphorus implantation (Schedule 111, Table 3.3) The result of that series of implants is a relatively flat Pconcentration of about 1.7~10~0cm-3, as is shown in Fig. 3.13 (calculated LSS-profile). This concentration would enhance the epitaxial growth rate by a factor of four at 475OC 13.301. The laser scanning was performed under exactly the same conditions as in the previous section, using two different laser powers yielding the maximum annealing temperatures of 117OoC and 1270°C. The results are given in Table 3 . 6 and Fig. 3.14, and show that little improvement in the regrowth has taken place compared to undoped samples.
ART0 LIETOILA AND JAMES F. GIBBONS
136
102'
1
I
CALCULATED 155 PROFILE
FIGURE 3.13. The calculated (LSS) 31P-profile for the multiple-energy phosphorus implant (Schedule 111, Table 3.3).
Thus the doping will not markedly increase the thickness which can be regrown with a scanning cw laser. This is not surprising, since the doping will enhance the growth of the polycrystallites as much as the epitaxial regrowth, and the impurity atoms might also serve as additional nucleation sites. 3.3.6
Summary
The solid phase epitaxial regrowth of amorphous Si can proceed under typical laser annealing conditions at rates which are at least an order of magnitude higher than those extrapolated from low temperature furnace annealing data. TABLE 3.6. Results of Laser Annealing f o r the P-Implantation Amorphized Samples.
PfPmelt
0.78 0.84
Tmax
117OOC 1240" C
Amorphous1 Polysilicon Layer ( A ) 3360 710
470
Regrown byer(A)
2650 289 0
3 ION IMPLANTED CRYSTALLINE SILICON I
137
DEPTH ( p m )
0.8
0l
0.4 I
SCHEDULE
IU
l
Tmax
--- 117OoC -.-
0.8
1 .o
1.2
1240°C
1.4
1
6
ENERGY (MeV)
FIGURE 3.14. Channeling spectra for the 3lP-implanted samples (Schedule 111, Table 3 . 3 ) . The solid line is the asimplanted spectrum (from Ref. 3 . 3 3 ) . Attempts to laser anneal thick amorphous layers result in formation of polycrystalline silicon which stops the epitaxial regrowth. Comparison of regrowth data for different amorphous layer thicknesses would indicate that the spontaneous nucleation of polycrystallites takes place in the bulk of the amorphous material rather than at the sample surface. Doping does not markedly increase the maximum thickness of regrowth. In the experiments reported in the previous sections, the maximum thickness of regrowth was between 2300 and 2900 8 , depending on laser power. This number is, however, likely to vary with sample preparation and other experimental conditions. Thus the results presented here give only the order of magnitude of the maximum regrowth. 3.4 3.4.1
ANNEALING OF 75As IMPLANTED Si: C
<
Css
Introduction
Medium dose 75As implanted Si samples are suitable for studies of the physical phenomena related to cw beam annealing. A 75As dose of -1~1015cm-2 is sufficient to create an amorphous layer on the sample surface, and yet the as-implanted concentration is below the solid solubility, Css, of A s in Si at typical processing temperatures. The benefit of using As is that RBS
ART0 LIETOILA AND JAMES F. GIBBONS
138
and channeling techniques can be applied. In this section, results are presented for the use of cw lasers and e-beams to anneal these medium dose 75As implants. 3.4.2
Use of
a
CW Laser
A cw Ar ion laser has been applied to annealing of medium dose As implants by Gat et al. [3.11, Williams et al. [3.2] and Auston et al. 13.341. The principal conclusion of the experiment by Williams et al. (3.21 was that the implantation amorphized layer regrows completely via SPE. This was established by channeling technique, as described in the previous section. In the following, we report on the other two experiments. 3.4.2.1
TEM,
SIMS
and
Electrical Analysis of Annealed
Samples TEM, SIMS and electrical studies were performed on ptype si samples which were 75As implanted at 100 keV to a dose of 5x10L4 cm-2 [3.1]. This implant creates an amorphous layer about 1000 vm thick. The samples were laser annealed using the system described in Chapter 1 with a 79 mu lens. The sample holder was at room temperature; the laser power providing optimum results was 7 W. Both single lines and overlapping areas were annealed. Control samples were annealed thermally at 1000°C for 30 min. in flowing N2 gas. Both laser and furnace annealed samples were analyzed by the The thermally annealed sample [Fig. 3.15(a)] contains a relatively high concentration (- 1010/cm2) of residual damage in the form of dislocation loops (- 300 average image diameter). In comparison, the concentration of defects in the laser annealed samples is considerably reduced. In the laser annealed samples, single crystal diffraction patterns [left inset Fig. 3.15(b)] were obtained. Thus the results indicate complete recrystallization of the amorphous material during laser treatment. TEM technique with the results given in Fig. 3.15.
An interesting feature in Fig. 3.15(b) is the presence of a reasonably well defined boundary (dashed line) between the recrystallized (C) and amorphous (A) zones at the edge of the laser scan line. The 75As distributions in the as-implanted as well as laser The and thermally annealed samples were determined by SIMS. results are given in Fig. 3.16 and show that, while thermal annealing has caused considerable diffusion of A s , the impurity profile in the laser annealed sample is identical to the as-implanted distribution.
3.
ION IMPLANTED CRYSTALLINE SILICON
139
FIGURE 3.15. Electron micrographs of 75As implanted Si subjected to thermal annealing at 1000°C for 30 min. (a) and laser annealing (b); inserts show diffraction patterns of the different regions (from Ref. 3.1).
140
ART0 LIETOILA AND JAMES F. GIBBONS
176
528
DEPTH (1) 1232 1584 1936
880
0 As IMPLANTED oLASER ANNEAL A THERMAL ANNEAL (1OOO"C 30min)
AA A
DISTRIBUTION WITH LSS RANGE STATISTICS
'@
40
80 120 160 200 SPUTTERING TIME (*)
240
2 0
FIGURE 3.16. R e l a t i v e c o n c e n t r a t i o n of 75As a s measured by SIMS i n t h e implanted samples a f t e r thermal and laser a n n e a l i n g (from Ref. 3.1). The s h e e t r e s i s t a n c e s i n t h e laser and t h e r m a l l y annealed i n d i c a t i n g essensamples were found t o be t h e same (180 Sa/O), t i a l l y 100% e l e c t r i c a l a c t i v a t i o n i n t h e l a s e r annealed samples. The r e s u l t s described above revealed the most s t r i k i n g complete f e a t u r e s of cw l a s e r a n n e a l i n g of i o n implanted Si: recovery of t h e c r y s t a l l i n e s t r u c t u r e and 100% e l e c t r i c a l a c t i v a t i o n w i t h no r e d i s t r i b u t i o n of dopants. 3.4.2.2
Channeling Measurements on
and
<111>S i
Auston e t a l . [ 3 . 3 4 ] used both and <111>S i samples i n t h e i r s t u d i e s . The <111>S i s a m l e s were As-implanted a t 50 keV t o doses of 1015 and l 0 l g cm-2. The (100) samples were implanted t o a dose of 1015 cm-2 a t 3 0 keV. The samples were laser annealed u s i n g an A r i o n l a s e r (Spect r a Physics Model 171) and a 100 mm f o c u s i n g l e n s . The sample w a s h e l d a t room temperature on a moving t a b l e . The a n n e a l i n g beam w a s s t a t i o n a r y . The l a s e r powers used f o r d i f f e r e n t samples a r e g i v e n i n Table 3 . 7 .
3. TABLE 3.7.
141
ION IMPLANTED CRYSTALLINE SILICON
CW Argon-ion Laser Annealing of As-implanted Si.
Orientation
As dose (cm-2)
Laser Power
<111>
1014 1015 1016 1015 10l6
9.0 9.5 9.5 8.6 8.0
Anneal strip width Xmin(Si)
Xmin(As)
Resistivity
(wn)
(%)
PQ/
37 31 30 26 30
3.5 63 13 3.5 3.5
10 61 40 6.8 10
930 375 60 131
46
Laser annealed Sam les of both crystalline orientations were analyzed by 1.9 MeV 4Hee channeling. Figure 3.17 shows representative spectra illustrating the orientation dependence of the annealing behavior; those spectra are from the samples receiving a dose of 1015 As/cm2. For the (100) sample, the strong reduction in Si yield (xmin 3.5%) shows essentially complete recovery of the crystalline structure, which is in agreement with the TEM observations of Gat et al. f3.11 (see section 3.4.2.1 above). On the other hand, poor recovery of the <111> sample is indicated 63%). by the large residual damage peak near the surface (xdn However, the spectrum given in Fig. 3.17 clearly shows that some epitaxial growth has occurred even for the <111> sample, since the crystalline-amorphous interface has moved towards sample surface.
-
-
A summary of the channeling and resistivity measurements for all laser annealed samples is given in Table 3.7. Of the three <111> samples, only the one implanted with the lowest dose annealed satisfactorily, giving a minimum Si yield of 3.5% and 90% substitutionality for As. On the other hand, both samples regrew completely, as expected.
The annealing behavior of <111> silicon can be well understood on the basis of the results given in this chapter above. It was shown in Section 3.3.2 that the low temperature epitaxial regrowth rate in the <111> direction is significantly slower than in the direction. A similar difference is expected to exist also at laser annealing temperatures. The results given in Section 3.3.5 indicated that, simultaneously with the regrowth, nucleation and growth of polycrystallites takes place in the bulk of the amorphous material. The latter process eventually stops the epitaxial growth, when the density and size of the
142
p o l y c r y s t a l l i t e s have grown l a r g e enough. The growth of p o l y s i l i c o n should n o t depend on c r y s t a l l i n e o r i e n t a t i o n , s i n c e t h e c r y s t a l s t r u c t u r e i s completely d e s t r o y e d when t h e m a t e r i a l i s amorphized. Thus, t h e amount of maximum regrowth i s expected t o Si. Indeed, Fig. 3.17 shows be s i g n i f i c a n t l y less f o r (111) t h a t only about 360 A of < 1 1 1 > S i h a s regrown, whereas t h e maximum t h i c k n e s s of regrowth f o r < l o o > S i w a s shown i n S e c t i o n 3.3.5.2 t o be almost 3000 8. The r e s u l t s p r e s e n t e d above c l e a r l y i n d i c a t e t h a t , i n p r a c t i c a l cases, only S i can be s u c c e s s f u l l y annealed w i t h a cw l a s e r , i f t h e implant dose i s heavy enough t o c r e a t e an amorphous s u r f a c e l a y e r . But i f t h e dose i s so l i g h t t h a t no amorphous l a y e r is formed, both and (111) Si can be cw l a s e r annealed.
3.4.3
Use of E l e c t r o n Beam Sources
Two t y p e s of cw e-beam s o u r c e s have been used f o r anneali n g purposes, namely an e-beam welder [3.35] and a modified SEM ( 3 . 3 6 , 3.371. The d i s t i n c t i v e d i f f e r e n c e between t h e s e two
3.
ION IMPLANTED CRYSTALLINE SILICON
143
systems is in the operating power: the welder is designed to yield currents high enough to melt sizable metal structures, whereas a SEM operates at very low currents. Thus, the beam diameter of the welder is typically about 0.1 nun or more, whereas the SEM beam has to be focused down to a few pm. The tight focusing in the SEM makes it possible to selectively anneal very small structures. In what follows, annealing results obtained with the two types of e-beam sources are reported. 3.4.3.1
The E-beam Welder
Regolini et al., [3.35] have used an e-beam welder [Hamilton Model EBW (7.5)] to anneal ion implanted Si. The samples were p-type Si, which were 75As implanted at 100 keV to dose of 1.5~105 cm-2. The wafers were held at LN2 temperature during implants. Typical e-beam parameters used for annealing were 30 keV at 0.5 mA with the beam focused to a spot approximately 300 pm in diameter. The scanning speed was 2.5 cmls, and the temperature of the sample holder remained below 5OoC during the anneal. Control samples were thermally annealed at 575OC for 30 min. Results of channeling studies for as-implanted and annealed samples are given in Fig. 3.18. The as-implanted spectrum was used to determine the implanted dose (1.5~1015 cm-2) and the thickness of the amorphous layer. The channeling spectrum for the e-beam annealed sample shows that a large number of As atoms (-25%) are on interstitial sites. 'he minimum Si yield for that sample is about 5%, indicating some residual damage in the lattice. The results for the thermally annealed sample are similar: 70% As substitutionality and 5% minimum Si yield. The modest results for the control samples can be explained by the low furnace annealing temperature, 575OC. Figure 3.19 shows the results of differential Van der Pauw measurements for the e-beam annealed samples. Plotted in that figure is also the calculated Pearson impurity distribution, from which it is apparent that no diffusion of impurities has taken place during annealing. However, the electrical activation is incomplete at the peak of the concentration profile. The total active dose is 1.02~1015 cm-2, which is 67% of the implanted dose. This number agrees well with the result from channeling, 75% substitutionality.
-
144
ART0 LIETOILA AND JAMES F. GIBBONS
RP-~’A; IMPLANT
ENERGY (MeV)
FIGURE 3.18. Aligned channeling spectra for 1.5 MeV 4He ions incident on Si implanted with 75As to a dose of 1.5~1015 cm-2 at 100 keV and then subjected to different anneals (from Ref. 3.35).
Regolini et al., [3.35] also performed TEM studies on the e-beam annealed samples, with the results given in Fig. 3.20. Selected area diffraction pattern (left insert) indicates that complete recrystallization of the amorphous layer has occurred. However, the bright field micrograph indicates that residual defects in the form of small dislocation loops (40-100 a image diameter) remain. The average loop concentration was estimated to be -1014 cm-2, which is significantly more than for the cw laser annealed samples studied in section 3.4.2.1. Removal of 280 8, of Si by anodic oxidation and etching reduced the loop density by about an order of magnitude, indicating that these loops are mostly associated with surface damage. After removal of the high concentration of loops, a distribution of rod-like structures -500 a in average length was observed. It is possible that those rods serve as precipitation sites for the As atoms. 3.4.3.2
The Modified SEX
The SEM was first applied by Ratnakumar et al., l3.371 to anneal ion implanted Si. They used a Cambridge Stereoscan S-4 scanning electron microscope which was modified f o r high current
3.
ION IMPLANTED CRYSTALLINE SILICON
145
loz1
I. 1.5~10"'5A3/cm'f
100 keV
(UNANNEALED)
FWHM.460
A
?- 10
z
w
u
-
A
L
I
I
200 400
A
I
I
I
I
6 0 0 SO; 1000 1200 DEPTH ( A )
J
FIGURE 3.19. Results of differential Van der Pauw measurements for a sample which was implanted to a 75As dose of 1 . 5 ~ 1 0 ~ ~ at 100 keV and then e-beam annealed. The solid line shows the calculated LSS profile for the 75As implant (from Ref. 3.35). operation. Even in the modified version, the current was limited to a range of 15-30 PA, the voltage being 10-30 kV. The principal results of Ratnakumar et al., r3.371 were that complete recrystallization and good electrical activity was obtained for silicon implanted to an 75As dose of 4 ~ 1 0 ~ cm-2 5 at 150 keV
.
The SEM described above was also used by Re olini et al., silicon samples which were 55As implanted to a dose of 1 . 1 ~ 1 0 ~ ~ at 100 keV. The SEM annealing conditions in this case were 50-70 M at 20 kV, with a scanning speed of 10 cmls. The estimated beam radius was 6 pm. 13.361 to anneal
Results of differential Van der Pauw measurements for the SEM annealed samples are shown in Fig. 3.21. The total active dose obtained from stripping measurement is 9 . 6 ~ 1 0 1 ~cm-2, which means that 14% of As is inactive. The partial inactivity of As was also confirmed by channeling, (see Fig. 3.22): the fraction of As atoms on non-interstitial sites was found to be 12%. Thus the activation obtained with the SEM is better than that achieved by the welder (67% activity, see Section 3.4.3.1).
146
ART0 LIETOILA AND JAMES F. GIBBONS
FIGURE 3.20. TEM micrograph of an e-beam (welder) annealed, 75As implanted sample (from Ref. 3.35).
The crystalline quality of the SEM annealed samples was studied by channeling and "EM. The minimum silicon yield (see Fig. 3.22) for the SEM annealed samples was found to be the same as for a virgin control sample, 4%. This indicates a perfect recovery of the crystalline structure. And indeed, the TEM analysis, with the results given in Fig. 3.23 shows that the annealed layer is completely crystalline with few remaining defects. 'Ihe inset in Fig. 3.22(b) shows the electron-diffraction pattern for the dark-field image; the sharpness of the Kikuchi line confirms the good crystalline quality of the layer. The bright-field micrograph [Fig. 3.22(a)] reveals, however, the presence of some dislocation loops. The largest loops are < 200 A in diameter, and the estimated area density of the loops is 1010 This is four orders of magnitude less than the defect density produced by the e-beam welder (Section 3.4.3.1). However, the residual defect density in the cw laser annealed samples is still lower (Section 3.4.2.1).
3.
ION IMPLANTED CRYSTALLINE SILICON
147
FIGURE 3.21. Electron concentration and mobility vs. depth in a 75As implanted, SEM-annealed sample. The solid lines show the calculated (LSS) 75 As distribution of the implant and the Irvin mobility calculated assuming complete activation of the asimplanted As distribution (from Ref. 3.36). 3.5 3.5.1
ANNEALING OF 75As IMPLANTED Si: C>Css Introduction
Numerous authors i3.39-3.441 have reported that pulsed laser annealing of ion implanted semiconductors can incorporate impurities on substitutional lattice sites in concentrations which greatly exceed the solid solubilities. Some authors have also observed that the supersaturated impurity concentrations are electrically active and thermally unstable [3.39,3.42-3.441. The enhanced solubilities are commonly explained by "solute trapping" of impurities [3.45]. This explanation relies on the melting hypothesis for pulsed laser annealing: when the molten layer resolidifies via liquid phase epitaxy at high speed, impurities would be trapped on substitutional sites. Since cw beam annealing of ion implanted Si takes place in solid phase, the formation of supersaturated impurity concentrations seems to be a great deal less likely. However, if the
148
ART0 LIETOILA AND JAMES F. GIBBONS
-DEPTH 4000 2000
(A) 0
1 1 x 1015 A;/cm2/100
0
+
kV
RANDOM SEM ANNEALED UNIMPLANTED SINGLE CRYSTAL
ENERGY (MeV)
FIGURE 3.22. Aligned spectra of 2.2 MeV 4He ions for Si which was implanted with As and then e-beam (SEM) Annealed (from Ref. 3.36).
FIGURE 3.23. TEM micrographs for 75As implanted (100 keV, 1.1~1015 cm-z), SEM annealed Si: (a) bright-field image and (b) weak beam dark-field image (g, 3g, s+ conditions) and
electron diffraction pattern with spot sequence 000 (large spot), 220, 440 and 660 (from Ref. 3.36).
3.
ION IMPLANTED CRYSTALLINE SILICON
149
implant produces an amorphous layer, the circumstances are more favorable. The mechanism of SPE is broadly similar to liquid phase epitaxy : in both processes, there is an interface betweeen two phases which proceeds towards the sample surface.
In Section 3.3, it was shown that the rate of SPE under typical laser annealing conditions is quite high. The dwell time of the laser is typically short (<1 ms) and the cooling rate high, so one might expect that even cw beam processing is capable of forming supersaturated, metastable impurity concentrations. In this section we will show that cw laser annealing of As and P implants (which form an amorphous layer) indeed results in electrically active concentrations exceeding the solid solubilities. Furnace annealing causes those concentrations to relax to equilibrium values: this feature can be used to measure the solubilities of As and P in Si as electrically active dopants. We will also report on measurement of the electron mobilities in both metastable and equilibrium concentrations and on a study of the relaxation process of the metastable As concentrations. Finally, it is shown that cw e-beam annealing is also capable of forming metastable As concentrations. 3.5.2
Cw Laser Annealing of Metastable 75As and 31P Concentrations
Lietoila et al., [3.46-491 have used a cw Ar ion laser to form metastable As and P concentration in Si. In what follows, a review of their experimental procedures and annealing results is given. 3.5.2.1
Sample Preparation, Experimental Methods.
The samples were of 10-20 Q. cm p-type silicon. The implants were carried out at 0°C; the 75As dose was 7.3~1015 cm-2 at 100 keV and the 31P dose 9x1015 cm-2 at 45 keV. Both of these implants produce a maximum as-implanted concentration in excess of 1021 cm-3. The laser annealing was performed using the system described in Chapter I with a 250 mm focusing lens and a sample holder (substrate) temperature of To = 332OC. The beam radius was measured to be 60 um by the technique described in Section 3.3.4.2 above. The laser scanning speed was 7 cm/s, and the distance between adjacent lines was 30 pm. The laser power was 86% of the power required for melting; this results in a calculated (see Chapter 11) maximum temperature of about 12OO0C.
ART0 LIETOILA AND JAMES F. GIBBONS
150 3.5.2.2
Results of Laser Annealing
Laser annealing, as performed under conditions described in Section 3.5.2.1 above, results in practically complete electrical activation of both As and P dopants. This is shown by the sheet Van der Pauw measurements, as compared with the calculated (see Section 3.2.1.1) values o f the sheet resistivity, R,, average given in mobility, !J, and sheet carrier concentration, N,, Table 3.8. The activation is further confirmed by differential Van der Pauw measurements, the results of which are given in Figs. 3.24 (As) and 3.25 (P). Figures 3.24 and 3.25 show that the electron concentrations follow closely the calculated LSSprofiles; and the electron mobilities agree well with the values obtained from the Irvin curve 13.501. We would like to point out that the actual measured quantities are the resistivity and electron concentration versus depth; the mobility is calculated using those two quantities, and the result may not represent the carrier mobility in the conventional sense of the word at these high concentrations. (This point is further discussed in Section 3.5.4 of this chapter). TABLE 3.8. Results from Sheet Van der Pauw Measurements for Laser Annealed Samples. The As implantation dose was 7.3~1015 cm-2 and the energy 100 keV; the P dose was 9x1015 cm-2 at 45 keV. For comparison we give the calculated values for Rs, l~ and Ns for those implants, assuming complete activation and Irvin mobility. (after Ref. 3.46).
Ion
Measurements RS
(n/o) As P
33 25
!J
(cm2/vs) 36 37
Calculations NS
(cm-2)
lJ
NS
(cm2/vs)
(cm-2)
RS (CUD)
5 . 3 ~ 1 0 ~ ~ 33 7.0~1015 28
36
5.3~10~5
33
6.6~1015
The effect of laser power on the electrical activation was studied for the As-implant. Sheet Van der Pauw measurements revealed that the activation was complete even at a power which was 75% of the melting power, corresponding to a maximum annealing temperature of 1O5O0C. Below this, unannealed areas were visible on the sample. Above the power normally used (P/Pmelt = 0.86), slip lines started to form.
3.
ION IMPLANTED CRYSTALLINE SILICON
X
151
HALL MOBILITY
4 I
A5 in Si I00 Lev Z 3 ~ 1 0 ~ ~ c r n ~ ~
lot7
500
1000
2000
1500
DEPTH (A)
FIGURE 3 . 2 4 . Results from differential Van der Pauw measurements for the laser annealed As-implanted sample (fromRef. 3 . 4 6 ) .
X
ELECTRON CONCENTRATION HALL MOBILITY
i X
/Ix
W u
0
P in
lo1'
si 45 kev 9 I 1 0 ~ ~ ~ ~ - ~ 500
1000 DEPTH
1500
(A)
2000
FIGURE 3 . 2 5 . Results from differential Van der Pauw measurements for the laser annealed, P-implanted sample (see Ref. 3 . 4 6 ) .
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ART0 LIETOILA AND JAMES F. GIBBONS
3.5.3
Thermal Treatment of Laser Annealed Concentrations
The active concentrations achieved by laser annealing are well below the values reported by Trumbore for both As and P (3.511. However, the maximum active concentrations achieved by thermal diffusion are about 3 x 1020 cm-3 for As [3.52] and 4 x 1020cm-3 for P [3.53]. Therefore, one could expect thermal annealing to cause deactivation of dopants in the laser annealed samples. Lietoila et al., [3.46-3.49) have shown that deactivation does occur. Long enough thermal annealing makes the active concentration relax to an equilibrium value, which allows the measurement of solubilities of .As and P in Si as electrically active dopants. 'Ihese observations 13.46-3.491 are reported below. 3.5.3.1
Electrical Deactivation
Thermal annealing of the laser annealed, As-implanted samples causes a significant decrease in the sheet carrier concentration, Ns, even at 400°C. This is illustrated in Fig. 3.26 where a plot of Ns vs. annealing temperature for isochronal, 5 minute anneals, is given.
1 ISOCHRONAL ANNEALS FOR 5 min 0
0 0 0
't
AS
in SI
0
0
7.3x t015 cm-'
FIGURE 3.26. The sheet carrier concentration, Ns, in '58simplanted, laser annealed samples as a function of temperature for isochronal (5 min.) anneals (from Ref. 3.46).
Figure 3.26 shows that when the temperature of isochronal annealing is increased, the sheet carrier concentration goes through a minimum, and then starts to increase. The increase is
3.
ION IMPLANTED CRYSTALLINE SILICON
153
due to impurity diffusion, which tends to lower the concentration so that the non-active dopants are activated again. This phenomenon was confirmed by isothermal annealing at different temperatures for both the As- and P- implanted laser annealed samples. The results are given in Figs. 3.27 (As) and 3.28 (P), and show that after a rapid initial drop, the sheet carrier concentrations start to slowly increase. (The exception was the lowest temperature for the P-samples, which gave a monotonic decrease of N,; this is probably due to out-diffusion of the impurities).
A 707OC 0 BOBOC 0 904OC x 1000L
LASER ANNEAL
A s in Si
ISOTHERMAL ANNEALS 100
10'
I
I
I
102
I
,
,
to3
I
s
, ,
to4
ANNEAL TIME (min)
FIGURE 3.27. The sheet carrier concentration, N,, in laser annealed, 75As-implanted samples as a function of annealing time for different temperatures (from Ref. 3.46). Since the initial decrease in the sheet concentration is so much more rapid than the diffusion-assisted reactivation, the behavior of N, vs. time can be divided into two distinct regimes. Lietoila [3.46] proposed the following model, illustrated in Fig. 3.29. While the dopants are deactivated [Fig. 3.29(a)], little diffusion takes place. Eventually the equilibrium active concentration is reached, and only after that, significant diffusion occurs [Fig. 3.29(b)]. During diffusion the maximum active concentration remains unchanged, which results in an increase in the sheet carrier concentration. Testing of this model is described in the following section 3.5.3.2. Metoila et al., [3.47,3.49] laser annealed a second time a sample which was thermally annealed at 900°C for 2 min. The maximum concentration showed partial recovery to 6x1020 ~ m - ~ , suggesting that the solubility at the laser annealing temperature (1200°C) is higher than at 1000°C.
-
154
ART0 LIETOILA AND JAMES F. GIBBONS
I
a 707%
i-1
GS'C 904%
0 0
x 99s-c 5 t 0 0
0
0 0
0
0
0
A ' A A
A 8
P In Si 100
10'
102
10'
104
ANNEAL TIME (min)
FIGURE 3.28. The sheet carrier concentration, N,, in laser annealed, P-implanted samples as a function of annealing time for different temperatures (from Ref. 3 . 4 6 ) .
FIGURE 3 . 2 9 . Schematic presentations of evolution of the active doping concentration. At to, the activation is complete in the laser annealed sample. During thermal annealing, the active concentration reaches thermal equilibrium by the time t2. Diffusion, which is much slower, is significant only for times which are much greater than t2. While diffusion takes place, nonactive dopants are reactivated, and the active concentration remains unchanged (from Ref. 3 . 4 6 ) .
3. 3.5.3.2
ION IMPLANTED CRYSTALLINE SILICON
155
Measurement of Solid Solubilities
The results described above would suggest that the solubility of As and P as electrically active, non-complexed dopants in Si can be determined by letting the laser annealed concentrations relax to thermal equilibrium values. However, to assure that the measurement gives correct results, it has to be determined whether the amount of supersaturation affects the equilibrium concentration. Also, one has to test the assumption presented in the previous section that when diffusion broadens the doping profile the maximum active concentration does not change. Both of those features mean, of course, that the amount of non-active dopant does not affect the maximum active concentration. Lietoila 13.461 studied the effect of the amount of supersaturation by laser and thermal annealing of different As doses. In addition to the dose of 7.3 x 1015 As/cm2, 3 x 1015 and 1 x 1015 As/crnZ were used (both at 100 keV), which give the maximum as-implanted concentrations of 6 x lO2ocm-3 and 2 x 1020 cm-3, respectively. All samples were first laser annealed. The two higher doses (7.3 x lOI5 and 3 x 1015) were thermally annealed for 15 minutes at 904OC. That thermal treatment is, according to Fig. 3.27, enough to pass the minimum in the sheet carrier concentration, which would indicate that the equilibrium concentration is reached. Results of differential Van der Pauw measurements are given in Fig. 3.30, and show that the maximum active concentrations for both doses are identical (2.3 x 1020 cm-3). lhus, the amount of original supersaturation does not have an appreciable effect on the equilibrium concentration. The low dose sample (1 x 1015 As/cm2) was thermally annealed at 904°C for 5 and 10 minutes (the time was short to minimize diffusion effects), with the results of sheet Van der Pauw measurements given in Table 3.9. The maximum as-implanted concentration, 2 x 1020 cm-3, is only slightly below the equilibrium value obtained above, and yet there is no decrease in the sheet concentration during thermal annealing. This further confirms that the maximum active concentration has a unique value which does not depend on the total As-concentration. Lietoila [3.46] tested the assumption that once the minimum in sheet carrier concentration is passed, the equilibrium active concentration is reached and further diffusion does not change that concentration as long as there is inactive As left which can be reactivated. A laser annealed sample with the dose of 7.3 x 1015 As/crnZ was annealed for 2 hours at 904OC. Figure 3.27 shows that this annealing results in a significant increase in sheet carrier concentration. However, differential Van der Pauw measurements (see Fig. 3.30) indicate that the maximum active
ART0 LIETOILA AND JAMES F. GIBBONS
156
1 o2
DOSE
n
A
'
I
TIME
7 3 x 1015 15 mm
0 0
A
73 x 3x
lo1' lo"
2 hr 15 rnm
ANNEALS AT 900OC
-
?
5
- lo3
z
2
z
.
z V w
j!
->t
n
0 0 2
e
loz
-102
5
P
uf U . I
Lu
_I
U I
10'
500
1000
lo1
1500
FIGURE 3.30. Results from differential Van der Pauw measurements of As-implanted, laser annealed samples which were then thermally annealed at 900°C (from Ref. 3.46).
TABLE 3 . 9 Results of Sheet Van der Pauw Measurements After Laser and Thermal Annealing of a Sample Which Was As-implanted to a Dose of 1x1015 cm-2 at 100 keV. This implant gives a maximum concentration of 2x1020 cm-3 (after Ref. 3.46). Annealing
Laser only Laser + 5 min. 900°C Laser + 10 min. 900°C
100 90
89
76 86 86
8.2x10I4 8.2~1014 8.2~1014
concentration is unchanged compared with annealing for the shorter time. These results show that the solubilities of A s and P as electrically active, noncomplexed dopants in Si can be measured by letting the metastable concentrations relax to equilibrium values. Figures 3.27 and 3.28 are used to determine the times required to reach the equilibrium. The results of differential Van der Pauw measurements for samples which were thermally annealed to equilibrium are given in Figs. 3.31 ( A s ) and 3.32 ( P ) .
3.
c
E
ION IMPLANTED CRYSTALLINE SILICON
157
x LASER ONLY o I I OO "C t min1000°C 6min: A 904 O C 16min 1 o 808 a h 707OC76h --
lo2'
V
z
e
-
L
kt
---
1020
w
--
-
0
z
0 0
I
I 019
0
I000 DEPTH
2000
(8)
FIGURE 3.31. Results from differential Van der Pauw measurements for 75As-implanted, laser annealed samples which were thermally annealed to equilibrium (from Ref. 3.48). Figure 3.33 summarizes the results of the solubility measurements. The values for As at 1100°C and 1200°C are subject to some error, however. At llOO°C, the deactivation rate is so high that partial deactivation has probably taken place during the cooling down of the sample. The solubility value at 1200°C was obtained with the second laser scan (see Section 3.5.2.2 above), and it is not certain whether an equilibrium concentration was achieved during that scan. We have included in Fig. 3.33 the solubility data originally reported by Trumbore in 1960 [3.51]. The difference in As solubilities is quite remarkable at about 1000°C. It should be noticed, however, that the actual measurement points in Ref. i3.511 are scattered and exist only at 1200'C and above. The curve quoted in many reference books is obtained by extrapolation of the few high temperature data points. As for the phosphorus solubilities, the disagreement seems to increase with temperature. One possible reason is that phosphorus is deactivated fast enough to cause significant decrease in the active concentration during cooling from even 900'C. This does not seem very likely since Figs. 3.27 and 3.28 show that the rate of deactivation is about the same for As and P. Also, we would like to point out that the P-concentrations measured here are higher than those obtained by thermal diffusion 13.531.
ART0 LIETOILA AND JAMES F. GIBBONS
158
10'
n p TEMPERATURE 998T 904'C 808OC
0
A A 0
0 -
-5
102
L
0
5 -[L t
w z U
z
0 lJ
z
0
EU lo2 w -I w
P in
to'
~i
45 lcev 9 I
500
1015cm-2 I
I
1000 DEPTH
(A)
1500
FIGURE 3 . 3 2 . Results from differential Van der Pauw measurements for P-implanted laser annealed samples which were thermally annealed to equilibrium (from Ref. 3 . 4 6 ) .
3.5.3.3
Discussion
The solubilities of As and P reported above were i 3 . 4 6 - 3 . 4 9 1 inferred from electrical measurements, and represent thus the maximum electrically active concentrations. This information is the most relevant for device processing applications. As was mentioned in the previous section, the solubility values determined by thermal annealing may be too low at the highest temperatures because deactivation may have taken place during cooling of the samples. However, the values given in Fig. 3 . 3 3 are reliable upper limits for active concentrations which in practice can be achieved by any thermal processes. In fact, a production scale process would be expected to yield even lower maximum concentrations than those given here, since cooling is markedly slowed down by the thermal mass of the heavy furnace boats and by the large number of wafers processed simultaneously.
3 . ION IMPLANTED CRYSTALLINE SILICON
--TRUMBORE:
159
0 AS
O P
-THIS WORK : X 0
b
. P I
800
I
I
I
1000
1
I
1200
1400
TEMPERATURE ('C)
FIGURE 3.33. Measured solubilities for As (solid circles) and P (solid squares) in Si as electrically active dopants (Refs. 3.46, 3.48). The point at 1200°C (cross) was obtained by laser scanning a second time a thermally annealed sample (Ref. 3.51). Included also are the original data points given by Trumbore and the extrapolated lines presented in Ref. 3.51 (dashed lines). 3.5.4
Measurement of Hall Mobilities
The laser annealed samples provide a possibility to measure the electron mobilities at concentrations previously unattainable. Furthermore, the effect of nonactive dopants on the mobility can be studied in the concentrations which have relaxed to thermal equilibrium. Lietoila [3.46] has performed this study, with the results reported below. The electron mobilities obtained by the Van der Pauw measurements performed for the laser and thermally annealed, As and P doped samples are given in Fig. 3.34. For comparison, the mobilities calculated from the Irvin resistivity curve [3.50] are included. We would like to point out that the Van der Pauw measurements give the mobility as a function of electron concentration, whereas the resistivity in the Irvin curve is given versus total impurity concentration. Therefore, it is most surprising that the agreement between the two measurements is so good. Moreover, the Van der Pauw measurements show that the mobility depends only on the electron concentration. The amount of nonactive dopant does not have any appreciable effect on the
ART0 LIETOILA AND JAMES F. GIBBONS
160
mobility, as can be seen in Fig. 3.34 by comparing the data obtained from laser annealed and laser and thermally annealed samples.
As METASTABLE mp -,,o As EQUlLlERlUM CONCENTRATION 0 p
--
-
I,
-It-
_I
40I 20
I
I
1
I
2
t019
4
6
1
1
8
I
2
4
6
1020
8
I
I
2
102'
CARRIER CONCENTRATION ( ~ r n - ~ )
FIGURE 3 . 3 4 . Hall mobility of electrons in A s and P implanted Si samples which were either laser annealed (metastable) or laser and thermally annealed (equilibrium) (from Ref. 3 . 4 6 ) .
3.5.5
Characterization of Deactivation of A s
The qualitative kinetics of A s deactivation was reported in Section 3 . 5 . 3 above. In this section, results are given of some more detailed investigations of the deactivation process ( 3 . 4 6 3 . 4 9 1 . The samples for this study are the same as described in Sections 3 . 5 . 2 - 3 . 5 . 4 above. 3.5.5.1
TEM Studies
Results of T E M analysis [3.471 performed on laser and thermally annealed samples are given in Fig. 3 . 3 5 . Laser annealing [Fig. 3 . 3 5 ( a ) ] results, as expected, in complete crystallization of the amorphous material. However, a relatively high concentration of crystal defects is left, mainly in the form of dislocation loops. Thermal annealing at 900°C results in a significant increase in the amount of crystalline damage. This is most likely due to the release of stress present in the supersaturated layer.
3,
ION IMPLANTED CRYSTALLINE SILICON
161
The amount of damage is markedly smaller in samples annealed at 1000°C after laser annealing [Fig. 3.35(d)]. This can be explained by defect annealing at 1000°C. The annealing at 400°C [Fig. 3.35(b)] results in a formation of rod-shaped structures which are about 500 A long and aligned along the <002> direction. These rods might be A s precipitates, even though their volume is t o o small for direct As detection. 3.5.5.2
Channeling Measurements
Results from channeling measurements are given in Fig. 3.36 and Table 3.10. The sample which was only laser annealed shows a relatively high minimum yield of 9.7%: this is due to the remaining crystal damage also observed with TEM. Thermal annealing at 500°C and 900°C increase the minimum yield quite substantially. The result at 900°C is consistent with the TEM observations [see Fig. 3.35(c)]. However, at 400°C-500°C, TEM does not show such a high density of defects, indicating that the defect structures forming at those temperatures are mostly too small to be seen with TEM. Comparison of the fractions of nonsubstitutional and electrically inactive As (see Table 3.10) reveals a surprising fact. At 5OO0C, the off axis A s fraction is smaller than the inactive fraction. But at 900°C the amount of nonsubstitutional As is less than at 5OO0C, while the electrically inactive fraction has substantially increased. The gross difference in those fractions indicates that no precipitation in the usual sense of the word has occurred at 900°C. It is likely that deactivation has taken place via formation of complexes where A s atoms are displaced o f f lattice sites only slightly, so that they cannot be detected with channeling without angular scan. This phenomenon has first been observed by Chu and Masters l3.541 in pulsed laser annealed samples. They used the angular scan to measure the displacement, which was found to be 0.15 A after thermal annealing at 950°C. The angular scan method would not, however, yield any information about the structure of the complexes.
-
The fact that the off-axis A s fraction is larger at 500OC than at 900°C, but still less than the electrical fraction, suggests that deactivation at 500°C may start with the formation of the complexes described above. The complexes would be, however, unstable at the lower temperatures, and would coalesce to form the rods seen by TEM.
162
ART0 LIETOILA AND JAMES F. GIBBONS
FIGURE 3.35. TEM micrographs of the high dose (7.3~1015 cm-2) As implanted samples subjected to laser anneal (a) and laser + thermal annealing (b)-(d) (from Refs. 3.46, 3.47). TABLE 3.10. Comparison ,Between Channeling and Electrical Measurements for As Implanted, Laser and Thermally Annealed Samples. Also given is the minimum yield, xmin, for each sample (after Ref. 3.49). ~~
Annealing
Xmin
Off-axis As ~
Laser only Laser + 5 mins. 500°C Laser + 16 mins. 900°C
9.7% 23.5% 20.2%
7.8% 30.3% 18.9%
-
~~~~
Electrically inactive As -~
5.5% 50.0%
71.0%
3.
ION IMPLANTED CRYSTALLINE SILICON
I
2 . 2 MeV 4He+100,uC ALIGNED
-AS
163
IMPLANTED
LASER ANNEALED ----_LASER+ 16' 900°C -- LASER+ 5' 500°C
-
m
2-
ENERGY [MeV)
FIGURE 3.36. Channeling spectra o f the high dose (7.3~1015 cm-2) 75As implanted samples which were laser and thermally annealed (from Ref. 3.48). 3.5.5.3
Activation Energy of the As Deactivation
The activation energy for the As deactivation process was determined by Lietoila et al. t3.48-3.491 by measuring the time required to reduce the sheet concentration of a laser annealed sample by 7% at various temperatures. For practical reasons, the measurement had to be restricted to a temperature range of 350-410°C; above this range the process is too quick to allow accurate measurements. Results o f the activation energy measurements are given in Fig. 3.37, and yield Ea = 2.0 eV. This value is close to the energy of vacancy formation in Si 13.551 which would support the assumption that As deactivation is a reaction which involves vacancies. 3.5.6 3.5.6.1
E-beam Annealing o f Metastable 75As Concentrations Use of an 2-Beam Welder
The e-beam welder described in Section 3.4.3.1 has been used by Regolini et a1 [3.56] to anneal metastable As concentrations. The samples were of Si implanted to an As dose of 1016/cm-2 at 100 keV. The e-beam annealing was performed at 31 kV using a current of 0.45 ma; the estimated beam diameter was 300 pm.
164
ART0 LIETOILA AND JAMES F. GIBBONS
I / kT l e v - ' )
FIGURE 3.37. The Arrhenius plot of the time required to decrease the sheet concentration by 7% in the A s implanted laser annealed samples (from Ref. 3.48).
Control samples were annealed thermally at 500°C for 30 min. and also using a cw Ar ion laser. Results of the differential Van der Pauw measurements performed on annealed samples [3.56] are given in Fig. 3.38. Both beam annealing methods have resulted in maximum carrier concentrations of about 1.1~1021 cm-3. However, annealing is not complete at the peak of the as-implanted concentration (the calculated LSS As expected on profile is given with a solid line in Fig. 3.38). the basis of the results given in the previous sections, thermal annealing has resulted in poor activation, the maximum electron concentration being only 3x1020 cm-3. None of the three annealing methods had caused any appreciable impurity diffusion. 3.5.6.2
Use of a Modified SEM
A s mentioned in Section 3.4.3.2, a modified SEM was used by Ratnakumar et al. [3.37] to anneal a heavy As implant of 4x1015 cm-2 at 150 keV in Si. Control samples were annealed in a furnace at 900°C and 1000°C for 30 min.
The electron concentrations in the e-beam and thermally annealed samples L3.371 were determined by spreading resistance measurements [3.38], with the results given in Fig. 3.39. E-beam annealing (curve labeled "SEBA") has produced a maximum electron concentration at 7.5~1020 c K 3 , which clearly exceeds the concentrations achievable by thermal annealing I3.521. On the other
3.
ION IMPLANTED CRYSTALLINE SILICON
165
10; 10'' "As'lcmf 100 keV, 0 "C IMPLANT CARRIER CONCENTRATION AN0 MOBILITY C W LASER ANNEALED ANNEALED
OD
r\
0 0 C W. e-BEAM
10"
I
I
600
I
I
1200 DEPTH
(a)
I
I \
,
1500
FIGURE 3.38. Results of differential Van der Pauw measurements for Si samples which wer 75As implanted to a dose of 1016 cm-2 and then e-beam (welder), cw Ar+ laser or thermally annealed (from Ref. 3.56). hand, both thermal anneals have caused significant dopant redistribution which is why the electrons concentrations are below 1020 cm-3. 3.6
HIGH RESOLUTION SELECTIVE E-BEAM ANNEALING OF AS IMPLANTED SILICON
The modified SEM described in Section 3.4.3.2 was used by Ratnakumar et al. [3.37] to demonstrate the potential for a high resolution localized annealing. The samples were Si implanted to a dose of 5x1014 at 100 keV. Results of selective e-beam annealing [3.37] are shown in Fig. 3.40, which is an optical micrograph of an annealed area. The area was scanned twice, with the line and frame scans interchanged for the second scan. The separation between successive scan lines is -2 pm. Due to localized annealing by the finely focussed beam, the two scans produce a grid of annealed areas. In the picture, the light areas correspond to the amorphous material and the dark areas to recrystallized silicon. Annealed islands less than 1 pm2 are shown in the figure. Sheet resistance measurements also indicate
166
ART0 LIETOILA AND JAMES F. GIBBONS
10"
10'1
0
FIGURE 3.39. Doping profiles from spreading resistance measurements for e-beam (SEM) and thermally annealed, 75As implanted Si (from Ref. 3.37).
the very narrow widths of the annealed stripes. Linewidths in the range of 40-100 mm were observed [3.371. Hence scanning e-beam annealing using a regular SEM introduces the intriguing possibility of selective annealing of submicron areas. 3.7
ANNEALING OF B IMPLANTED S I
In contrast to the thorough studies performed on cw beam annealed As-implanted Si, very little work on boron implanted Si has been reported in the literature. One of the earliest papers on cw laser annealing was that of Gat et al. [3.571 reporting the use of a Kr ion laser to anneal boron implanted Si, but it remained for a long time the only one on that subject. Much later, Yep et al. i3.581 used a scanning e-beam for the same task. 'Ihe principal results of those two experiments are reported in the following sections 3.7.1 and 3.7.2. 3.7.1
Use of a CW Kr Ion Laser
Gat et al. i3.571 implanted n-type Si with boron to a dose of 2x1015 cm-2 at 35 keV. %is dose is not sufficient to drive the silicon amorphous but damages the surface layer to a depth of 0.25 tim. Laser annealing was performed using a Kr
3. ION IMPLANTED CRYSTALLINE SILICON
167
ion laser in a multimode operation (the main output is at A = 0.647 pm). The laser power was 6 W and the estimated beam radius was 11 pin (79 mm focussing lens). The scan rate was 9.8 cmlsec and the sample holder was held at 178OC. Control samples were thermally annealed at 1000°C for 30 min. The annealed samples were then analyzed by electrical measurements, SIMS and TEM.
FIGURE 3.40. Micrograph of submicron single crystal islands in amorphous Si formed by annealingwith a finely focused electron beam (SEM) scanning in a grid pattern of 2x2 um spacing (from Ref. 3.37). 3.7.1.1
Electrical Measurements
The electrical sheet resistance measurements i3.571 were performed using a 2 point probe station. The station was originally designed for spreading resistance measurements [3.38], which allows well controlled stepping of the probes across the sample surface. Results of this analysis are given in Fig. 3.41 for one thermally annealed sample (SIB 155) and two laser annealed samples ( 7 S 36 and 7 S 36). The thermally annealed sample shows a probe to probe resistance of 70 52, and very good lateral uniformity. In the laser annealed sample (7 S 37), the probes were stepped from a non-annealed area to the annealed regions in a direction perpendicular to the scan lines. The average probe to probe resistance in the annealed area is 100 52 , with some periodic fluctuation. 'Ihe fluctuation is due to nonuniformity in the temperature produced by the raster scanned beam. The areas under the center of the beam are exposed to higher temperature than the areas between two scans. Stepping of the probes in sample 7 S 36 was performed along the scan lines, and the resistivity is now constant, 120 52.
-
ART0 LIETOILA AND JAMES E GIBBONS
168
....."
0
THERMAL ANNEAL LASER ANNEAL
I~
$200
L
wg 1 0 0 7537
7936
t LASER ANNEALED AREA
FIGURE 3.41. Surface spreading resistance measurements performed on IlB implanted [2xlO15 cm-2, 35 keV] Si subjected to cw Ar+ laser and thermal annealing (from Ref. 3.57).
The fact that the probe to probe resistance in the laser annealed samples is up to 50% higher than in the thermally annealed control sample indicates that the electrical activation of impurities is not quite complete, but about 70-80%. This is further confirmed by the non-uniformity of the resistivity in the laser annealed sample 7 S 37. 3.7.1.2
SIMS Measurements
The boron distributions in the laser and thermally annealed samples were determined by SIMS, and as a calibration, an asimplanted sample was also analyzed in the same run. The results are given in Fig. 3.42 and show that, while thermal annealing has caused very significant motion of boron atoms, laser annealing has completely preserved the as-implanted profile. 3.7.1.3
TEN Analysis
The crystalline quality of laser and thermally annealed samples was studied by TEM, with the results given in Figs. 3.43 and 3 . 4 4 . It should be emphasized that the samples were crystal___ line, but damaged after implantation. The as-implanted damage
3.
ION IMPLANTED CRYSTALLINE SILICON
169
is not, however, resolvable by TEM. 'Ihe thermally annealed samples (Fig. 3.43) contains large defects such as loops and dipoles, which are formed during annealing by coalescence of the small defects created by the implant.
-8
+-
0 0
z lo2-
+
m
* t*
0
k c
*++
0
e
>. Y, Z W
+*+
b
++
0
++ ++
e 0
% e
+
' 0
D. 0
FIGURE 3.42. Relative boron profiles determined by SIMS in as-implanted, laser annealed and thermally annealed, IlB implanted [2xlO15, 35 keV] Si (from Ref. 3.57). The laser annealed sample, Fig. 3.44 shows two distinct regions. The center of the scan (labelled "Laser scanned region") is completely free of defects. There is no trace of rod-shaped structures or dislocation loops, typically found in boron implanted samples thermally annealed at temperatures > 700°C. However, at the edge of a scan line (labelled '' unscanned region" in Fig. 3.44), we see small dislocation loops whose concentration increases with increasing distance from the center of the scan. This means that the laser induced temperature in the edge region was such as to allow some growth of the defects during the scan, but was too low to produce a defect-free crystalline structure. 3.7.2
Use of a Scanning Low Resolution E-Beam
Yep et-al. [3.58] have used a broadly rastered, low resolution electron beam to anneal 1lB-implanted diode structures. The
170
ART0 LIETOILA AND JAMES E GIBBONS
FIGURE 3.43. Transmission electron micrograph of a Bimplanted sample which was thermally annealed at 1000°C for 30 min. (from Ref. 3.57). device properties are beyond the scope of this chapter, but since the experimental apparatus of Yep et al. incorporates some interesting features, we report on their experiments. Diodes were fabricated on (100) n-t e Si wafers (2" diameat 50 keV. The ter) by implanting B to a dose of 1xlOl' diode areas were defined to be 16 mil (0.41 mm) diameter dots by a 0.63 pm thick Si02 layer. The electron beam annealing was performed using an accelerating voltage of 5.2 keV and a current of 52 mA. The annealing system is illustrated in Fig. 3.45. The electron beam was deflected electrostatically k1.5" in the x-direction and 0.75" in the y-direction. In addition, the sample stage was mechanically scanned at 0.13 inchis in the y-direction. This arrangement was effectively a 1.5"~ 3" bar of electrons annealing the wafer. The annealing temperature was measured t o be 850 +2OoC by an optical pyrometer, and the exposure time for a given point was 3-4 s. Annealing was completed by a single pass across the whole wafer, which only took about 20 s. Control wafers were annealed at 850OC and 1000°C for 30 min in flowing nitrogen. The B doping profiles in the annealed samples were determined by SIMS, with the results given in Fig. 3.46. While both furnace
3.
ION IMPLANTED CRYSTALLINE SILICON
171
FIGURE 3 . 4 4 . Transmission electron micrograph of B-implanted, laser annealed Si. The area labelled "laser scanned" is at the center of the scan, and the "unscanned region" refers to the edge of a single scan (from Ref. 3 . 5 7 ) .
X A N D Y SCANNED 5.2 keV E . B E A M
SPEED 0.13 IN/SEC
\\-
FIGURE 3 . 4 5 . Experimental setup f o r the low resolution ebeam annealing system (from Ref. 3.58). 9 9 8 1 Elsevier Science Publishing CO.
ART0 LIETOILA AND JAMES F. GIBBONS
172
anneals have broadened t h e B p r o f i l e by 1200-1600 8 , t h e p r o f i l e has s h i f t e d only about 450 8 d u r i n g e-beam annealing. Figure 3.46 a l s o shows t h e e l e c t r i c a l l y a c t i v e B p r o f i l e , as determined by s p r e a d i n g r e s i s t a n c e measurements. The measured area under t h i s curve shows t h a t t h e a c t i v a t i o n of boron by e-beam a n n e a l i n g i s about 7 5 % .
t FIGURE 3.46. SIMS l l B p r o f i l e f o r t h e sample as-implanted, a f t e r e-beam a n n e a l and a f t e r f u r a n c e a n n e a l a t 1000°C f o r 30 min. For comparison, the e l e c t r i c a l l y a c t i v e l l B p r o f i l e obtained by s p r e a d i n g r e s i s t a n c e a f t e r e-beam a n n e a l i s shown (Ref. 3 . 5 8 ) . 01981 Elsevier Science Publishing Co.
I n conclusion,
the
e-beam
a n n e a l i n g method
by Yep e t a l .
(3.581 d i f f e r s s i g n i f i c a n t l y from those p r e s e n t e d i n S e c t i o n s 3 . 4 . 3 , 3.5.6 and 3.6. The s c a n frequency w a s 550 Hz i n t h e xd i r e c t i o n and 330 Hz i n the y - d i r e c t i o n . While t h e l a t t e r systems f e a t u r e a narrow, low power beam, l e a d i n g t o s h o r t dwell t i m e s but long p r o c e s s i n g t i m e f o r a wafer, Yep e t a l . u t i l i z e a broad beam w i t h a long exposure t i m e ( 4 s ) and f a s t wafer proc e s s i n g . The long dwell t i m e permits the temperature t o be Even though t h e r e s u l t s of Yep e t r e l a t i v e l y low, about 850°C. a l . [ 3 . 5 8 ] a r e obviously n o t y e t o p t i m a l ( t h e e l e c t r i c a l a c t i v a 7 5 x 1 , t h e i r experiment demonstrates t h e p o t e n t i a l t i o n was only of a wide beam t o combine d i f f u s i o n f r e e a n n e a l i n g and a high throughput c a p a c i t y .
-
-
3
ION IMPLANTED CRYSTALLINE SILICON
173
REFERENCES 3.1 3.2 3.3 3.4
3.5 3.6 3.7 3.8 3.9 3.10
Gat, A., Gibbons, J. F., Magee, T. J., Peng, J., Deline, V. R., Williams, P., and Evans, C. A., Jr., Appl. Phys. Lett. 32, 276 (1978). Williams, J.S., Brown, W. L., Leamy, H. J., Poate, J. M., Rodgers, J. W., Rousseau, D., Rozgonyi, G. A., Shelnutt, J. A., and Sheng, T. T., Appl. Phys. Lett. 33, 542 (1978). Gat, A., Lietoila, A., and Gibbons, J. F., J. Appl. Phys. 50, 2926 (1979). Olson, G. L., kkorowski, S. A., Roth, J. A., and Hess, L. D., in "Laser and Electron-Beam Solid Interactions and Materials Processing" (J. F. Gibbons, L. D. Hess and T. W. Sigmon, eds.), p. 125. North-Holland (1981). Gat, A., Gibbons, J. F., Magee, T. J., Peng, J., Williams, P., Deline, V., and Evans, C. A., Jr., Appl. Phys. Lett. 33, 389 (1978). Lietoila, A . , Gibbons, J. F. , Magee, T. J. , Peng, J., and Hong, J. D., Appl. Phys. Lett. 35, 532 (1979). Lietoila, A . , Gibbons, J. F., and Sigmon, T. W., Appl. Phys. Lett. 36, 765 (1980). Regolini, J. L., Sigmon, T. W., and Gibbons, J. F., Appl. Phys. Lett. 35, 114 (1979). See, e.g., White, C. W., Christie, W. H., Appleton, B. R., Wilson, S. R., and Pronko, P. P., Appl. Phys. Lett. 33, 662 (1978). Thomas, G., "Transmission Electron Microscopy of Materials , Wiley (1971). Van der Pauw, L. J., Philips Technical Review 20, 220 (1958/1959). Johansson, N. G. E., Mayer, J. W., and Marsh, 0. J., Solid State Electron. 13, 317 (1970). Bube, R.H., "Electronic Properties of Crystalline Solids," p. 379, Academic Press (1974). Barber, H. D., Lo, H. B., and Jones, J.E., J. Electrochem. SOC. 123, 1404 (1976). Regolini, J. L., unpublished. Dearhaley, G., Freeman, J. H . , Nelson, R. S., and Stephen, J., "Ion Implantation," p. 539, North-Holland (1973). Hofker, W.K., Werner, H.W., Oosthoek, D.P., and de Grefte, H. A. M., Appl. Phys. 2, 265 (1973). Williams, P., Lewis, R. K., Evans, C. A., Jr., and Harley, P. R., Analytical Chem. 49, 1399 (1977). Williams, P., and Evans, C. A., Jr., Appl. Phys. Lett. 30, 559 (1977). Dobrott, R. D., NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices, held at NBS, Gaithersburg, Md., 23-24 April 1975, p. 31. Ibid, p. 45 Chu, W.-K., Mayer, J.W., and Nicolet, M.-A,, "Backscattering Spectrometry," Academic Press (1978). "
3.11 3.12 3.13 3.14 3.15 3.16. 3.17 3.18 3.19 3.20 3.21 3.22
174 3.23 3.24 3.25 3.26 3.27 3.28
3.29 3.30 3.31 3.32 3.33 3.34
3.35
3.36
3.37
3.38 3.39 3.40
3.41
A R T 0 LIETOILA AND JAMES F. GIBBONS
Nuclear I n s t r u m e n t s and Methods 126, 205 Williams, J . S . , (1975). Chu, W.-K., Muller, H., and Mayer, J. W., Sigmon, T. W., J. Appl. Phys. 5, 347 (1975). C s e p r e g i , L., Mayer, J. W., and Sigmon, T. W., Phys. L e t t . 54A, 157 (1975). German, P., Squelard, S . , Bourgoin, J. C., Zellama, K., and Thomas, P. A,, J. Appl. Phys. 50, 6995 (1979). Gat, A., Ph.D. T h e s i s , S t a n f o r d U n i v e r s i t y , Department of E l e c t r i c a l Engineering (1979). Roth, J. A., Anderson, C. L., and Dunlap, Hess, L. D., H. L., "Laser-Solid I n t e r a c t i o n s and Laser P r o c e s s i n g 1978," ( S . D. F e r r i s , H. J. Leamy, and J. H. P o a t e , eds.) A I P Conf. Proc. 50, 496, American I n s t i t u t e of P h y s i c s (1979). Olson, G. L., Kokorowski, S. A., McFarlane, R. A., and Hess, L. D., Appl. Phys. L e t t . 37, 1019 (1980). Kennedy, E. F., G a l l a g h e r , T. J., Mayer, Csepregi, L., J . W., and Sigmon, T. W., J. Appl. Phys. 48, 4234 (1977). Kennedy, E. F., Mayer, J. W., and Sigmon, Csepregi, L., T. W., 3. Appl. Phys. 49, 3906 (1978). L i e t o i l a , A., Gold, R. B., and Gibbons, J . F., Appl. Phys. L e t t . 53, 1169 (1982). Ph.D. T h e s i s , S t a n f o r d U n i v e r s i t y , DepartL i e t o i l a , A., ment of Applied Physics (19811, Chapter IV. Auston, D. H., Golovchenko, J. A . , Smith, P. R., Surko, C. M., and Venkatesen, T. N. C., Appl. Phys. Lett. 33, 539 (1978). Gibbons, J. F., Sigmon, T. W., Pease, R e g o l i n i , J. L., R. F. W., Magee, T. J., and Peng, J., Appl. Phys. Lett. 34, 410 (1979). Johnson, N. M., S i n c l a i r , R., Sigmon, R e g o l i n i , J. L., T. W., and Gibbons, J. F., i n "Laser and Electron-Beam P r o c e s s i n g of Materials," (C. W. White and P. S . Peercy, eds.), p. 297, Academic P r e s s (1980). Pease, R. F. W., B a r t e l i n k , D. J., Ratnakumar, K. N., Johnson, N. M., and Meindl, J.D., Appl. Phys. L e t t . 35, 463 (1979). and Dutton, R. W., S t a n f o r d D'Avanzo, D. C., Rung, R. D., 5013-2 Electronics Laboratories T e c h n i c a l Report No. (1977). Narayan, J., and Young, R. T., AIP ConferWhite, C. W., ence Proceedings 50, 275 (1979). Grob, J. J., Grob, A., S i f f e r t , Fogarassy, E., Stuck, R., P., i n "Laser and Electron-Beam P r o c e s s i n g of Materials," (C. W. White and P. S. Peercy, eds.), p. 117, Academic P r e s s (1980). i n "Laser and E l e c t r o n B e a m P r o c e s s i n g of Rimini, E., E l e c t r o n i c Materials," ECS Proc. 80, 270, E l e c t r o c h e m i c a l S o c i e t y (1980).
3 - ION IMPLANTED CRYSTALLINE SILICON 3.42 3.43 3.44
3.45 3.46 3.47 3.48 3.49 3.50 3.51 3.52 3.53 3.54
3.55 3.56 3.57 3.58
3.59 3.60 3.61
175
Tamura, M., Natsuaki, N., and Tokuyama, T., in "Laser and Electron-Beam Processing of Materials," (C. W. White and p. 247, Academic Press (1980). P. S. Peercy, eds.), Chu, W. K., Mader, S . R., and Rimini, E., in "Laser and Electron-Beam Processing of Materials," (C. W. White and P. S . Peercy, eds.) p. 253, Academic Press (1980). Pianetta, P., Amano, J., Woolhouse, G., and Stolte, C. A., "Laser and Electron-Beam Solid Interactions and Materials Processing," (J. F. Gibbons, L. D. Hess and T. W. Sigmon, eds.), p. 239, North Holland (1981). Jackson, K. A., Gilmer, G. H., and Leamy, H. J., in "Laser and Electron-Beam Processing of Materials," (C. W. White and P. S . Peercy, eds.), p. 104, Academic Press (1980). Lietoila, A,, Ph.D. Thesis, Stanford University, Department of Applied Physics (1981). Chapter V. Lietoila, A., Gibbons, J. F., Magee, T. J., Peng, J., and Hong, J . D., Appl. Phys. Lett. 35, 532 (1979). Lietoila, A,, Gibbons, J. F., and Sigmon, T. W., Appl. Phys. Lett. 36, 765 (1980). Lietoila, A., Gibbons, J. F., Sigmon, T. W., Magee, T. J., Peng, and Hong, J . D., [Ref. 3.411, p. 350. Sze, S . M., "Physics of Semiconductor Devices," p. 43, Wiley (1969). Trumbore, F. A., Bell Syst. Tech. J. 39, 205 (1960). Tsai, M.Y., Morehead, F.F., Baglin, J.E.E., and Michael, A. E., J. Appl. Phys. 51, 3230 (1980). Fair, R. B., J. Electrochem. SOC. 125, 323 (1978). Chu, W. K., and Masters, B. J., "Laser-Solid Interactions and Laser Processing - 1978," ( S . D. Ferris, H. J. Leamy, and J. H. Poate, eds.) AIP Conf. Proc. 50, 305, American Institute of Physics (1979). See, e.g., Bennerman, K. H., Phys. Rev. 137A, 1497 (1965); and Van Vechten, J. A., and Turmond, C.D., Phys. Rev. B14, 3551 (1976). Regolini, J. L., Sigmon, T. W., and Gibbons, J. F., Appl. Phys. Lett. 35, 114 (1979). Gat, A., Gibbons, J. F., Magee, T. J., Peng, J., Williams, P., Deline, V., and Evans, C. A,, Jr., Appl. Phys. Lett. 33, 389 (1978). Yep, T. O., Fulks, R. T., and Powell, R. A., in "Laser and Electron-Beam Solid Interactions and Materials Processing" (J. F. Gibbons, L. D. Hess and T. W. Sigmon, eds.) p. 345, North-Holland ( 1981) Lietoila, A., Wakita, A., Sigmon, T. W., and Gibbons, J.F. J. Appl. Phys. 53, 4399 (1982). Lidow, A., Ph.D. Thesis, Stanford University, Dept. of Applied Physics (1977). Nissim, Y. I., Ph.D. Thesis, Stanford University, Dept. of Electrical Engineering (1981).
.
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CHAPTER 4
Electronic Defects in CW Transient Thermal Processed Silicon N . M . Johnson XEROX CORPORATION PALO ALTO RESEARCH CENTER PALO ALTO, CALIFORNIA
4.1 4.2 4.3 4.4 4.5 4.6 4.7 4.8
INTRODUCTION . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . DEVICE PROCESSING AND MEASUREMENT TECHNIQUES ....................................... CW BEAMPROCESSED BULKSILICON. . . . . . . . . . . . . . . . . LATERAL NONUNIFORMITIES IN SCANNED-BEAM ANNEALING. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . ISOTHERMAL TRANSIENT ANNEALING. . . . . . . . . . . . . . . . . . POST-RECRYSTALLIZATION PROCESSING. . . . . . . . . . . . . . . . BEAM-CRYSTALLIZED SILICONTHINFILMS. . . . . . . . . . . . . . SUMMARY AND CONCLUSIONS
........................
REFERENCES. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.1
177 179 184 195 200 204 21 1 220 223
INTRODUCTION
Directed energy sources such as lasers and electron beams and incoherent light sources have been proposed as alternatives to conventional furnace annealing for recrystallizing the amorphous layer created by high-dose ion implantation in singlecrystal silicon. These energy sources have also been used to crystallize silicon thin films on insulating amorphous substrates in order to obtain semiconducting material for electronic device fabrication. In both applications the purpose of transient thermal processing is to produce single-crystal material of high crystalline perfection. With cw energy sources, recrystallization of implanted amorphous layers occurs by solid-phase epitaxy, with lower densities of extended defects (e.g., stacking faults) than can be achieved by conventional furnace annealing. However, materials studies with techniques such as capacitance transient spectroscopy and luminescence reveal high densities of residual defects in and near the recrystallized layers, which give rise to deep levels in the silicon forbidden-energy band. This chapter reviews the current state of knowledge in the ongoing study of residual defects in recrystallized bulk silicon and introduces the subject of electronic defect evaluation in crystallizedsilicon thin films. The high degree of perfection and control which has been achieved in silicon integrated-circuit fabrication must be retained while realizing the potential advantages of beam recrysSEMICONDUCTORS AND SEMIMETALS, VOL. 17
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Copyright 0 1984 by Academic Press, Inc. All rights of reproduction in any form reserved. ISBN 0-12-752117-8
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tallization. Only then will the unique features of spatial and temporal localization of energy be appreciated. The stringency of this requirement becomes evident when it is noted that one of the major achievements underlying the success of present-day silicon integrated-circuit processing is the minimization and control of electrically active defects in single-crystal silicon. Impurities such as oxygen and carbon can be tolerated at concentrations of the order of 10l8 ~ m - ~which , is typical for Czochralski-grown silicon, because they can be rendered electrically inactive. On the other hand, a metallic impurity such as gold can significantly alter the electrical pro erties of a device at concentrations of less than loll cm-' through the introduction of deep traps and recombination centers. For transient thermal processing of silicon, the issue centers on the removal of electronic defects associated with lattice damage, such as that created by ion implantation. This chapter first examines the problem of residual defects which result from the use of cw transient thermal sources to recrystallize the amorphous layer created by high-dose ionimplantation in silicon. This topic is of central importance in silicon integrated-circuit processing due to the extensive use of ion implantation to introduce dopant impurities (e.g., As, P, and B) for spatially controlling the free-carrier concentration. The displacement damage created by ion implantation is conventionally removed by furnace annealing, which also electrically activates the implanted impurities through their placement on substitutional lattice sites. As discussed in Chapter 3, directed energy sources such as laser and electron beams can be used for this purpose. The potential advantages of energy beams are derived from their spatial and temporal features. For example, regions of silicon can be selectively annealed thereby preventing possible electrical degradation of other regions not requiring thermal cycling. It has been shown that implanted layers can be recrystallized with low densities of extended defects and with complete electrical activation of the dopant. Cw transient thermal sources, including incoherent light sources, offer the additional feature that the implanted amorphous layer is recrystallized without dopant redistribution. This is a consequence of recrystallizing the amorphous layer by solid-phase epitaxial regrowth on a time scale of seconds or less, which precludes significant dopant diffusion. In addition, it is generally found that free-carrier mobilities in recrystallized layers equal their values in conventionally processed bulk single-crystal silicon. The above features of cw transient thermal processing are particularly applicable in silicon submicron device fabrication and have motivated numerous studies of residual electronic defects and their dependence on transient annealing conditions. Beam-crystallized silicon thin films are of technological interest as a materials/processing system for fabricating highperformance thin-film transistors (TFT) and circuits. As discus-
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sed in Chapter 5 , such devices have been fabricated on insulating layers over single-crystal silicon wafers and could lead to vertical integration in silicon integrated-circuit technology. High-performance TFTs have also been fabricated on bulk glass substrates and may be applicable for switching and logic circuitry in large-area displays. For a TFT technology, crystalline silicon thin films offer the significant feature of utilizing conventional silicon microelectronic processing techniques in the fabrication of thin-film devices and circuits. The potential impact of these rapidly emerging device technologies underscores the need for comprehensive evaluation of the electronic properties of crystallized-silicon thin films, and in particular, characterization of electronic defects and their correlation with materials processing and device operation. The chapter is organized as follows: general experimental considerations of device fabrication and measurements of elecElectronic defects tronic defects are discussed in Section 4.2. in cw beam annealed bulk silicon are reviewed in Section 4 . 3 . This section is concerned both with residual defects arising from the incomplete removal of ion-implantation damage by beam annealing and with the generation of defects under such annealing conditions. Laterally distributed defects are examined in Section 4.4, which includes consideration of the effect of patterned dielectric overlayers in scanned beam annealing. Section 4.5 presents results from deep level studies in isothermal transient annealed silicon. Section 4.6 reviews post-recrystallization processing techniques for further removal of residual defects. The emerging topic of beam-crystallized silicon thin films is discussed in Section 4.7 with results for electronic defects in silicon layers on both thin film and bulk amorphous substrates. A summary with conclusions is presented in Section 4.8. 4.2
DEVICE PROCESSING AND MEASUREMENT TECHNIQUES
In bulk single-crystal silicon, several forms of energy beams have been used to recrystallize the amorphous layer created by high-dose ion implantation. The features of cw beam processing include spatial selection of annealed regions, negligible dopant redistribution, and retention of majority-carrier transport properties. As discussed in Chapter 1, these features have been demonstrated with scanned cw lasers, scanned electron beams, incoherent light sources, and line-source (shaped) electron beams. However, from electrical measurements that are specifically sensitive to minority-carrier transport and from direct measurements of deep levels it has been found that all of the above forms of beam annealing leave residual electronic defects, with densities in excess of those obtained by conventional furnace annealing. This section reviews general considerations in the design and fabrication of test devices for the characterization of electronic defects in transient thermal annealed silicon.
N. M.JOHNSON
4.2.1
CW Transient Thermal Processing
The principal diagnostic technique that has been used for electronic defect evaluation is capacitance transient spectroscopy performed on current-rectifying devices. Both Schottkybarrier and p-n junction diodes have been used €or this purpose. Schematic cross-sectional diagrams of these structures are shown in Fig. 4.1, with the amorphized layer resulting from high-dose ion implantation represented by the cross-hatch patterns. The Schottky diode (Fig. 4.la) is the simplest device to fabricate, requiring only the deposition of a metallic rectifying contact onto the silicon surface. However, the Schottky contact cannot be used to directly evaluate implanted dopants because the high doses required to amorphize silicon (>1014 produce a degenerately doped surface layer which destroys the current rectifying feature of the Schottky barrier. To utilize this structure, self implantation may be used to create lattice displacement damage, of the form accompanying implantation of the commonly used dopants in silicon, without altering the shallow dopant concentration. Alternatively, the Schottky diode is well suited for studying induced defects in the absence of implantation. Dopant implants are conveniently studied with p-n junction test devices in which the heavily doped surface layer permits the fabrication of structures which closely approximate one-sided step junctions. The p-n junction structure can be realized i n two configurations: the mesa diode (Fig. 4.lb) and the oxidepassivated planar diode (Fig. 4.1~). As test devices, these two configurations differ solely in the form of the junction perimeter. And while the planar p-n junction is the more technologically relevant configuration, the mesa structure is often the more convenient to fabricate for defect studies. For capacitance transient spectroscopy it is essential to have back electrical contacts which remain ohmic at low temperatures, since the measurement is performed over a range of temperatures typically from 78 K to above room temperature. This i s readily achieved through the use of semiconducting epilayers on degenerately doped substrates. On bulk single-crystal silicon the back surface may be implantedwith the substrate dopant, which is then electrically activated with a furnace anneal (e.g., 800°C, 30 min, N2), in order to produce a degenerately doped layer for ohmic contact. To further facilitate electrical connection to the back contact, a metallic thin film (e.g., aluminum) can be deposited over the back surface of the wafer and then sintered In studies of scanned(at, e.g., 450°C, 30 min, forming gas). beam annealing, preparation of back contacts can be completed prior to front surface processing. Transient thermal processing is used to recrystallize the amorphous layer created by ion implantation and to electrically
4. ELECTRONIC DEFECTS IN SILICON
n -type Si
METAL
RECRYSTALLIZED LAYER
-’ p-type Si
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\-RECRYSTALLIZED LAYER
I
n-type si METALLURGICAL JUNCTION
METAL
LL
RECRYSTALLIZED J’ LAYER
n-type Si METALLURGICAL JUNCTION
p-type Si
FIGURE 4.1. Schematic diagrams of two-terminal current rectifying devices for electronic defect evaluation in recrystallized bulk silicon: (a) Schottky-barrier diode, (b) p-n junction (mesa) diode, and (c) oxide-passivated planar p-n junction diode. The cross-hatch areas represent material driven amorphous by ion implantation and subsequently recrystallized. activate implanted dopants. Cw beam processing is performed with a focused laser or electron beam by scanning the energy beam over the silicon surface in a raster pattern. Energy beams shaped to form a scanned line source permit recrystallization over large areas in a single pass. In scanned beam processing, the silicon wafer is locally heated with dwell times of the deposited energy typically of the order of milliseconds. With isothermal transient annealing (Section 4 . 5 ) , the entirewafer is annealed at an essentially constant temperature €or times on the order of seconds.
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T e s t d e v i c e s are completed by vacuum e v a p o r a t i n g m e t a l l i c t h i n f i l m s o v e r t h e f r o n t s u r f a c e of t h e wafer and u s i n g conFor p-n ventional photolithography t o define the electrodes j u n c t i o n d i o d e s , plasma e t c h i n g o r w e t chemical e t c h i n g i s used t o form t h e mesa s t r u c t u r e , and on p l a n a r diodes t h e j u n c t i o n i s defined by chemically e t c h e d windows i n t h e s i l i c o n - d i o x i d e passivation layer. For c o n t r o l d e v i c e s , conventional f u r n a c e a n n e a l i n g r e p l a c e s t r a n s i e n t thermal p r o c e s s i n g i n t h e f a b r i c a t i o n of Schottky o r p-n j u n c t i o n test d e v i c e s .
.
4.2.2
Measurement Techniques
S e v e r a l e l e c t r i c a l and p h y s i c a l measurement t e c h n i q u e s have been used t o e v a l u a t e r e s i d u a l e l e c t r o n i c d e f e c t s , and t h e i r possible c o r r e l a t i o n with s t r u c t u r a l imperfections, i n t r a n s i e n t thermal r e c r y s t a l l i z e d s i l i c o n . Both c a p a c i t a n c e t r a n s i e n t spectroscopy and luminescence have been used t o d i r e c t l y probe deep l e v e l s i n t h e s i l i c o n bandgap. However, i t s high s e n s i t i v i t y and d i r e c t e l e c t r i c a l d e t e c t i o n of r e s i d u a l deep l e v e l s have rendered c a p a c i t a n c e - t r a n s i e n t spectroscopy t h e most widely used diagnostic technique f o r electronic-defect e v a l u a t i o n . I n d i r e c t o b s e r v a t i o n s of r e s i d u a l d e f e c t s , by d e t e c t i o n of m i n o r i t y - c a r r i e r recombination, have been made w i t h current/photocurrent-voltage measurements and w i t h electron-beam-induced c u r r e n t s . S t r u c t u r a l information f o r c o r r e l a t i o n w i t h e l e c t r i c a l measurements has been obtained from Nomarski-interference microscopy, t r a n s m i s s i o n e l e c t r o n microscopy, and x-ray topography. Most of t h e r e s u l t s reviewed i n t h i s c h a p t e r were o b t a i n e d by c a p a c i t a n c e t r a n s i e n t s p e c t r o s c o p y . The b a s i c method is termed deep-level t r a n s i e n t spectroscopy (DLTS) [ 4 .l] and y i e l d s t h e energy l e v e l s and d e n s i t i e s of e l e c t r o n i c d e f e c t s . Many of t h e r e s u l t s p r e s e n t e d h e r e were o b t a i n e d from DLTS measurements performed i n t h e constant-capacitance mode, which i s p a r t i c u l a r l y a p p l i c a b l e when measuring t r a p d e n s i t i e s comparable t o t h e shallow dopant c o n c e n t r a t i o n . A d e t a i l e d a n a l y s i s of constant-capacitance DLTS (CC-DLTS) measurements of bulk-semiconductor d e f e c t s h a s I n a d d i t i o n , v a r i o u s DLTS techbeen p r e s e n t e d elsewhere [ 4 . 2 ] niques have been developed t o o b t a i n s p a t i a l depth p r o f i l e s of d e f e c t l e v e l s [4.1-4.31
.
.
The e s s e n t i a l f e a t u r e s of t h e DLTS technique are i l l u s t r a t e d i n F i g . 4.2 w i t h an energy-band diagram f o r a S c h o t t k y - b a r r i e r s t r u c t u r e w i t h a s i n g l e t r a p l e v e l a t energy Et i n t h e bandgap of an n-type semiconductor. The DLTS measurement i n v o l v e s f i r s t t h e a p p l i c a t i o n of a r e v e r s e b i a s VR t o e s t a b l i s h a d e p l e t i o n l a y e r of width WD. The t r a p l e v e l i n t e r s e c t s t h e quasi-Fermi energy EFS a t XD. Traps l o c a t e d below EFS are f i l l e d w i t h elect r o n s and t h o s e above a r e empty. Voltage p u l s e s p e r i o d i c a l l y
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183
Semiconductor
FIGURE 4.2. Schematic energy-band diagram for a Schottkybarrier structure with a single discrete defect level at an energy Et in the silicon bandgap. The distance WD is the depletion width for a reverse bias VR, and XD is the depth below the silicon surface at which the defect level intersects the quasi-Fermi level. The distance X is the depth at which the trap level intersects the quasi-Fermf energy during a pulse bias Vp. reduce the depletion width in order to populate additional traps. After returning to the quiescent bias, these trapped electrons are thermally emitted to the conduction band and swept out of the depletion layer. This gives rise t o a capacitance transient in the transient-capacitance mode or a voltage transient in the constant-capacitance mode of measurement. Analysis of the DLTS signal provides a characterization of deep levels in the semiconductor. The DLTS signal is obtained by forming the difference of the transient response (either capacitance or voltage transient) measured at two delay times ti and t2 after a charging pulse. For an exponential transient, the delay times define an emission rate constant eo which is given by the expression [4.1]
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A DLTS spectrum is obtained by recording the DLTS signal over a range of temperatures. An emission peak appears at that temperature for which the emission rate of a given trap, en, equals eo. The emission rate constant can be varied to obtain the emission rate over a range of temperatures. Then the activation energy for thermal emission is obtained from an Arrhenius analysis of the emission rate, which from considerations of detailed balance may be expressed as
where un is the electron capture cross section, vn is the mean thermal velocity for electrons, Nc is the effective density of states in the semiconductor conduction band, Ec is the conductionband minimum, k is Boltzmann's constant, and T is the absolute temperature. The DLTS analysis thus yields the activation energy for thermal emission of charge carriers and the cross-section for capturing free carriers. In addition, the amplitude of the DLTS signal is proportional to the trap concentration and can be used to measure the spatial depth profile of deep levels.
4.3
4.3.1
CW BEAM PROCESSED BULK SILICON Recrystallized Silicon
Electronic defects remaining after beam recrystallization of a n implanted amorphous layer may be considered to arise from either of two origins. Ion implantation introduces displacement damage to depths beyond the amorphized surface layer. A beam annealing schedule can be sufficient to recrystallize the amorphous layer by solid phase epitaxy while only partially removing this more deeply penetrating lattice damage. On the other hand, beam irradiation can generate both electronic and extended structural defects as a consequence of the high thermal gradients and rapid quenching which are consequences of the spatial and temporal features of beam recrystallization. A clear distinction between these two origins can be obscured by defect interactions; that is, the form, distribution, and density of induced defects may be altered by the presence of displacement damage and impurities. Residual electronic defects in cw beam-recrystallized silicon were first evaluated in Schottky-barrier diodes which were processed with a scanned cw Ar-ion laser i4.41. Epitaxial silicon
4. ELECTRONIC DEFECTS IN SILICON
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wafers with n-type conductivity were implanted at room temperature with SiH' to create displacement damage, of the form accompanying the implantatlon of dopants, without altering the shallow dopant concentration. The ions were implanted at 80 keV to a dose of 2x1015 cm'2, which is sufficient to drive the silicon amorphous to a depth of -120 nm. The im lanted species was chosen to avoid nitrogen contamination (N24) which can occur during 28Si+ implantation. CC-DLTS emission spectra are presented in Fig. 4 . 3 for an as-laser-annealed diode. The two emission spectra were recorded with different combinations of reverse and pulse biases, which define different spatial observation windows for trap detection. The bounds of the spatial intervals quoted in Fig. 4 . 3 (i.e., AWD) are the steady-state depletion depths under reverse and pulse biases, which were computed from the measured device I
2a
-
I
I
Epitaxial Silicon N-Type, (100) Self Implanted en = 347 sec-'
I
I
(- 0.49 eV)
->-E 15 ID
5
iij v)
5
9 0
10
0
5
0 -1
I t 100
I 150
I
I
I
200
250
300
0
T(K) FIGURE 4 . 3 . CC-DLTS spectra for electron emission in selfimplanted cw laser-annealed silicon and from a furnace-annealed control device.
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capacitance. The i n t e r v a l s i n d i c a t e t h e r e l a t i v e d e p t h s of t h e The e l e c t r o n e m i s s i o n o b s e r v a t i o n windows f o r t h e two s p e c t r a . s p e c t r a are dominated by two d e f e c t l e v e l s which are l o c a t e d n e a r t h e middle of t h e s i l i c o n band gap. The e m i s s i o n s p e c t r a recorded w i t h d i f f e r e n t o b s e r v a t i o n windows i n d i c a t e t h a t t h e r a t i o of t h e d e n s i t i e s of t h e two d e f e c t l e v e l s v a r i e s r a p i d l y w i t h depth. The a c t i v a t i o n e n e r g i e s were o b t a i n e d from measurements over s p a t i a l i n t e r v a l s i n which a s i n g l e e m i s s i o n peak dominated t h e spectrum. The a c t i v a t i o n energy f o r t h e hight e m p e r a t u r e peak i s 0.56 eV. The s p a t i a l s e p a r a t i o n of t h e two t r a p s i s p r i m a r i l y due t o a r a p i d d e c r e a s e w i t h d i s t a n c e i n t h e The s p a t i a l d e n s i t y of t h e s h a l l o w e r of t h e two midgap l e v e l s . s e p a r a t i o n i s less complete f o r t h i s s h a l l o w e r l e v e l , g i v i n g rise t o an u n c e r t a i n t y i n t h e a c t i v a t i o n energy which i s found A broad emission s i g n a l of r e v e r s e p o l a r i t y t o be -0.49 e V . a t i n t e r m e d i a t e t e m p e r a t u r e s i s a s c r i b e d t o a h i g h t r a p concent r a t i o n n e a r t h e s i l i c o n s u r f a c e which c o n t r i b u t e s t o a h i g h A t low t e m p e r a t u r e s a t h i r d degree of charge compensation. e m i s s i o n peak of low i n t e n s i t y i s d e t e c t a b l e , w i t h a n a c t i v a t i o n energy of 0.19 eV; t h i s l e v e l is removed by a forming-gas a n n e a l as d i s c u s s e d below. (See S e c t i o n 4.3.2 f o r further discussion of t h i s l e v e l . ) Also shown i n Fig. 4.3 i s t h e f e a t u r e l e s s emiss i o n spectrum f o r a c o n t r o l d i o d e which was f u r n a c e annealed a t 1000°C (30 min, N 2 ) . CC-DLTS emission s p e c t r a are shown i n Fig. 4.4 f o r a diode which r e c e i v e d a forming-gas a n n e a l (15% H2, 85% Np, 450 C , 30 min.). These s p e c t r a are dominated by t h e same two midgap d e f e c t l e v e l s a s found above, w i t h the l e v e l a t -0.49 eV decaying more r a p i d l y t h a n t h e 0.56 e V l e v e l w i t h d e p t h i n t o t h e s i l i c o n subs t r a t e . However, t h e r e is an a d d i t i o n a l e m i s s i o n peak w i t h an a c t i v a t i o n energy of 0.28 e V which i s comparable i n magnitude t o This l e v e l i s w e l l r e s o l v e d i n t h e e m i s s i o n t h e midgap peaks. spectrum and can be unambiguously p r o f i l e d by t h e double-correlat i o n CC-DLTS technique. The d e f e c t - d e n s i t y d i s t r i b u t i o n i s shown i n Fig. 4.5. The h o r i z o n t a l b a r s mark t h e s p a t i a l i n t e r v a l s over which average d e f e c t d e n s i t i e s were measured, and t h e v e r t i c a l b a r s d e n o t e t h e u n c e r t a i n t i e s i n t h e d e n s i t i e s due t o u n c e r t a i n t y i n t h e measurement of t h e n e t e m i s s i o n s i g n a l . The d e f e c t d e n s i t y d e c r e a s e s monotonically w i t h d e p t h below t h e l a s e r - i r r a d i a t e d s u r f a c e , over t h e i n v e s t i g a t e d i n t e r v a l . Also shown i s t h e p r o j e c t e d This approximates t h e range f o r Si' implanted a t 77 keV [4.5]. p r o j e c t e d range of s i l i c o n f o r t h e a c t u a l l y implanted s p e c i e s , SiH+ i f i t is assumed t h a t t h e i o n i z e d molecule d i s s o c i a t e s i n t o s i l i c o n and hydrogen atoms a t t h e s i l i c o n s u r f a c e , w i t h equa p a r t i c l e v e l o c i t i e s a t t h e i n s t a n t of s e p a r a t i o n .
ing and
High d e n s i t i e s of deep l e v e l s remain a f t e r cw laser anneal. ntroduced a n n e a l e f f e c t i v e l y removed t h e l e v e l a t 0.19 eV The a n o t h e r l e v e l a t 0.28 e V . The d e n s i t i e s of a l l 450-C
4. ELECTRONIC DEFECTS IN SILICON
187
25
20
Epitaxial Silicon N-TYw, (100) Self Implanted CW Laser Annealed (1,= 350 C) ' Furnace Annealed: 450 C, 30 rnin. eo = 347 sec-'
I
> E
1 15 6 i7j v)
< io
I-
V
0
5
C
T(K)
FIGURE 4.4. CC-DLTS emission s p e c t r a f o r a self-implanted laser-annealed diode which a l s o r e c e i v e d a 45OoC f u r n a c e anneal. l e v e l s d e c r e a s e w i t h depth i n t o t h e s i l i c o n s u b s t r a t e from > IOl5 i n t h e near-surface r e g i o n , as shown f o r t h e l e v e l a t 0.28 e V i n Fig. 4.5. The p r o j e c t e d range f o r Si' r e v e a l s t h a t t h e meas u r e d d e f e c t s r e s i d e i n material t h a t was n o t d r i v e n amorphous by i o n i m p l a n t a t i o n . Diodes f a b r i c a t e d w i t h cw laser-processed Czochralski-grown s i l i c o n produce t h e same r e s u l t s as shown above f o r e p i t a x i a l material [4.61. E l e c t r o n i c d e f e c t l e v e l s i n scanned electron-beam annealed The test d e v i c e s w e r p-n (SEBA) s i l i c o n are shown i n Fig. 4.6. j u n c t i o n d i o d e s f a b r i c a t e d by i m p l a n t i n g As' (100 keV, 5x101 Z
(B,1-3 a-cm) s i l i c o n . The i o n i n t o Czochralski-grown p-type energy and dose were s u f f i c i e n t t o d r i v e t h e s i l i c o n s u r f a c e amorScanned electron-beam a n n e a l i n g phous t o a depth of -100 nm. was performed w i t h both a commercial electron-beam welder i4.7 1 and a scanning e l e c t r o n microscopy 14.81, w i t h t h e s u b s t r a t e s a t room temperature. The p a j u n c t i o n diodes were completed by d e p o s i t i n g aluminum over t h e r e c r y s t a l l i z e d s u f a c e , u s i n g photol i t h o g r a p h y and plasma e t c h i n g t o d e f i n e t h e mesa s t r u c t u r e , and CC-DLTS s p e c t r a i n Fig. s i n t e r i n g (45OoC, 30 min. forming g a s ) . 4.6 are f o r h o l e emission i n t h e d e p l e t i o n l a y e r of t h e p-type s u b s t r a t e . The h o l e t r a p a t 0.28 e V dominates t h e emission
N. M.JOHNSON
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Epitaxial Silicon N Type. (100) Self Implanted CW L a w A n d & (T, = 350 Cl Furnace Anne* 450 C, 30 min.
1075
t
\
3 -1
Depth (10' \ )
E,-0.28
FIGURE 4.5. S p a t i a l depth p r o f i l e of t h e d e f e c t l e v e l a t eV i n self-implanted laser-annealed s i l i c o n .
s p e c t r a , w i t h an a d d i t i o n a l peak a t 0.36 e V i n t h e SEM-annealed device. These are a c t i v a t i o n e n e r g i e s f o r thermal emission of h o l e s t o t h e s i l i c o n v a l e n c e band. Neither l e v e l i s detectable i n t h e furnace-annealed c o n t r o l diode. The d e f e c t d e n s i t i e s were found t o be i n t h e 1013 cm-3 range t o g r e a t e r d e p t h s t h a n a micron below t h e m e t a l l u r g i c a l j u n c t i o n [4.9]. R e s u l t s f o r cw Ar-ion l a s e r - f a b r i c a t e d p-n j u n c t i o n s are shown i n Fig. 4.7 [4.10]. Boron-doped s i l i c o n wafers were implanted w i t h As' (100 keV, 4 ~ 1 0 1 ~ cw laser annealed w i t h t h e s u b s t r a t e a t 350 C , and f a b r i c a t e d i n t o mesa d i o d e s ; I n Fig. 4.7 f o u r h o l e t r a p s are t h e d i o d e s were n o t s i n t e r e d . d e t e c t a b l e w i t h a c t i v a t i o n e n e r g i e s as f o l l o w s : 0.10 e V , 0.20 e V , 0.28 e V , and 0.45 e V ; t h e s e are denoted as H(0.10), H(0.20), H(0.28), and H(0.451, respectively. The H(0.28) l e v e l is t h e same l e v e l shown i n Fig. 4 . 6 f o r s i n t e r e d d i o d e s and w a s not p r e s e n t when t h e s u b s t r a t e t e m p e r a t u r e was s e t t o 250 C d u r i n g l a s e r annealing. The dominant h o l e t r a p a t 0.45 e V w a s found t o be u n s t a b l e a t room temperature. The d e f e c t d e n s i t y decays w i t h t i m e , w i t h t h e l e v e l t r a n s m u t i n g t o t h e s h a l l o w e r l e v e l a t 0.10 eV. By comparing these d e f e c t s i n laser-annealed s i l i c o n w i t h
" 4. ELECTRONIC DEFECTS IN SILICON
Cz Silicon P-Type, (100) As+ Implanted
(0.28 eV)
n
-> E
-m
1
189
3-
SEBA (SEM)
P-N Junction Hole Traps eo = 347 ~ c - l
-
SEBA (Welder)
C
en iij 2 v)
Furnace Annealed (1000 C, 30 min, N,
4 0 0
1
1-
80 100
150
200
250
300
330
FIGURE 4.6. CC-DLTS spectra for hole emission in As+implanted scanned electron-beam annealed silicon. The spectrum for a furnace-annealed p-n junction diode is also shown. defects produced by furnace annealing followed by rapid cooling, and with other published results, the laser induced defects were identified as interstitial Fe and Fe-B pairs. A comparison of electronic defect levels in laser-annealed and scanned-electron-beam-annealed silicon is shown in Fi 4.8. t4.111. The test devices were As+-implanted (100 keV, 4~101~'cm-~) p-n junction diodes, fabricated with the mesa structure, which were annealed with either a cw Ar-ion laser at a substrate temperature of 250' C or with a scanning electron microscope. Figure 4.8 shows a DLTS spectrum immediately after diode preparation. For cw laser annealing, the hole emission spectrum is dominated by a level at 0.45 eV, as seen previously. As shown with spectra (a) and (b), the density of this level increases rapidly with laser power P, where PM is the power required to melt the silicon.
N. M. JOHNSON
190 1
1
I
I
I
1
1
I
1
I
H(0.281
Si:B (100) 10-20S1-crn As+ 100 KeV 4 x
lo” ~ r n - ~
CW Ar+ LASER ANNEALED
~ = 3 rn 9s BIAS PULSE 0
-
10 V
H(0.45)
H(0.I) I
t
3 5 0 75
1
100
I
125
1
150
I
175
I
XK, TEMPERATURE (K)
t
I
I
225
250
275
FIGURE 4.7. Typical capacitance transient spectrum of ptype si implanted with As+ and annealed by cw Ar-ion laser. (N. H. Sheng and J. L. Merz, Ref. 4.10.) These laser powers are below that required to generate slip When the laser dislocations as discussed in Section 4.3.3. power i s sufficiently high to produce slip dislocation, the dominant defect level is found t o depend on substrate temperature and appears at 0.43 eV for a temperature of 250° C. The SEBA results for slip-free annealing reveal only a weak signal at E,++0.40 eV, which does not depend strongly on electron-beam power. A s discussed above, the 0.45 eV level i s not stable at room temperature and i s associated with interstitial iron. 4.3.2
Ream-Induced Defects
Recrystallization of implanted amorphous layers on singlecrystal silicon involves the epitaxial regrowth of the amorphous layers, with electrical activation of the implanted dopant, and removal of displacement damage in the near-surface region which remains crystalline during ion implantation. It has also been observed that beam irradiation of single-crystal silicon can induce or quench in electronic defects [ 4 . 9 , 4 . 1 2 , 4 . 1 3 ] . This is illustrated in Fig. 4.9 with the electron-emission spectrum for an electron-beam annealed Schottky-barrier diode. The diode was not implanted and received the same beam-annealing schedule used to recrystallize the As+-implanted layer in Fig. 4 . 6 . The electron-emission spectrum for a SEBA-annealed diode is compared
4. ELECTRONIC DEFECTS IN SILICON
I
151
I
100
I
I
125
I
I
150
I
I
1
I
I
1
200 225 TEMPERATURE (K) 175
I
I
250
191
I
I
275
I
FIGURE 4.8. Capacitance transient spectrum of p-type Si implanted with As+ and annealed by either a cw Ar-ion laser (LA), using three different values of laser power, or a scanning elecTS is the substrate temperature during beam tron beam (SEBA). annealing. (N. H. Sheng et al., Ref. 4.11.)
to that for unannealed material. The spectrum for beam-annealed silicon is dominated by two emission peaks with activation energies of 0.19 and 0.44 eV. Shoulders on both peaks indicate additional unresolved emission centers, In the unannealed control, no emission peaks are detectable on the indicated scale of sensitivity. Similar quenched-in deep levels have been observed in cw laser irradiated silicon, where it was further determined that the low-temperature emission peak has an electric-field dependent activation energy with a zero-field value of 0.22 eV [4.14]. The 0.44 eV level has been ascribed to the phosphorus-vacancy complex [4.9,4.13-4.151. The spatial distributions of the quenched-in defects differ in electron-beam irradiated t4.91 and cw laserirradiated silicon t4.131, but in both cases the defects are located in the near-surface region of the silicon. The essential observation is that even without displacement damage from ion implantation the localized heating and rapid quenching, which are inherent features of beam recrystallization, can provide the stimulus for defect formation during laser or electron-beam annealing. However, it has been further observed, with cw laser annealing, that beam-induced damage can be strongly influenced by the crystalline quality of the substrate before laser annealing t4.121. In addition, near the point of melting,
N. M.JOHNSON
192 1.1
I
I
I
I
I
1
I
Schottky Diode Electron Traps h
> E
v
SEBA (Welder)
Unannealed
80 100
150
200
250
300 330
T (K) FIGURE 4 . 9 . CC-DLTS spectrum for electron emission in unimplanted scanned electron-beam annealed silicon. Also shown is the spectrum for the unannealed material. beam-induced slip dislocations are generated as discussed in the next section. 4.3.3
Minimization of Beam Recrystallization Defects
Critical optimization and control of annealing conditions To illustrate, are required for successful beam processing. with a scanned electron beam too low a beam current (for a given accelerating voltage) yields incomplete recrystallization of the implanted amorphous layer, while beam currents near that required for melting generate slip dislocations, which can be readily viewed in an optical microscope as a cross grid of slip lines over the annealed area. These extremes have also been documented in scanned laser annealed silicon. The photoluminescence efficiency of the recrystallized material varies dramatically depending upon the creation or suppression of dislocations. For example, beyond a critical exposure time, at a constant power near that required
4.
ELECTRONIC DEFECTS IN SILICON
193
-400 \
c: Q
-300
2
I-
-200
?? ;
t; W -100
01
I
I
I
I
(x6
0.7
0.8
a9
zj
10
ANNEALING LASER POWER, PIP,
FIGURE 4.1 0. Dependence of short-circuit photocurrent I hot0 in arbitrary units and sheet resistivity p on laser power P !or a diode implanted with 4x1Ol5 As+/cm2 and annealed with a 100 urn laser spot and 6 pm step. Laser power is normalized to the power required for melting, Po. (M. Muzita et al., Ref. 4.17.) for melting, slip appears and the luminescence is dominated by Such studies have dislocation-related defect levels [4.16]. further revealed that the range of laser power over which good quality annealed material can be achieved is limited [4.11,4.17], as further discussed below. Finally, electronic deep levels associated with dlslocations in deformed silicon have been detected by capacitance transient spectroscopy 14.181. Optimum annealing conditions for a cw Ar-ion laser are shown in Fig. 4.10 [4.17,4.19]. Single-crystal silicon with ptype conductivity was implanted with As+ at 100 keV to a dose of 4x1015 The laser annealing conditions were selected €or laterally uniform recrystallization, with laser power a variable. After annealing, the sheet resistance p and shortcircuit photocurrent Iphoto of mesa p-n junction diodes were measured over a range of laser powers P up to that required for melting, Po. The sheet resistance depends on majority-carrier properties, while the photocurrent is sensitive to those of the minority carriers. As shown in Fig. 4.10, for P > O.68Po the sheet resistivity achieves a minimum value of 45 O/square, corresponding to 100% electrical activation of the implanted arsenic.
N. M.JOHNSON
On the other hand, the relative magnitude of the photocurrent increases with laser power to a peak value at 0.76Po and then decreases at higher laser powers. The effect of laser power on electronic-defect density has been directly evaluated with DLTS i4.111. In Fig. 4.11 are plotted defect concentrations versus beam power for scanned-laser and scanned-electron beam annealing. The defect concentration of scanned-electron-beam-annealed diodes is lower than that of laserannealed diodes and is comparatively insensitive to electron-beam power. However, for laser-annealed diodes, the defect concentration increases more than one order of magnitude with increasing laser power. For high-power laser-annealed devices, the defect concentration is comparable to the substrate dopant Concentration. Results from EBIC are also shown in Fig. 4.11. The EBIC chargecollection efficiency decreases as the defect concentration increases. It was found in this study [4.111 that laser-induced defects can extend several microns below the implanted junction. I
L
o
t
I
I
A'"
I 4
V
z
\ A
\
W
w
-0.6 L L
LL W
-
= I
w U w
"
>-
\
2
0
k V
0 LA Ts=250"c H(0.45)
+-04
I
0.6
I
2
R
LA T,~250"c H(0.43) A SEBA Ts:25"c H(0.40) A EBlC ACCELERATING VOLTAGE : 30 KeV 0
I
a8 ANNEALING BEAM POWER, PIP,
I
I .o
FIGURE 4.11. Defect concentration measured by DLTS, as a function of annealing beam power. The broken curve displays the EBIC charge-collection efficiency as a function of laser power. [ N . H. Sheng et al., Ref. 4.11.1
4. 4.4
ELECTRONIC DEFECTS IN SILICON
195
LATERAL NONUNIFORMITIES IN SCANNED-BEAM ANNEALING This section discusses some practical limits in optimizing
cw beam annealing conditions for minimum densities of residual
defects in the recrystallized silicon layer. First, the specific issue of lateral variations in defect density due to scan line separation is discussed with results from luminescence and electronbeam induced current measurements. Then the topic of scanned-beam annealing of silicon wafers with patterned dielectric overlayers is addressed with results from transmission electron microscopy. 4.4.1
Scan Line Separation
Defect luminescence has been used to evaluate cw laserannealed silicon [4.16,4.20,4.21]. In particular, it has been found that when the lateral separation between adjacent laser-scan lines is reduced, the luminescence intensity decreases [4.201. This is demonstrated in Fig. 4.12 with luminescence spectra for n-type Czochralski-grown silicon which was implanted with SiH+ at 80 keV to a fluence of 2x1015 cm-2 and recrystallized at a substrate temperature of 350' C with the scanned beam of a cw Ar-ion laser. Sample A was annealed at 5.2 W with a 40 pm diameter spot scanned at 12 cm/sec and with a 15 pm separation between lines. The two lower spectra correspond to samples laser annealed at 11.5 W, with an 80 pm diameter spot and a scan speed of 5 cm/sec. For sample B the separation of the scan lines was 30 um and for sample C it was 15 pm. Although the spectra are similar, the defect luminescence intensity varies by a factor of 30, with sample C giving the weakest signal. Inspection of the samples in an optical microscope showed that A and B have discernible lateral nonuniformity consisting of stripes between adjacent scan lines, which is not present in sample C . These samples were also evaluated with electron spin resonance. In each sample the characteristic dangling bond resonance of amorphous silicon [4.211 was observed with volume densities (calculated with an estimated thickness of 180 nm for the initial1 amorphized layer) of 4.I x ~ O ~1.~lx1017, , and 0.4x101x1 O1 cm-s for samples A, B, and C , respectively. Thus, as the lateral separation between adjacent scan lines was reduced, the luminescence intensity and spin density decreased and the nonuniformity of the surface was also reduced. These results indicate that the defects are associated with an imperfectly crystallized interface between adjacent laser annealed zones. It was suggested that the luminescence is possible due to extended defects such as dislocations or grain boundaries. Lateral nonuniformities in cw laser annealed silicon have been directly demonstrated with electron beam-induced currents
196
N . M. JOHNSON
FIGURE 4 . 1 2 . Luminescence s p e c t r a of n-type s i l i c o n , showi n g the e f f e c t of changing the l a s e r a n n e a l i n g c o n d i t i o n s . Details of t h e p r e p a r a t i o n of samples A, B y and C are g i v e n i n t h e text.
f 4 . 1 7 1 . A t y p i c a l EBIC image i s shown i n Fig. 4 . 1 3 . The sample was implanted w i t h As' a t 100 keV t o a f l u e n c e of 6x1Ol4 and annealed wich an Ar-ion l a s e r beam w i t h a 40 v m diameter s p o t A f t e r a n n e a l i n g , mesa diodes were f a b r i c a t e d and 30 vm s t e p . by s t a n d a r d p h o t o l i t h o g r a p h i c t e c h n i q u e s , and ohmic metal c o n t a c t s were formed on both s i d e s of t h e diode. The m e t a l l u r g i c a l p-n j u n c t i o n w a s l o c a t i o n 0.2-0.3 u m below t h e implanted s u r f a c e . The E B I C measurement was performed w i t h a 5 keV a c c e l e r a t i n g The E B I C v o l t a g e which r e s u l t e d i n a n e l e c t r o n range of -0.3 vm. image arises from t h e d i f f u s i o n of beam-generated m i n o r i t y c a r riers t o t h e j u n c t i o n where they d r i f t i n t h e e l e c t r i c f i e l d of t h e d e p l e t i o n l a y e r , y i e l d i n g an induced c u r r e n t i n t h e e x t e r n a l c i r c u i t . This c u r r e n t is monitored as the e l e c t r o n beam i s r a s t e r e d over t h e specimen t o produce an image which i s d i s p l a y e d on a cathode r a y tube. The upper and lower p o r t i o n s of Fig. 4 . 1 3 d i s p l a y t h e s t a n d a r d secondary e l e c t r o n image of a diode. In t h i s mode t h e laser-annealed s u r f a c e appears t o be smooth and defect free. However, when t h e E B I C mode i s switched on ( c e n t e r
4. ELECTRONIC DEFECTS IN SILICON
197
FIGURE 4.13. EBIC micrograph ( c e n t e r ) and secondary emiss i o n image ( t o p and bottom) e x c i t e d by 5 keV beam. The sample w a s implanted w i t h 6x1014 As+/cm2 and annealed w i t h 40 pm s p o t I n t h e E B I C mode, dark regions correspond t o and 30 pm s t e p . Ref. 4.17 .] low charge c o l l e c t i o n . [M. Mizuta e t a1
.,
of F i g . 4.13), h i g h - c o n t r a s t dark l i n e s appear which are p a r a l l e l t o t h e d i r e c t i o n of t h e laser s c a n . The EBIC dark-line s p a c i n g i s of t h e o r d e r of t h e l a s e r - s c a n s t e p ; t h e d e v i a t i o n s from cons t a n t s p a c i n g are considered t o be a s s o c i a t e d w i t h complex o v e r l a p e f f e c t s [4.171. F u r t h e r , i t was found t h a t improved r e s u l t s were o b t a i n e d f o r a large-diameter beam (100 pm), small s c a n s t e p (6 pm), and slow scan speed (6 cm/sec), which i s c o n s i s t e n t w i t h t h e photoluminescence r e s u l t s d e s c r i b e d above. The d a r k - s t r i p e f e a t u r e s have a l s o been observed i n x-ray topographs of cw Ar-ion l a s e r - a n n e a l e d , s l i p - f r e e s i l i c o n , w i t h a one-to-one correspondence between t h e dark s t r i p e s i n t h e EBIC image and i n t h e x-ray topograph [4.22]. With TEM i t has been determined t h a t t h e d a r k - s t r i p e r e g i o n s c o n t a i n a h i g h d e n s i t y of It h a s been d i s l o c a t i o n loops due t o i n t e r s t i t i a l atoms [4.23]. f u r t h e r determined t h a t C02 laser a n n e a l i n g (at 10.6 pm wavel e n g t h ) does n o t d i f f e r s i g n i f i c a n t l y from cw Ar-ion laser anneali n g (0.514 pm wavelength) i n t h a t t h e u s e of a scanned h o t s p o t t o a n n e a l i m p l a n t a t i o n damage seems always t o be accompanied by t h e p r o d u c t i o n of r e s i d u a l i n t e r s t i t i a l d i s l o c a t i o n l o o p s , which act t o d e c r e a s e m i n o r i t y - c a r r i e r l i f e t i m e [4.24]; t h i s i s s u e i s f u r t h e r d i s c u s s e d i n S e c t i o n 4.4.2. The above i d e n t i f i e d l i m i t a t i o n s of raster-scanned p o i n t s o u r c e s can be overcome by u s i n g scanned-line s o u r c e s , where t h e l e n g t h of t h e l i n e s o u r c e i s g r e a t e r t h a n t h e maximum l a t e r a l dimension of t h e s u r f a c e area t o be r e c r y s t a l l i z e d .
198 4.4.2
N. M. JOHNSON
P a t t e r n e d Dielectric Overlayers
Most s t u d i e s of e l e c t r o n i c d e f e c t s i n cw beam-recrystall i z e d s i l i c o n have been conducted on material which was uniformly amorphized by i o n i m p l a n t a t i o n and r e c r y s t a l l i z e d over l a t e r a l dimensions l a r g e compared t o t h e examined area. Isolating the r e c r y s t a l l i z a t i o n process i n t h i s way s e r v e s t o f o c u s a t t e n t i o n on b a s i c m a t e r i a l s i s s u e s . However, t h e a c t u a l a p p l i c a t i o n t o integrated-circuit fabrication introduces additional constraints which must be addressed i f beam r e c r y s t a l l i z a t i o n i s t o r e p l a c e conventional furnace annealing. For example, i n i n t e g r a t e d c i r c u i t p r o c e s s i n g only small well-defined areas of a s i l i c o n wafer a r e i o n implanted t o h i g h doses, such as t h e s o u r c e and d r a i n c o n t a c t s i n a metal-oxide-semiconductor f i e l d - e f f e c t t r a n s i s t o r (MOSFET). The implanted area i s o f t e n d e f i n e d by a window etched i n a d i e l e c t r i c t h i n - f i l m o v e r l a y e r . An example i s shown i n Fig. 4.l(c). The p l a n a r p n j u n c t i o n diode i s bounded by a s i l i c o n - d i o x i d e l a y e r , which e l e c t r i c a l l y p a s s i v a t e s t h e s i l i c o n s u r f a c e . This is a more complex s t r u c t u r e f o r beam r e c r y s t a l l i z a t i o n s i n c e t h e oxide s t e p a t t h e edge of t h e implanted amorphous l a y e r a l t e r s both t h e s p a t i a l d i s t r i b u t i o n of energy d e p o s i t e d by t h e scanned beam and t h e temperature p r o f i l e ( o r heat: d i s s i p a t i o n ) d u r i n g r e c r y s t a l l i z a t i o n . Oxide-passivated p - n j u n c t i o n diodes have been examined a f t e r scanned electron-beam a n n e a l i n g w i t h a modified SEM [4.25]. In F i g . 4.14 i s shown a Nomarski i n t e r f e r e n c e o p t i c a l micrograph f o r an SEM-annealed r e g i o n of a n implanted amorphous l a y e r bounded by a s i l i c o n - d i o x i d e o v e r l a y e r . The s t a r t i n g material was Czochralski grown -oriented s i l i c o n w i t h p-type c o n d u c t i v i t y ( B , 1+3 Q-cm). A s i l i c o n - d i o x i d e l a y e r w a s thermally grown t o a t h i c k n e s s of 260 nm and p h o t o l i t h o g r a p h i c a l l y p a t t e r n e d and chemically e t c h e d t o expose areas of s i l i c o n f o r i o n i m p l a n t a t i o n . The e n t i r e wafer w a s t h e n implanted w i t h As+ at 100 keV t o a dose of lX1015 The SEM a n n e a l i n g c o n d i t i o n s were as f o l l o w s : beam v o l t a g e of 20 kV, l i n e - s c a n r a t e of 1 0 msec/line, frame-scan rate of 40 sec/frame, r a s t e r e d area of 1.10x1.15 mm, beam d i a m e t e r of 12+14 pm, s u b s t r a t e temperature of 25OC, and a beam c u r r e n t of 78 PA. The beam c u r r e n t w a s -90% of t h a t r e q u i r e d t o m e l t t h e exposed s i l i c o n . I n F i g . 4.14 t h e SEM-recrystallized r e g i o n of t h e implanted s i l i c o n is d i s t i n g u i s h e d from t h e amorphized a r e a by t h e d i f f e r e n c e i n r e f l e c t i v i t y . I n F i g . 4.15 i s shown a b r i g h t - f i e l d TEU micrograph of t h e boundary between a r e g i o n of implanted and r e c r y s t a l l i z e d s i l i c o n and an o x i d i z e d r e g i o n ; t h e r a s t e r e d e l e c t r o n beam overlapped t h e oxide l a y e r , and t h e oxide was removed f o r TEM. The dark band running d i a g o n a l l y a c r o s s t h e micrograph l o c a t e s t h e boundary between t h e r e c r y s t a l l i z e d and o x i d e - p a s s i v a t e d r e g i o n s and w a s i d e n t i f i e d a s a zone of amorphous material.
4.
ELECTRONIC DEFECTS IN SILICON
199
FIGURE 4.14. Nomarski interference optical micrograph of an oxide-passivated SEM-annealed p-n junction in silicon.
In samples annealed at higher beam currents, the boundary contained a high density of dislocations. In a MOSFET such a defective zone would be particularly detrimental to device operation since it would be situated at the critical interface between the source or drain and the channel. Diodes fabricated in the above material were also evaluated by DLTS, a spectrum from which is The dominant hole trap has an activation shown in Fig. 4.16. ener y of 0.28 eV, as discussed in Section 4.3, with a density of -l0lf cm-3 at a depth of 1 pm below the metallurgical junction. No hole emission peak is detectable on the same scale of sensitivity in the furnace-annealed diode. The results presented above illustrate the added difficulties which can be encountered in the practical application of beam recrystallization in silicon integrated-circuit processing. It is not anticipated that nonuniform recrystallization arising from patterned dielectric overlayers will be alleviated with scannedline sources. However, this is an area of cw beam processing requiring further investigation.
200
N. M . JOHNSON
FIGURE 4 . 1 5 . Bright-field TEM micrograph of the boundary in single-crystal silicon between a region of exposed silicon ~ ~ and a region which was implanted with As+ (100 keV, I x ~ OcmV2) which was covered with a thermally grown layer of silicon dioxide (260 nm thick) during implantation and SEM recrystallization. The silicon-dioxide layer was removed for ‘ E M . 4.5
ISOTHERMAL TRANSIENT ANNEALING
A s reviewed in Chapter 1 , transient isothermal annealing techniques have recently been introduced to activate ion implanted dopants through solid-phase regrowth. The implanted silicon wafer is irradiated by an incoherent light source (e.g., tungsten halogen lamp, arc lamp, or graphite heater), and the damage is annealed on a time scale of seconds. The combination of largearea illumination and long dwell times (as compared to fast pulse or scanned laser annealing) insure that the wafer is heated essentially uniformly. The technique results in minimum redistribution of implanted dopants, complete electrical activation, and uniform annealing across the wafer. Residual electronic defects after transient isothermal annealing have also been investigated r 4 . 2 6 ,
4.271.
4.
I
I
ELECTRONIC DEFECTS IN SILICON
I
I
I
E, = 0.28 eV (N,) = 1.6 X 1014 cm-3 (7100A < X, < 11,200A)
201
I
Cz Silicon P T W e *(loo) As+ Implanted Oxide-Passivated P-N Junctia Hole Traps e, = 347 sec-'
Furnace Anneal
(1000C, 30 min., N,)
I
1
0 100
I
150
I
200
I
250
I 300
1
3:
T (K)
FIGURE 4 . 1 6 . CC-DLTS spectrum for hole emission in an oxide-passivated SEM-annealed p-n junction diode. Also shown is the spectrum for a furnace-annealed control diode. A comprehensive study of electronic defects in silicon after transient isothermal annealing has recently been reported which includes a comparison of different annealing systems and clearly demonstrates the importance of quench rate on the density of residual defects [4.271. A typical trap-level spectrum for a boron-doped epitaxial silicon wafer after transient thermal annealing is shown in Fig. 4.17; the wafer was not ion implanted. After an anneal at 900°C for 20 seconds, a large DLTS peak appears at a temperature of -1 7 K. From the DLTS analysis, the trap concentration NT = 10" ~ m - ~the , trap ionization energy E E + 300 meV, and and the capture cross-section u = 2x10ZlS cg2; this appears to be the same as the % + 0.28 eq level discussed in Section 4.3. The defect level is not detectable in the starting (reference) material, but is present in low concentration after a conventional furnace anneal. In unimplanted Czochralskigrown and float zone grown silicon, the level was also not detectable after the transient thermal annealing (i.e., NT < 1x10'' 100 keV) p-n while in arsenic-implanted ( 5 ~ 1 0 1As+ ~ junction diodes fabricated in Czochralski-grown silicon, the level is resent after transient annealing at a concentration of 5x1Ol2 e m ' . Transient thermal annealing of n-type (phosphorus-doped)
N. M.JOHNSON
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Epi -Silicon (111) PF
p = 2 Qcm
0.9
*
(3
-
0.7
-
NT= 1 x 1 015cm-3 ET = E, +300me V
cn
0.5
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-
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a
DLTS : tl =lms t, = 2 m s
p -Type / Boron
furnace anneaiing (900O C , 30 min)
0.3 0.1
0
Reference
I
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l
r
i
160
l
r
r
l
r
l
l
r
l
200 240 TEMPERATURE
l
l
l
280
l
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K
FIGURE 4.17. DLTS spectrum f o r a boron-doped e p i t a x i a l Also s i l i c o n sample a f t e r t r a n s i e n t thermal annealing ( I T A ) . shown a r e s p e c t r a f o r a f u r n a c e annealed c o n t r o l sample and f o r t h e unannealed ( r e f e r e n c e ) m a t e r i a l . No s i g n i f i c a n t DLTS s i g n a l s i l i c o n wafers was a l s o i n v e s t i g a t e d . w a s d e t e c t e d between 50 K and 300 K, from which i t w a s concluded t h a t no e l e c t r i c a l l y a c t i v e t r a p l e v e l s w i t h c o n c e n t r a t i o n s cm-3 a r e induced i n t h e upper h a l f of t h e g r e a t e r than s i l i c o n bandgap under i s o t h e r m a l t r a n s i e n t a n n e a l i n g c o n d i t i o n s . The dependence of t r a p c o n c e n t r a t i o n on t h e p r i n c i p a l In Fig. 4.18 annealing parameters was a l s o i n v e s t i g a t e d [4.27]. i s shown t h e e f f e c t of a n n e a l i n g temperature on t h e t r a p concent r a t i o n f o r t h r e e d i f f e r e n t a n n e a l i n g systems. The samples i n t h e SEL and AG 210 annealing systems were cooled a t a r a t e of -lOO°C/sec; t h e cooling rate i n t h e Varian I A 200 system was not c o n t r o l l a b l e . A l l t h r e e curves peak at a temperature of -lOOO°C. The r e l a t i v e displacements of t h e t h r e e curves i s considered t o be p r i m a r i l y due t o d i f f e r i n g c o o l i n g rates f o r samples i n t h e t h r e e systems, even f o r nominally t h e same s p e c i f i e d r a t e . This i s f u r t h e r d i s c u s s e d below. The t r a p c o n c e n t r a t i o n s induced a f t e r i s o t h e r m a l t r a n s i e n t annealing i n a s i n g l e system w i t h d i f f e r e n t ramp down rates is
4.
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SEL
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-
2 oz
A G 2'0
I-
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t =20s fast ramp down (50 - 150 'C/S)
iSEL AG 210 1
ramp down not controllable
VA IAN IA200
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203
Epi-Silicon (111)
-
Z
ELECTRONIC DEFECTS IN SILICON
\ \
1
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0 0
\
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n 4 oz
-
I-
-
VARIAN I A 2 0 0 4
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1 \ \
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13
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FIGURE 4.18. Trap c o n c e n t r a t i o n of t h e 300 meV l e v e l v e r s u s a n n e a l i n g t e m p e r a t u r e . The t h r e e curves are measured w i t h d i f f e r e n t systems as i n d i c a t e d . shown i n F i g . 4.19. The rates of sample cooling are not w e l l d e f i n e d because of problems i n t h e measurement of sample temp e r a t u r e under t r a n s i e n t c o n d i t i o n s . The dashed curve ( s o l i d circles) and s o l i d curve (open c i r c l e s ) were taken on s e p a r a t e samples from t h e same w a f e r . The open circles r e p r e s e n t measurement p o i n t s taken a t a slow ramp down of 5OC/sec, and t h e s o l i d circles are p o i n t s t a k e n at a t e n times f a s t e r r a t e . Additiona l l y , f u r n a c e a n n e a l s were performed f o r comparison. The r e s u l t s r e v e a l t h a t t h e production of t r a p s i s s t r o n g l y reduced by a slow ramp down of t h e temperature, independent of t h e a n n e a l method. F u r t h e r o b s e r v a t i o n s on t h e behavior of t h e 300 meV l e v e l are as f o l l o w s [4.27] : The d e f e c t c o n c e n t r a t i o n d e c r e a s e s w i t h i n c r e a s i n g d i s t a n c e from t h e s u r f a c e , which may be caused by t h e temperature g r a d i e n t d u r i n g c o o l down o r may i n d i c a t e t h a t t h e s u r f a c e i s a p o s s i b l e s o u r c e of t h e d e f e c t . The l e v e l can be removed by a f u r n a c e a n n e a l a t temperatures above 400OC. The t r a p l e v e l i s induced by i s o t h e r m a l t r a n s i e n t a n n e a l i n g a f t e r
N . M. JOHNSON
204
Epi-Siiicon (111) p -Type/ Boron p = 2 Qcm
I-
z z
n: t 520s * = fast ram down (SO0 C I S ! o = slow ramp down
F
-
0
(5"C/S)
I-
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#
furnace anneal: t = 30 min A = fast ramp down A = slow ramp down
--0
0
IT
tz w 0 Z
-- -
4-
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4 ; --A\ \
a a
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-
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13
ANNEALING TEMPERATURE TA FIGURE 4.19. Comparison of residual trap concentrations of the 300 meV level after slow and fast ramp down rates. The circles and triangles represent points of transient thermal annealing and furnace annealing, respectively.
, ' s A B + , and Ar+ implantation into boron-doped silicon and also after B+ implantation in Al-doped silicon and transient annealing. It is not present after transient annealing of 30Si+ implanted samples. It was proposed that the defect giving rise to the 300 meV level is a complex involving boron and the lattice vacancy. 4.6
POST-RECRYSTALLIZATION PROCESSING
Can residual electronic defects from transient thermal processing be removed with other techniques which are compatible with integrated-circuit processing and which retain the inherent features of beam recrystallization (as enumerated in Section 4.1)? This possibility may be anticipated since the defects arise from incomplete removal of displacement damage from ion implantation and the rapid quenching which is characteristic of transient thermal annealing. As a consequence, the annealing kinetics of
4.
ELECTRONIC DEFECTS IN SILICON
205
the residual deep levels should be similar to those found for centers introduced by ionizing radiation where it has been shown that most of the defects can be removed at low temperatures (e.g., < 600°C) [4.15,4.28]. 4.6.1
Furnace Anneals
A low-temperature furnace anneal is commonly used for contact alloying in integrated-circuit processing. In Section 4.3 it was noted that an anneal at 450°C was effective in removing the electron trap at 0.19 eV in self-implanted Schottky diodes. The effect has also been demonstrated with hole traps in p-n junction diodes [4.29]. Mesa diodes [Fig. 4.l(b)] were fabricated by implanting As+ into B-doped (1-3 Q-cm) oriented Czochralski-grown silicon and recrystallizing with a scanning cw Ar-ion laser; the substrate was maintained at room temperature (T, = 25°C) during beam annealing. Charge-emission spectra are shown in Fig. 4.20. The spectrum in Fig. 4.20(a) is for a diode which received no further thermal processing after laser recrystallization. The spectrum is dominated by the hole emission center at 0.28 eV, which was discussed previously. Also evident is a broad emission peak of reverse polarity near room temperature. In a p-n junction, the reverse polarity signifies minority-carrier emission [4.11, which in this case arises from electron emission from traps in the upper half of the silicon bandgap. Similar broad emission peaks of reverse polarity were observed in furnace-annealed mesa diodes, while such a feature is absent in oxide-passivated diodes with the planar structure (see Fig. 4.16). These observations suggest that the electron emission is largely due to edge effects in the mesa structure. Figure 4.20(b) shows the spectrum for a diode which was processed simultaneously with that of Fig. 4.20(a) except for the addition of a 450°C furnace anneal which was performed immediately after laser annealing. Both diodes were then aluminum metallized and plasma etched to define the mesa structure. In Fig. 4.20(b) the hole trap at 0.28 eV is absent and the broad emission signal near room temperature has been significantly reduced in intensity. In addition, a new electron emission peak appears corresponding to an electron trap at E, - 0.28 eV; that is, the narrow peak of reverse polarity identifies a deep level with an activation energy of 0.28 eV for thermal emission of electrons to the silicon conduction band. The electron emission intensity is slightly enhanced by using a forward-bias voltage pulse for trap filling which injects electrons (minority carriers) into the p-region of This suggests the contact from the n+ layer [see Fig. 4.l(b)]. that the electron traps are located near the metallurgical junction. This new electron emission peak appears at nearly the same temperature as the hole emission peak and is apparently masked by the more intense hole emission in Fig. 4.20(a). The
N. M.JOHNSON
206
(a) CW Laser Anneal (Ts = 25 E, + 0 . 2 8 eV Emission
h
> E
v
cn
5
0 0 0
I
C) Cz Silicon P-Type, (100) As+ Imp la nted P - N Junction
( b ) CW Laser Anneal (T,= 2 5 C )
Plus 4 5 0 - C Furnace Anneal
Ec- 0.28 eV
o----------v-Electron I
I
Injection Pulse I
I
A
FIGURE 4.20. CC-DLTS emission spectra for cw laser-annealed silicon diodes. In (a) the diode received only a laser appeal, and in (b) the diode also received a 45OOC furnace anneal. removal of the hole trap uncovers the residual electron trap, which requires higher anneal temperatures for removal, as demonstrated below. It was found that the observed reduction in defect density after the low-temperature furnace anneal is accompanied by a decrease in the reverse-bias leakage current in the above diodes. In addition, the observed reduction in defect density realized by a 45OoC furnace anneal after laser recrystallization suggests that a similar reduction should be achievable by laser annealing with a heated substrate. However, diodes which were cw laser annealed with a substrate temperature of 35OoC display the same dominant hole trap at 0.28 eV, with defect densities comparable t o those of diodes with substrates at room temperature during laser recrystallization. The effect of high-temperature furnace annealing after cw laser recrystallization is illustrated in Fig. 4.21 14.61. The results are for Schottky-barrier diodes on epitaxial silicon
4. ELECTRONIC DEFECTS IN SILICON
207
Epitaxial Silicon N-Type, f1OO) Self Implanted CW Laser Annealed Furnace Annealed
10
FIGURE 4.21. E l e c t r o n emission s p e c t r a f o r e p i t a x i a l s i l i con which w a s f u r n a c e annealed a f t e r s e l f i m p l a n t a t i o n and cw laser annealing. The i n s e t shows t h e d e f e c t d i s t r i b u t i o n f o r t h e 0.28 e V i n t h e specimen annealed a t 60OoC. l e v e l EC
-
which were f i r s t self-implanted and laser annealed and then In f u r n a c e annealed a t 6OO0C and 8OO0C b e f o r e m e t a l l i z a t i o n . specimens annealed a t 6OO0C, t h e emission spectrum is dominated by t h e same t h r e e l e v e l s shown p r e v i o u s l y i n Fig. 4.4. Except f o r t h e l e v e l a t -0.49 e V t h e r e i s l i t t l e change i n t h e t r a p This l e v e l has s i g n i f i c o n c e n t r a t i o n s a f t e r a 600°C anneal. c a n t l y d e c r e a s e d i n d e n s i t y r e l a t i v e t o t h e o t h e r s , which s u g g e s t s a low d i s s o c i a t i o n energy f o r t h e corresponding ( u n i d e n t i f i e d ) d e f e c t complex. A f t e r an 8OOOC a n n e a l a l l d e f e c t d e n s i t i e s are s i g n i f i c a n t l y reduced. The i n s e t i n Fig. 4.21 shows t h e s p a t i a l d i s t r i b u t i o n of t h e e l e c t r o n t r a p a t 0.28 e V a f t e r t h e 6OOOC anneal. The p r o f i l e was measured by t h e d o u b l e - c o r r e l a t i o n CC-DLTS technique [4.21, and t h e h o r i z o n t a l b a r s mark t h e s p a t i a l i n t e r v a l s over which t h e l o c a l average d e f e c t c o n c e n t r a t i o n s were measured. Also shown i n t h e i n s e t i s t h e p r o j e c t e d range of 28Si+ [4.5], which approximates t h e p r o j e c t e d range of s i l i c o n f o r t h e a c t u a l l y implanted s p e c i e s , Si$. The d e f e c t d e n s i t y decreases r a p i d l y with depth, again i l l u s t r a t i n g t h a t t h e r e s i d u a l d e f e c t s are l o c a l i z e d n e a r t h e r e c r y s t a l l i z e d s u r f a c e b u t extend i n t o material n o t d r i v e n amorphous by i o n i m p l a n t a t i o n .
K. M. JOHNSON
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P u l s e d - l a s e r a n n e a l i n g s t u d i e s o f f e r f u r t h e r examples of high-temperature f u r n a c e a n n e a l i n g of r e s i d u a l d e f e c t s . A capacit a n c e - t r a n s i e n t spectrum t y p i c a l of t h e d e f e c t s t a t e s i n t r o d u c e d by m e l t i n g s i l i c o n w i t h e i t h e r a 1.06 pm o r 0.53 um Nd:YAG The two major peaks l a b e l e d A laser i s shown i n Fig. 4.22 i4.301. and B r e p r e s e n t s t a t e s i n t h e s i l i c o n bandgap a t 0.19 eV and The 0.33 e V , r e s p e c t i v e l y , below t h e conduction-band minimum. r e l a t i v e c o n c e n t r a t i o n s of t h e s e d e f e c t s were found t o v a r y i n d i f f e r e n t samples, but t h e same d e f e c t states were c o n s i s t e n t l y observed w i t h i n a s u r f a c e l a y e r of t h i c k n e s s e q u a l t o t h e d e p t h A d e n s i t y maximum was observed for the 0.33 e V l e v e l of melting. a t high l a s e r powers which s u g g e s t e d t h a t i t i s quenched-in d u r i n g It was a l s o s u g g e s t e d t h a t t h e 0.19 t h e p u l s e d laser p r o c e s s i n g . e V l e v e l may be due t o sodium i n - d i f f u s i o n from t h e s u r f a c e f4.311. The s t a b i l i t y of t h e s e prominant d e f e c t s t a t e s was examined w i t h a series of 20 min i s o c h r o n a l anneals. Defect s t a t e c o n c e n t r a t i o n (averaged o v e r a depth of 0.7 urn) as a f u n c t i o n of a n n e a l i n g temperature i n a hydrogen atmosphere i s p l o t t e d i n Fig. 4.23. Annealing begins a t 4OOOC and i s complete by 700°C. The r e s u l t s i l l u s t r a t e t h a t c o n v e n t i o n a l f u r n a c e a n n e a l s of 500-700°C can remove t h e dominant r e s i d u a l d e f e c t s a f t e r t r a n s i e n t thermal processing.
3
'c
Sl P-CG-SB n= t .5 x 10'6 cm-3
rt=sms 0.53p,4.4J/cm2
5
200
250
TEMPERATURE ( O K )
FIGURE 4.22. t y p e s i l i c o n melted o r 0.53 wn. Defect i s E(0.33 e V ) . [J.
Capacitance t r a n s i e n t spectrum t y p i c a l of nw i t h a Nd:Yag l a s e r a t wavelengths of 1.06 pm s t a t e A i s i d e n t i f i e d as E(0.19 e V ) , and B L. Benton e t a l . , Ref. 4.30.1
4. ELECTRONIC DEFECTS IN SILICON
209
SI-P-SB X = O 53p.44JIcm2 H2
ANNEALING TEMPERATURE ('C 1
FIGURE 4.23. Schematic representation of 20 min, isochronal annealing in hydrogen ambient of the A and B defect states. [J. L. Benton et al., Ref. 4.30.1 4.6.2
Hydrogenation
Hydrogen passivation of pulsed-laser induced defects is illustrated in Fig. 4.24 [4.32]. Spectrum (a) shows post-recrystallization defect concentrations of Superimposed on the DLTS spectrum is a thermally stimulated capacitance (TSCAP) scan. The steps in capacitance indicate a high concentration of defects and appear when the thermal energy is sufficient t o empty the traps. The increase in capacitance over the original substrate value of 50 pF indicates an increase in carrier concentration corresponding to the introduction of donors by the laser processing. There is no significant change in these measurements after a 4 hr anneal at 3OOOC in molecular hydrogen [Fig. 4.24(b)]. In contrast, monatomic hydrogen neutralizes the electrically active defects at low temperatures. The DLTS spectrum in Fig. 4.24(c) is featureless and the TSCAF' scan is constant and equal to the bulk-silicon value. Both measurements indicate that the hydrogen plasma passivates the quenched-in recombination centers. Equivalent results were obtained with plasma anneals at 100°C and 300°C and times of 10 min to 4 hrs. If the defect reaction
N. M.JOHNSON
210
0"
-
4 -
(X100)
2-
0
4 I
hr. W - C , h.400.C.
'
- 200 - 100
--0
n PUSMA 0.002
Tom
6
TEMPERATURE CK 1
FIGURE 4.24. ( a ) DLTS spectrum and TSCAP s c a n f o r lasermelted phosphorus-doped c r y s t a l l i n e s i l i c o n . ( b ) DLTS spectrum and TSCAP s c a n of laser-melted s i l i c o n taken a f t e r a subsequent 4 h r , 3OO0C a n n e a l i n H2 showing no change i n d e f e c t - s t a t e d e n s i t y . ( c ) DLTS and %CAP s p e c t r a showing p a s s i v a t i o n of e l e c t r i c a l l y a c t i v e d e f e c t s a f t e r t r e a t m e n t i n a 0.38 Torr hydrogen plasma a t 200°C f o r 4 h r . (d) DLTS and TSCAP s p e c t r a produced f o l l o w i n g hydrogen plasma t r e a t m e n t and a subsequent hydrogen e v o l u t i o n , .002 Torr vacuum a n n e a l a t 400°C f o r 1 h r . [ J . L. Benton e t a l . , Ref. 4.32.1 is l i m i t e d by hydrogen i n - d i f f u s i o n , it should be e f f e c t i v e l y A i n s t a n t a n e o u s a t t h e s e temperatures f o r a depth of 1 pm. f i n a l experiment t o evolve t h e hydrogen a f t e r plasma t r e a t m e n t is shown i n Pig. 4.24(d). The reappearance of e l e c t r i c a l l y a c t i v e d e f e c t s a f t e r a vacuum a n n e a l a t 4OO0C for 1 h r a t .002 Torr i s e v i d e n t . The d a t a a g a i n r e v e a l d e f e c t d e n s i t i e s of The changes i n t h e s p e c t r a l f e a t u r e s r e l a t i v e t o the pre-annealed s t a t e a r e not y e t understood. However, t h e r e s u l t s demonstrate t h a t hydrogen p a s s i v a t i o n can be used t o n e u t r a l i z e r e s i d u a l e l e c t r o n i c d e f e c t s a t temperatures of 100-300°C.
4. ELECTRONIC DEFECTS IN SILICON 4.7
211
BEAM-CRY STALLIZED SILICON THIN FILMS
As reviewed in Chapter 5, crystallization techniques such as laser melting are used to increase the grain size of deposited silicon films to dimensions that are comparable to or greater than those of the active device, thereby minimizing the detrimental effects of grain boundaries on device operation. Residual grain and subgrain boundaries intersecting metallurgical p-n junctions in beam-crystallized silicon films have been shown to provide efficient paths for dopant diffusion during silicon processing [4.33,4.341, and in thin-film transistors grain boundaries have been shown to reduce the channel mobility [4.35-4.381. However, crystallization techniques currently under development should permit large-scale device fabrication in single-crystal silicon thin films (see, for example, Refs. 4.38-4.44).
In this section results are presented from investigations of residual electronic defects in beam-crystallized silicon films on amorphous substrates. To date three studies have been reported which specifically focus on the characterization of electronic defects in crystallized silicon thin films [4.45-4.471. All three used the metal-oxide-silicon structure and concluded that the Si-SiO2 interface is the dominant source of residual electronic defects in this material. Deep levels at this interface are generally found to be continuously distributed in energy. Since this feature has not been encountered in previous sections, the measurement of continuous deep-level distributions is introduced in this section and the arguments underlying the above conclusion are presented. 4.7.1
Ow Ar-Ion Laser Crystallized Films
Crystallized-silicon thin films are processed on insulating substrates and generally contain residual grain boundaries. These features introduce unique requirements for electronic defect evaluation. The test device must necessarily be a thin-film structure, and the semiconductor-insulator interface will generally be an integral part of the device. The conventional current-rectifying devices (i.e., Schottky-barrier and p-n junction diodes) used in previous sections are not as readily implemented in thin semiconducting films because lateral placement of electrical contacts generally results in a non-negligible series resistance 14.481. This effect can be minimized through proper selection of device geometry and semiconductor conductivity as discussed in Section 4.7.3. A unique approach compatible with beam crystallization is shown in Fig. 4.25. Since the semiconducting layer can be processed on an insulating thin film, a metal-oxide-semiconductor (MOS) structure can be fabricated which possesses negligible
N. M.JOHNSON
212
I I
I
AE (1%Si)
\
/
I
SiO, SingleCrystal silicon (n+)
FIGURE 4.25. Schematic cross-section of an inverted MOS capacitor consisting of a beam-crystallized silicon thin film on thermally grown silicon dioxide over single-crystal silicon.
series resistance. Such a configuration has been used for capacitance-voltage (C-V) measurements of cw laser-crystallized silicon thin films 14.491. The MOS structure in Fig. 4.25 has been used for DLTS evaluation of defects in cw Ar-ion laser-crystallized silicon films on oxidized bulk silicon f4.451. The device is an inverted MOS capacitor consisting of a silicon thin film deposited and crystallized on a thermally grown layer of silicon dioxide. The single-crystal silicon substrate was degenerately doped for n+ conductivity (resistivity < .003 Q cm) and provided the gate electrode in this inverted structure. For device fabrication, 0.65 um of polycrystalline silicon was deposited by low-pressure chemical vapor deposition onto a 200 nm-thick thermal oxide. Next, the silicon film was implanted with boron at 100 keV to a dose of lxlO1* The continuous film was then encapsulated with 30 nm of silicon nitride, also by low-pressure chemical vapor deposition, in preparation for laser processing. Crystallization was performed with a scanned cw Ar-ion laser with the substrate heated to 50OoC; typical beam conditions were as follows: laser beam power 5 W, beam diameter 80 um, scan speed 50 cmfs, and scan line separation 20 vm. Aftep crystallization, the siliconnitride layer was removed (with phosphoric acid at 180°C). Next, the silicon layer was implanted with boron (10 keV, 5x1014 cm-2) and furnace annealed (9OO0C, 30 min) t o form a shallow p+ layer for ohmic contact to the "backside" of the semiconducting film. The MOS capacitors were completed by vacuum depositing an A1-Si alloy (1% Si) over the crystallized silicon film, photolithographically patterning the metal layer to define back electrodes (500 dia.), plasma etching the exposed silicon film to form the mesa structure depicted in Fig. 4.25, and finally sintering the devices at 45OOC in forming gas (15% Hp and 85% N2) for 20 min. The high-dose boron implant was found to be essential in order to maintain an ohmic back contact over the wide temperature
4.
ELECTRONIC DEFECTS IN SILICON
213
range required for DLTS, and the A1-Si alloy was used to minimize contact alloying during sintering. A crystallized-silicon layer as shown in Fig. 4.25 offers three distinct sources of electronic defect: (1) point and extended defects within grains, (2) defects at grain boundaries, and ( 3 ) defects at the semiconductor-insulator interface. Only the first of these is generally encountered in recrystallized bulk silicon. The latter two differ from the first in that the deep levels are spatially localized and generally distributed continuously in energy in the silicon bandgap. Continuous deeplevel distributions have been demonstrated for grain boundaries in bulk polycrystalline silicon 14.50-4.511 and are generally found at the Si-SiO2 interface on bulk single-crystal silicon f4.521. In general, the above three souces of electronic defect will be detected in a DLTS measurement. As illustrated in previous sections, point defects in single-crystal silicon contribute emission peaks in a DLTS spectrum. On the other hand, a continuous deep-level distribution will produce a nonzero DLTS signal over a wide range of temperatures, corresponding to carrier emission from the distributed deep levels, which results in a smoothly varying, often featureless, DLTS spectrum. In general, the above three sources of emitted carriers will be superimposed in a given DLTS spectrum.
In Fig. 4.26 is shown a TEM micrograph of a cw lasercrystallized silicon layer in an inverted MOS capacitor. As is typical for cw laser crystallization of continuous silicon films, the grains are elongated in the direction of the laser scan with widths in the range of 1 to Spy [4.53,4.54].
In Fig. 4.27 is shown a CC-DLTS spectrum for the p-type laser-crystallized silicon film. The spectrum was obtained by biasing the capacitor in depletion and periodically pulsing the gate voltage into strong accumulation in order to populate deep levels throughout the film with majority carriers (in this case, holes). The emission signal over the entire investigated temperature range was found to be essentially saturated with respect to voltege-pulse amplitude. Failure to achieve saturation could result from partial band bending in the vicinity of grain boundaries during pulse biasing due to high defect densities. In this regard the test structure shown in Fig. 4.25 is nonideal for characterizing grain boundaries since they extend through the depletion layer of the silicon thin film and therefore do not experience a uniform trap filling condition. Saturation of the emission signal suggests that the grain boundaries are not the dominant contribution of emitted holes. In Fig. 4.27 the emission signal is nonzero from the lowest investigated temperature to above room temperature; the upper end of the temperature range generally corresponds to emission of holes from defect
N. M. JOHNSON
214
FIGURE 4.26. TEM micrograph of a s i l i c o n t h i n f i l m which was c r y s t a l l i z e d w i t h a scanned cw Ar-ion l a s e r .
I
-> -E
I
I
1
Polycrystaliine Silicon
I
-10
-
N
400 -CW Laser Crystallized MOS Capacitor, P-Type
I
300-<
.-m
VJ
200-
$
100-
oi
'
I
I
I
FIGURE 4.27 CC-DLTS spectrum and d e f e c t d i s t r i b u t i o n obtained from a n i n v e r t e d MOS c a p a c i t o r f a b r i c a t e d on a cw laserc r y s t a l l i z e d s i l i c o n t h i n f i l m . The s p e c t r o m e t e r d e l a y times f o r hole-emission d e t e c t i o n were t i = 2 m s and t 2 = 4 ms. The d e f e c t d i s t r i b u t i o n was obtained w i t h an assumed c a p t u r e c r o s s - s e c t i o n f o r h o l e s of cm2.
4.
ELECTRONIC DEFECTS IN SILICON
215
levels near midgap. Such spectra are representative of carrier emission from deep levels which are continuously distributed in energy. Such defect distributions exist at insulator-semiconductor interfaces and have been extensively investigated in MOS devices [4.551. These studies have shown that in an MOS capacitor a high-temperature emission peak, as displayed in Fig. 4.27 and corresponding to emission from states near midgap, need not arise from a discrete energy level but rather can result from the onset of minority-carrier emission (in the present case, electrons) through a slowly varying continuous distribution of states at the semiconductor midgap. Thus, over the entire temperature interval there is no definite evidence of carrier emission from discrete levels. It was therefore proposed from the above considerations that defect levels at the Si-SiO2 interface are the dominant source of emission centers in the crystallized silicon thin film, with possible secondary contributions from grain boundaries and defects within grains. This issue is discussed in more detail in Section 4.7.3. For the purpose of numerically estimating the defect density, if it is assumed that interface states are the sole contribution to the emission signal, the spectrum in Fig. 4.27 can be analyzed directly to obtain the defect distribution, which is shown in the same figure. The high-temperature peak in the emission spectrum is not included in the defect distribution because the analysis assumes that majority-carrier emission dominates the DLTS transient response [4.55]. The energy scale was calculated with an assumed value of lx10-14 cm-2 for the hole capture cross section, which is based on similar measurements of interface states in MOS capacitors on p-type bulk singlecrystal silicon i4.561. The defect density varies slowly in the mid-10" eV-l cmm2 range, with a slight peak in the distribution at approximately 0.3 eV above the silicon valence-band maximum. Against such a high background signal from interface defects, the density of a spatially uniform discrete level would have to be > lx1014 cm-3 to be detectable. 4.7.2
Strip-Heater Crystallized Silicon Films
Residual electronic defects have also been investigated in graphite strip-heater crystallized silicon thin films 14.461. For electrical characterization inverted MOS capacitors were fabricated as described above. High-frequency capacitance-voltage characteristics ielded effective fixed charge densities in the oxide of < 2x10g crnv2. Trap-emission spectra, recorded with CC-DLTS on both p-type and n-type capacitors, indicated a continuous distribution of deep levels throughout the silicon bandgap. Since the films were essentially single crystal with a low density of subgrain boundaries [4.40,4.41], the Si-SiO2 interface was
216
N. M. JOHNSON
considered t o be t h e p r i n c i p a l s o u r c e of t h i s deep-level continuum. The e f f e c t i v e i n t e r f a c e - s t a t e d e n s i t y w a s found t o cm-2 n e a r midgap which c l o s e l y approaches be < 2 . 5 ~ 1 0 ~ e' V - l the state-of-the-art f o r t h e r m a l l y o x i d i z e d bulk s i n g l e - c r y s t a l s i l i c o n . An e l e c t r o n emission peak w i t h an a c t i v a t i o n energy 0.35 e V was a l s o d e t e c t e d above t h e background continuum of and may be t h e s i g n a t u r e of a bulk d e f e c t w i t h an e q u i v a l e n t l x d 3 ~ m - ~ a, l t h o u g h t h e d e f e c t s p a t i a l l y uniform d e n s i t y of was not chemically o r m i c r o s c o p i c a l l y i d e n t i f i e d .
-
-
A contaminant which has been d e t e c t e d i n g r a p h i t e s t r i p h e a t e r c r y s t a l l i z e d s i l i c o n is carbon i4.571. It has been r e p o r t e d f4.281 t h a t t h i s impurity when i n c o r p o r a t e d i n t h e s i l i c o n l a t t i c e as an i n t e r s t i t i a l c o n t r i b u t e s a hole-emission deep l e v e l a t 0.27 eV and as a carbon i n t e r s t i t i a l - c a r b o n s u b s t i t u t i o n a l complex y i e l d s a hole-emission l e v e l a t 0.36 eV. The absence of any c l e a r l y d e f i n e d peaks i n DLTS hole-emission s p e c t r a f o r g r a p h i t e s t r i p - h e a t e r c r y s t a l l i z e d s i l i c o n i4.461 as summari z e d above, a r g u e s a g a i n s t t h e presence of h i g h c o n c e n t r a t i o n s of e l e c t r i c a l l y a c t i v e carbon i n such f i l m s .
4.7.3
Laser-Crystallized
S i l i c o n Films on Fused S i l i c a
R e s i d u a l e l e c t r o n i c deep l e v e l s i n f u l l y processed cw laserc r y s t a l l i z e d s i l i c o n t h i n f i l m s on f u s e d s i l i c a have been meas u r e d by t r a n s i e n t c a p a c i t a n c e s p e c t r o s c o p y , supplemented w i t h c a p a c i t a n c e - v o l t a g e t e c h n i q u e s [4.471. The t e s t d e v i c e s were MOS t h i n - f i l m c a p a c i t o r s w i t h p-type c o n d u c t i v i t y . A schematic c r o s s - s e c t i o n of t h e c a p a c i t o r s on f u s e d s i l i c a i s shown i n Fig. 4.28. A s c a n n i n g C02 laser was used t o c r y s t a l l i z e s i l i c o n f i l m s (0.5 pm t h i c k ) on bulk f u s e d s i l i c a s u b s t r a t e s w i t h t e c h n i q u e s p r e v i o u s l y shown t o produce s i n g l e - c r y s t a l material c o n t a i n i n g low d e n s i t i e s of l i n e a r d e f e c t s i n areas (termed " i s l a n d s " ) t h a t were p r e - s e l e c t e d f o r subsequent d e v i c e f a b r i c a t i o n [4.38,4.42, 4.43,4.581. The dopant f o r t h e p+ r e g i o n s beneath t h e s u b s t r a t e c o n t a c t s was i n t r o d u c e d by i o n i m p l a n t a t i o n ( B + , ~ x ~ O ~ ~ 50 keV) i n t o t h e c r y s t a l l i z e d s i l i c o n i s l a n d s p r i o r t o o x i d a t i o n . The i s l a n d s were t h e n o x i d i z e d i n d r y 02 a t 1000°C t o grow a 130nm t h i c k g a t e oxide; t h i s a l s o s e r v e d t o d r i v e i n and a c t i v a t e t h e c o n t a c t dopant. Dual i o n i m p l a n t a t i o n of boron ( i .e., 2 . 8 ~ 1 0 ~ ~ cm-2 a t 60 keV and 4x1Ol2 a t 170 keV) followed by a 900°C a n n e a l (30 min i n N2) w a s used t o dope t h e s i l i c o n f i l m p-type i n t h e gate r e g i o n w i t h a n e a r l y uniform boron p r o f i l e i n t h e s p a c e charge l a y e r immediately beneath t h e thermal oxide and a h i g h dopant d e n s i t y n e a r t h e back s u r f a c e of t h e f i l m . This dopant p r o f i l e was e s s e n t i a l t o minimize t h e series r e s i s t a n c e of t h e s i l i c o n s u b s t r a t e i n t h e s e l a t e r a l l y c o n t a c t e d d e v i c e s . A f t e r opening c o n t a c t windows i n t h e o x i d e , an A1-Si a l l o y ( 1 % S i ) w a s vacuum evaporated and p h o t o l i t h o g r a p h i c a l l y p a t t e r n e d
4. ELECTRONIC DEFECTS IN SILICON
Substrate Contact
217 Substrate
AQ (1 % Si) SiO,
P+
1
I
I
c - Si Island
I
P -type
-II
I
P+
I
I---------
P+
Fused Silica Substrate
FIGURE 4.28. Schematic diagram of an MOS c a p a c i t o r f a b r i c a t e d i n a l a s e r - c r y s t a l l i z e d s i l i c o n t h i n f i l m on bulk fused silica. t o d e f i n e g a t e and s u b s t r a t e c o n t a c t s . The g a t e area was d e t e r mined by an i s l a n d width of 40 um and a g a t e width of 50 pm. The l a s e r - c r y s t a l l i z a t i o n c o n d i t i o n s were n o t optimized f o r t h i s i s l a n d width, which i s l a r g e r than t h a t used i n t h e previous s t u d i e s r e f e r e n c e d above, and consequently t h e s i l i c o n f i l m s g e n e r a l l y contained some r e s i d u a l subgrain boundaries and microtwins. The MOS c a p a c i t o r s were completed w i t h a s i n t e r i n formi n g gas a t 45OOC f o r 30 min. The s e p a r a t e e l e c t r i c a l c o n t a c t s t o t h e s u b s t r a t e on opposite s i d e s of t h e g a t e e l e c t r o d e i n Fig. 4.28 were used t o v e r i f y t h e low r e s i s t a n c e of t h e l a t e r a l l y c o n t a c t e d s u b s t r a t e s and were commonly connected f o r c a p a c i t a n c e measurements. From C-V measurements t h e dopant c o n c e n t r a t i o n , i s i n good agreement w i t h was e s t i m a t e d t o be 2x1Ol6 ~ m - ~ which t h e average near-surface boron c o n c e n t r a t i o n a s determined from SUPREM 111 process s i m u l a t i o n s 14.591. The t e s t c a p a c i t o r s were e v a l u a t e d w i t h t r a n s i e n t capacit a n c e spectroscopy. Hole-emission DLTS s p e c t r a a r e shown i n The t r a n s i e n t c a p a c i t a n c e measurements were performed Fig. 4.29. A t each a t f i x e d temperatures over t h e range from 8 0 t o 300 K. temperature, t h e C-V c h a r a c t e r i s t i c was recorded and used t o s e l e c t t h e g a t e b i a s r e q u i r e d t o maintain t h e same e q u i l i b r i u m device c a p a c i t a n c e i n d e p l e t i o n , C D , over t h e e n t i r e temperature range, and t r a n s i e n t waveform averaging was r e q u i r e d t o a t t a i n a u s e f u l t r a p emission s e n s i t i v i t y . The gate-voltage adjustment was r e q u i r e d t o maintain a c o n s t a n t CD s i n c e t h e C-V c h a r a c t e r i s t i c of an MOS device s h i f t s w i t h temperature due t o t h e temperature dependence of t h e Fermi energy and t h e presence of i n t e r f a c e s t a t e s [ 4 . 5 2 ] . This s h i f t i s e s s e n t i a l l y p a r a l l e l t o t h e v o l t a g e a x i s and r e s u l t s i n t h e high-frequency c a p a c i t a n c e changing w i t h
N. M. JOHNSON
218 4
i I DLTS Delay Times:
I
75 100
I 150
I
I
I
200
250
300 325
T (K)
FIGURE 4.29. DLTS s p e c t r a f o r a l a s e r - c r y s t a l l i z e d t h i n f i l m s i l i c o n MOS c a p a c i t o r . The s p e c t r a were recorded w i t h two d i f f e r e n t e q u i l i b r i u m d e v i c e c a p a c i t a n c e s i n d e p l e t i o n , CD. The e r r o r b a r r e p r e s e n t s twice t h e s t a n d a r d d e v i a t i o n of t h e DLTS s i g n a l , which i s formed as t h e d i f f e r e n c e between t h e average c a p a c i t a n c e s measured a t d e l a y times ti and t 2 a f t e r a t r a p f i l l i n g voltage pulse. temperature i f t h e d e p l e t i o n g a t e b i a s i s h e l d c o n s t a n t ; t h i s i s avoided by h o l d i n g CD c o n s t a n t w i t h t e m p e r a t u r e , which s i m p l i f i e s t h e i n t e r f a c e - s t a t e a n a l y s i s 14.551. The two s p e c t r a i n F i g . 4.29, r e c o r d e d w i t h d i f f e r e n t v a l u e s of CD, have a f u r t h e r s i g n i f i c a n c e . The e m i s s i o n peak a t t h e high-temperature end of t h e DLTS s p e c t r a i n F i g . 4.29 could be t h e s i g n a t u r e of a d i s c r e t e energy l e v e l s i t u a t e d n e a r midgap, o r i t could a r i s e from t h e continuous d i s t r i b u t i o n of deep l e v e l s which i s r e v e a l e d by t h e smoothly v a r y i n g , non-zero s i g n a l over t h e r e s t of t h e spectrum. A s d i s c u s s e d p r e v i o u s l y , p o i n t d e f e c t s i n s i n g l e - c r y s t a l s i l i c o n c o n t r i b u t e emission peaks i n a DLTS spectrum, w h i l e a continuum w i l l produce a non-zero DLTS s i g n a l over a wide range of t e m p e r a t u r e s corresponding t o c a r r i e r emiss i o n from t h e d i s t r i b u t e d deep l e v e l s . The dominant s o u r c e of a continuous deep-level d i s t r i b u t i o n i n t h e p r e s e n t material i s t h e Si-SiO2 i n t e r f a c e . In o r d e r t o decide between bulk v e r s u s i n t e r f a c e d e f e c t s as t h e o r i g i n of t h e e m i s s i o n peak, t h e second spectrum w a s recorded w i t h i t s e q u i l i b r i u m d e v i c e c a p a c i t a n c e i n d e p l e t i o n , cD2, chosen such t h a t cD2 > cD1. A s s e e n i n F i g . 4.29, t h e r e s u l t s are ( 1 ) a lower peak h e i g h t and ( 2 ) t h e peak s h i f t s t o a lower t e m p e r a t u r e (i.e., TD2 < T D ~ ) . I n c r e a s i n g CD reduces t h e magnitude of t h e s i l i c o n s u r f a c e p o t e n t i a l w i t h c o r r e s p o n d i n g d e c r e a s e s i n t h e width of t h e s i l i c o n space-charge l a y e r and i n t h e magnitude of t h e e l e c t r i c - f i e l d i n t e n s i t y a t every p o i n t
4. ELECTRONIC DEFECTS IN SILICON
219
within this depletion layer 14.601. If the emission peak arose from a discrete level spatially distributed in the bulk of the silicon, then the reduced depletion width would mean fewer defect centers contributing to the emission signal and consequently a lower peak height, as is observed. However, this would not cause the temperature of the emission peak to shift. For hole emission from a bulk discrete level, the change in the electric field would have to mediate the shift of the peak temperature when CD is changed. This follows from the fact that the trap filling conditions were identical for the two spectra and from the experimental determination that the DLTS signals which comprise the spectra were saturated with respect to both gate-voltage pulse amplitude and pulse width. Thus, it is the emission cycle, rather than the capture cycle, that is being affected by changing CD. For hole emission from a discrete level, the peak appears at a temperature T which is related to the DLTS delay times and trap parameters as follows: (4.3) where e-l is the hole emission time, y is the attempt-to-escape P frequency for trapped holes, EA is t\e activation energy for thermal emission, and k is the Boltzmann constant. Since the right-hand side of Eq. (4.3), which is commonly termed the DLTS emission rate window (as discussed in Section 4.2.2), is a constant, a change in Tp must be compensated by a change in EA and/or up. Furthermore, since Tp depends logarithmically on yp, a strong functional dependence of yp on field would be required, which is not generally observed for discrete deep levels. However, electric-field modulation of the activation energy is a welldocumented phenomenon. Specifically, Poole-Frenkel emission from a neutral center, which for hole emission would be an acceptorlike defect, yields the following field-dependent activation energy (for a one-dimensional coulombic defect potential) 14.611:
where F is the magnitude of the electric field, Eo is the zerofield activation energy, q is electronic charge, and E~ is the silicon permittivity. (A defect center that is neutral when empty and chargedwhen occupied will not experience this effect, as discussed In Section 4.3.2.) Equation (4.4) predicts that the activation energy increases with decreasing electric field. From Eq. (4.31, this requires that the peak temperature increase, rather than decrease, with an increase in CD. Another mechanism which would yield the same qualitative conclusion is high-f ield tunnel emission which has been documented for several point defects in 111-V semiconductors (see, €or example, Ref. 4.62). Thus, the
N. M . JOHNSON experimental r e s u l t s are i n c o n s i s t e n t w i t h f i e l d - m o d u l a t e d emission from a d i s c r e t e l e v e l .
hole
The dependence on CD of t h e high-temperature emission peak i n Fig. 4.29 i s c o n s i s t e n t w i t h t h e e f f e c t of s u r f a c e g e n e r a t i o n on t h e DLTS measurement of i n t e r f a c e s t a t e s n e a r t h e s i l i c o n mid-gap [4.55]. In t h i s c a s e , a t t e m p e r a t u r e s corresponding t o t h e DLTS d e t e c t i o n of m a j o r i t y - c a r r i e r e m i s s i o n ( i . e . , h o l e s ) from i n t e r f a c e states n e a r midgap, m i n o r i t y - c a r r i e r emission from t h o s e s t a t e s becomes a c o m p e t i t i v e p r o c e s s and can i n t r o d u c e a peak i n a DLTS spectrum due t o s u r f a c e g e n e r a t i o n of charge through a slowly v a r y i n g continuous deep-level d i s t r i b u t i o n . The dependence on CD i s due t o t h e change of t h e s u r f a c e p o t e n t i a l w i t h gate voltage. The s u r f a c e p o t e n t i a l d e t e r m i n e s t h e i n t e r s e c t i o n of t h e Fermi l e v e l w i t h t h e i n t e r f a c e which s p e c i f i e s t h e e q u i l i b rium o c c u p a t i o n of t h e i n t e r f a c e s t a t e s ; s t a t e s above t h e Fermi l e v e l do not change occupancy d u r i n g t h e c a p a c i t a n c e t r a n s i e n t ( i n p-type c a p a c i t o r s ) and t h e r e f o r e cannot c o n t r i b u t e t o t h e DLTS s i g n a l . Reducing t h e magnitude of t h e s u r f a c e p o t e n t i a l , by i n c r e a s i n g t h e e q u i l i b r i u m d e p l e t i o n c a p a c i t a n c e , pinches o f f t h e s u r f a c e g e n e r a t i o n p r o c e s s which both a t t e n u a t e s t h e e m i s s i o n peak and s h i f t s i t t o lower t e m p e r a t u r e s , a s d i s p l a y e d i n Fig. 4.29. Based on t h e arguments p r e s e n t e d above, t h e e n t i r e DLTS spectrum i n Fig. 4.29 may be a s c r i b e d t o a continuous d i s t r i b u t i o n Then t h e i n t e r f a c e - s t a t e of deep l e v e l s a t t h e Si-SiO2 i n t e r f a c e . d e n s i t y can be e s t i m a t e d w i t h a s t a n d a r d a n a l y s i s from t h e DLTS s i g n a l a t t e m p e r a t u r e s below t h e emission peak where h o l e emission dominates 14. 5 51 . For example, t h e DLTS s i g n a l j u s t below t h e peak ( i . e . , n e a r 210 K) y i e l d s a v a l u e of approximately 6x1010 eV-l cm-2 f o r t h e d e n s i t y of i n t e r f a c e s t a t e s a t e n e r g i e s n e a r t h e s i l i c o n midgap. Against t h i s background s i g n a l from i n t e r f a c e s t a t e s , t h e d e n s i t y of a s p a t i a l l y uniform d i s c r e t e l e v e l would have t o be > lx1014 cm-3 t o be d e t e c t a b l e i n t h e c a p a c i t o r s examined i n t h i s study. To conclude, t h e r e s u l t s of t h i s s t u d y provide a d e f e c t - s p e c i f i c v e r i f i c a t i o n of t h e device-grade q u a l i t y of C02 l a s e r - c r y s t a l l i z e d s i l i c o n t h i n f i l m s on f u s e d s i l i c a .
4.8
SUMMARY AND CONCLUSIONS
From t h e e x p e r i m e n t a l s t u d i e s reviewed i n S e c t i o n s 4.34.6, t h e f o l l o w i n g g e n e r a l o b s e r v a t i o n s may be drawn r e g a r d i n g e l e c t r o n i c d e f e c t s i n cw t r a n s i e n t thermal annealed bulk s i n g l e crystal silicon: 1.
R e s i d u a l e l e c t r o n i c d e f e c t s remain a f t e r t r a n s i e n t t h e r m a l a n n e a l i n g w i t h a l l i n v e s t i g a t e d forms of energy sources. The d e f e c t d e n s i t i e s are i n e x c e s s of t h o s e o b t a i n e d by c o n v e n t i o n a l f u r n a c e a n n e a l i n g .
4. ELECTRONIC DEFECTS IN SILICON
221
2.
These defects are located in the near-surface region. In recrystallized silicon the defects extend with decreasing density into material not driven amorphous by ion implantation.
3.
Even in material which has not been ion implanted, electronic defects are quenched in during transient thermal annealing.
4.
Beam-annealing conditions (e.g., beam power and scan line separation) can be optimized for recrystallizing implanted amorphous layers with a minimum density of residual defects. However, the use of a raster-scanned point source to anneal implantation damage seems always to be accompanied by the production of laterally-distributed residual defects.
5.
Laterally nonuniform recrystallization can also result from wafer topography (e.g., patterned dielectric overlayers), which introduces lateral variations in energy deposition and heat dissipation during beam recrystallization.
6.
Post-recrystallization processing can be used to reduce the density of residual defects. Furnace anneals in the range of 400-700°C and hydrogen passivation at temperatures of 100-3OO0C have been shown to be effective in removing deep levels.
From the standpoint of electronic defects, of the several cw transient thermal processing techniques that have now been investigated, isothermal transient annealing is the most promising as a replacement for conventional furnace annealing in silicon integrated-circuit technology. Of course, this technique most closely approaches a furnace anneal in its thermal treatment of a silicon wafer. This is further reflected in the similarity of the deep levels which result from these two annealing procedures and in the convergence of their residual defect densities as the quench rate in isothermal transient annealing is reduced. On the other hand, directed energy sources, such as scanned laser and electron beams, may offer unique advantages for materials modifications other than recrystallization as well as provide a controlled means of introducing defects in silicon for fundamental studies. For beam-crystallized silicon thin films on amorphous substrates, the results presented in Section 4.7 indicate that the Si-SiO2 interface is the dominant source of residual electronicdefects in this materials system. Since the measurements were performed on MOS devices in which the interface was a source of
222
N. M.JOHNSON
d e t e c t a b l e d e f e c t s , i n t e r f a c e s t a t e s e s t a b l i s h e d a h i g h background s i g n a l f o r d e t e c t i n g d e f e c t s i n the bulk of t h e s i l i c o n t h i n f i l m s . This background r e q u i r e d t h a t t h e d e n s i t y of a s p a t i a l l y uniform bulk d e f e c t be g r e a t e r t h a n t y p i c a l l y lx1014 cm-3 t o be d e t e c t a b l e . T h i s e f f e c t i v e s e n s i t i v i t y t o bulk d e f e c t s i s r a t h e r low f o r s t u d i e s d i r e c t e d toward t h e c o r r e l a t i o n of r e s i d u a l Novel d e f e c t s w i t h materials p r o c e s s i n g and device o p e r a t i o n . deep-level s p e c t r o s c o p i c t e c h n i q u e s can be combined w i t h m u l t i p l e t r a n s i s t o r s ) t o a c h i e v e high s e n s i t e r m i n a l MOS d e v i c e s (e.g., t i v i t y f o r bulk d e f e c t d e t e c t i o n , i n p a r t by e x c l u d i n g any conIt may be a n t i c i p a t e d t h a t t r i b u t i o n from t h e i n t e r f a c e 1 4 . 6 3 1 . such t e c h n i q u e s and d e v i c e s w i l l be extended t o e l e c t r o n i c - d e f e c t c h a r a c t e r i z a t i o n i n beam-crystallized s i l i c o n t h i n f i l m s .
4. ELECTRONIC DEFECTS IN SILICON
223
REFERENCES 4.1 4.2
4.3
4.4 4.5
4.6
4.7
4.8 4.9
4.10 4.11 4.12
4.13 4.14
4.15 4.16 4.17 4.18 4.19
J. Appl. Phys. 45, 3014 and 3023 (1974). Lang, D. V., Johnson, N. M., B a r t e l i n k , D. J., Gold, R. B., and J . F. Gibbons, J. Appl. Phys. 50, (1979). L e f e v r e , H. and Schulz, H., Appl. Phys. 12, 45 (1977). Johnson, N . M., Gold, R. B., and J. F. Gibbons, Appl. Phys. L e t t . 34, 704 (1979). Gibbons, J. F., Johnson, W. S., and Mylroie, W. W., "Proj e c t e d Range S t a t i s t i c s , Semiconductors and R e l a t e d Mate r i a l s , " 2nd ed. Halstead, New York (1975). Gold, R. B., L i e t o i l a , A . , and Gibbons, Johnson, N. M., J. F., "Laser-Solid I n t e r a c t i o n s and Laser P r o c e s s i n g 1978," (S. D. F e r r i s , H. J. Leamy, and J. M. Poate, eds.) AIP Conference Proceedings, pp. 550-555. American I n s t i t u t e of P h y s i c s , New,York (1979). R e g o l i n i , J. F., Gibbona, J. F., Sigmon, T. W., Pease, R. F. W., Magee, T. J., and Peng, J., Appl. Phys. L e t t . 34, 410 (1979). B a r t e l i n k , D. J., Ratnakumar, K. N . , Pease, R. F. W., Johnson, N . M., and Meiqdl, J. D., Appl. Phys. L e t t . 35, 463 (1979). Johnson, N. M., R e g o l i n i , J. L., B a r t e l i n k , D. J., Gibbons, J. F., and Ratnakumar, K. N.,. Appl. Phys. L e t t . 36, 425 (1980). Sheng, N. M. and Merz, J. L., "Laser and E l e c t r o n Beam Int e r a c t i o n s w i t h S o l i d s , " (B. R. Appleton and G. K. Celler, eds.) pp. 313-318. E l s e v i e r , New York (1982). Sheng, N. H., Mizuta, M., and Merz, J. L., Appl. Phys. L e t t . 4 0 , 68 (1982). Mizuta, M., and Merz, J. L., "Laser and Sheng, N. H., Electron-Beam S o l i d I n t e r a c t i o n s and Materials Processi n g , " ( J . F. Gibbons, I,. D. Hess, and T. W. Sigmon, eds.) pp. 155-162. E l s e v i e r , New York (1981). Chantre, A . , Kechouane, M., and Bois, D., "Laser and Electron-Beam I n t e r a c t i o n s w i t h S o l i d s , " (B. R. Appleton and G. K. Celler, eds.) pp. 325-330. E l s e v i e r , New York (1982). Chantre, A., Kechouane, and Boise, L, "Defects i n Semiconductors 11," (S. Mahajan and J. W. C o r b e t t , eds.), pp. 547-551. E l s e v i e r , New York (1983). IEEE Trans. Nucl. S c i . NS-23, 1497 Kimerling, L. C., ( 1976). Uebbing, R. H., Wagner, P., Baumgart, H., and Q u e i s s e r , H. J., Appl. Phys. L e t t . 37, 1078 (1980). Mizuta, M., Sheng, N. H., Merz, J. L., L l e t o i l a , A., Gold, R. B., and Gibbons, J. F., Appl. Phys. L e t t . 37, 154 (1980). Klmerling, L. C. and P a t e l , Appl. Phys. L e t t . 34, 73 (1979). Sheng, N. H., and Merz, J. L., Appl. Phys. Mizuta, M., L e t t . 38, 453 (1981).
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N. M . JOHNSON S t r e e t , R. A , , Johnson, N. M., and L i e t o i o l a , A., "Laser and Electron-Beam P r o c e s s i n g of Materials," (C. W. White Academic P r e s s , New York and P. S. Peercy, eds.) p. 435. (1980). S t r e e t , R. A., Johnson, N. M., and Gibbons, J. F., J. Appl. Phys. 50, 8201 (1979). and P h i l l i p p , F., "Laser Rozgonyi, G. A., Baumgart, H., and Electron-Beam S o l i d I n t e r a c t i o n s and M a t e r i a l s Processi n g , " ( J . F. Gibbons, L. D . Hess, and T. W. Sigmon, e d s . ) , pp. 193-199. E l s e v i e r , New York (1981 ) Baumgart, H., Hildebrand, O., P h i l l i p p , F a , and Rozgonyi, G . A , , "Proceedings of 2nd Oxford Conference on Microscopy of Semiconductor Materials." I n s t i t u t e of P h y s i c s , London (1981). Baumgart, H., P h i l l i p p , F., and Leamy, H. J., "Laser and Electron-Beam I n t e r a c t i o n s w i t h S o l i d s , " (B. R. Appleton E l s e v i e r , New York (1982). and G. K. Celler) pp. 355-360. Johnson, N. M., S i n c l a i r , R., and Moyer, M. C., IEEE Trans. Elec. Dev. ED-27, 2199 (1980). Benton, J . L., C e l l e r , G. K., Jacobson, A. C., Kimerling, L. C., Lischner, D. J., Miller, G. L., and Robinson, McC., "Laser and Electron-Beam I n t e r a c t i o n s w i t h S o l i d s , " (R. R. Appleton and G. K. C e l l e r , eds.) pp 765-770. E l s e v i e r , New York (1982). P e n s l , G., S c h u l z , M., Johnson, N. M., Gibbons, J. F., and Hoyt , J., "Energy Beam-Solid I n t e r a c t i o n s and T r a n s i e n t Thermal P r o c e s s i n g , " (J. C. C. Fan and N. M. Johnson, eds.) i n p r e s s . E l s e v i e r , New York (1984). "Radiation E f f e c t s i n Semiconductors Kimerling, L. C., 1979." I n s t . Phys. Conf. S e r . 31, B r i s t o l (1977). Johnson, N. M., B a r t e l l i n k , D . J., Moyer, M. D., Gibbons, J. F., L i e t o i l a , A., Ratnakumar, K. N., and R e g o l i n i , J . L., "Laser and Electron-Beam P r o c e s s i n g of Materials," (C. W. White and P. S . Peercy, eds.) pp. 423-429. Academic Press, New York (1980). Benton, J. L., Doherty, C. J., F e r r i s , S . D., Kimerling, L. C., Leamy, H. J., and Celler, G. K., "Laser and E l e c t r o n Beam P r o c e s s i n g of Materials," (C. W. White and P. s. Peercy, eds.) pp. 430-434. Academic P r e s s , New York (1980). Kimerling, L. C. and Benton, J. L., "Laser and E l e c t r o n B e a m P r o c e s s i n g of Materials,'* (C. W. White and P. s. Peercy, eds.) pp. 385-396. Academic Press, New York (1980). Benton, J. L., Doherty, C. J., F e r r i s , S . D., Flamm, D. L., Kimerling, L. C., and Leamy, H. J., Appl. Phys. L e t t . 36, 670 (1980). Johnson, N. M., B i e g e l s e n , D . K., and Moyer, M. D., Appl. Phys. L e t t . 38, 900 (1981). Maby, E. W., Atwater, H. A., Keigler, A. L., and Johnson, N. M., Appl. Phys. Lett. 43, 482 (1983).
.
4. 4.35 4.36
4.37
4.38
4.39 4.40
4.41
4.42 4.43
4.44 4.45 4.46
4.47
4.48 4.49 4.50 4.51
4.52
ELECTRONIC DEFECTS IN SILICON
225
Kamins, T. I. and Von Herzen, B. P., IEEE Elec. Dev. L e t t . EDL-2, 313 ( 1 9 8 1 ) . Ng, K. K., Celler, G. K., P o v i l o n i s , E. I., Frye, R. C., Leamy, H. J., and Sze, S. M., IEEE Elec. Dev. L e t t . EDL-2, 316, ( 1 9 8 1 ) . Frye, R. C. and Ng, K. K., "Grain Boundaries i n Semicond u c t o r s , " (C. H. Seager, G. E. P i k e , and H. J . Leamy, eds.) E l s e v i e r , New York ( 1 9 8 2 ) . pp. 275-286. B i e g e l s e n , D. K., Johnson, N. M., Nemanich, R. J., Moyer, M. D., and Fennel, L. E., "Laser and Electron-Beam I n t e r a c t i o n s w i t h S o l i d s , " (B. R. Appleton and G. K. Celler, eds.) pp. 331-336. E l s e v i e r , New York ( 1 9 8 2 ) . Fan, J. C. C., Geis, M. W., and Tsaur, B.-Y., Appl. Phys. L e t t . 3 8 , 365 ( 1 9 8 1 ) . Maby, E. W., Geis, M. W., LeCoz, L. Y., S i l v e r s m i t h , D. J. Mountain, R. W., and A n t o n i a d i s , D. A., IEEE Elec. Dev. Lett. EDL-2, 241 ( 1 9 8 1 ) . Geis, M. W., Smith, H. I., Tsaur, B.-Y., Fan, J. C. C., Maby, E. W., and A n t o n i a d i s , D. A . , Appl. Phys. L e t t . 4 0 , 158 ( 1 9 8 1 ) . Hawkins, W. G., Black, J. G., and G r i f f i t h s , C. H., Appl. Phys. L e t t . 4 0 , 319 ( 1 9 8 2 ) . F e n n e l l , L. E., Moyer, M. D., B i e g e l s e n , D. K., Chiang, A., and Johnson, N. M., "Energy Beam-Solid I n t e r a c t i o n s and T r a n s i e n t Thermal P r o c e s s i n g , " (J. C. C. Fan and N. M. Johnson, eds.). E l s e v i e r , New York ( 1 9 8 4 ) i n p r e s s . A t w a t e r , H. A., Thompson, C. V., Smith, H. I., and Geis, M. W., Appl. Phys. L e t t . 4 3 , 112 ( 1 9 8 3 ) . Johnson, N. M., Moyer, M. D., and F e n n e l l , L. E., Appl. Phys. L e t t . 4 1 , 562 ( 1 9 8 2 ) . Johnson, N. M., Moyer, M. D., F e n n e l l , L. E., Maby, E. W., and A t w a t e r , H., "Laser-Solid ' I n t e r a c t i o n s and T r a n s i e n t Thermal P r o c e s s i n g of Materials," ( J . Narayan, W. L. Brown, E l s e v i e r , New York and R. A. Lemons, eds.) pp. 491-497. (1983). Johnson, N. M. and Moyer, M. D., "Comparison of Thin Film T r a n s i s t o r s and SO1 Technologies," (H. W. Lam and M. J. Thompson, eds.) E l s e v i e r , New York ( 1 9 8 4 ) i n p r e s s . Wiley, J. D. and Miller, G. L., IEEE Trans. Elec. Dev. EDL-22, 265 ( 1 9 7 5 ) . Kamins, T. I., Lee, K. F., and Gibbons, J. F., IEEE Elec. Dev. L e t t . EDL-1, 5 ( 1 9 8 0 ) . Seager, C. H., P i k e , G. E., and Ginley, D. S., Phys. Rev. L e t t . 4 3 , 532 ( 1 9 7 9 ) . Werner, J., J a n t s c h , W., Froehner, K. H., and Q u i e s s e r , H. J., "Grain Boundaries i n Semiconductors, " (C. H. Seager , G. E. P i k e , and H. J. Leamy, eds.). E l s e v i e r , New York (1982). N i c o l l i a n , E. H. and B r e w s , J. R., "MOS Physics and Technology." Wiley, New York ( 1 9 8 2 ) .
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4.53 4.54
4.55 4.56 4.57 4.58
4.59
4.60 4.61 4.62
4.63
G a t , A . , Gerzberg, L., Gibbons, J. F., Magee, T. J., Peng, Appl. Phys. L e t t . 33, 775 (1978). and Hong, J. D., Johnson, N. M., B i e g e l s e n , D. K., and Moyer, M. D., "Laser and Electron-Beam S o l i d I n t e r a c t i o n s and Materials Processi n g , " (J. F. Gibbons, L. D. Hess, and T. W. Sigmon, eds.) pp. 463-470. E l s e v i e r , New York (1981). J. Vac. Sci. Techn. 2 1 , 303 (1982). Johnson, N. M., Johnson, N. M., B i e g e l s e n , D. K., and Moyer, M. D., J. Vac. Sci. Techn. 1 9 , 390 (1981). P i n i z z o t t o , R. F., Lam, H. W., and Vaandrager, B. L., Appl. Phys. Lett. 40, 388 (1982). Johnson, N. M., B i e g e l s e n , D. K., Tuan, H. D., Moyer, M. D., and F e n n e l l , L. E., I E E E Elec. Dev. L e t t . EDL-3, 369 (1982). Chiang, A . , Z a r z y c k i , M. H., Meuli, W. P., and Johnson, N. M., "Energy B e a r S o l i d I n t e r a c t i o n s and T r a n s i e n t Thermal ElP r o c e s s i n g , " (J. C. C. Fan and N. M. Johnson, eds.). s e v i e r , New York (1984). "Physics of Semiconductor Devices," 2nd ed. Sze, S . M., Chap. 7. Wiley, New York (1981). Milnes, A. G., "Deep I m p u r i t i e s i n Semiconductors," Chap. 5. Wiley, New York (1973). Makram-Ebeid, S., "Defects i n Semiconductors," (J. Narayan, E l s e v i e r , New York and T. Y. Tan, eds.) pp. 495-501. (1981). C o l l e t , M. G., S o l i d - s t a t e Elec. 1 8 , 1077 (1975).
CHAPTER 5
Beam Recrystallized Polycrystalline Silicon: Properties, Applications, and Techniques K. F. Lee", T. J . Stultzt and James F. Gibbons STANFORD ELECTRONICS LABORATORIES STANFORD UNIVERSITY STANFORD. CALIFORNIA
5.1 RECRYSTALLIZATION OF THINPOLYCRYSTALLINE FILMSWITH A SCANNING cw LASER . . . . . . . . . . . . . . . . . . 5 . 2 RESISTIVITYREDUCTIONIN BEAM-RECRYSTALLIZED
228
.................................
247
PROPERTIES OF LASERRECRYSTALLIZED 5.3 ELECTRONIC FILMSINTENDED FOR MOS DEVICE FABRICATION ...................................... AND 5.4 CHARACTERISTICS OF MOS DEVICES INTEGRATED CIRCUITS FABRICATED ON LASER
250
POLYSILICON FILMS
RECRYSTALLIZEDPOLYSILICON FILMS. . . . . . . . . . . . . . . . . .
260
5.5 IMPROVEMENTS IN THE RECRYSTALLIZATION
PROCESS. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . REFERENCES. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
278 335
Heavily doped polycrystalline silicon (polysilicon) is a material that is widely used in present day silicon integrated circuit technology for gates and interconnection lines in MOS integrated circuits. Initial interest in the beam annealing of polysilicon arose because of its potential as a process that could be used to obtain resistivity reduction in as-deposited films. Experiments undertaken to explore that potential are described in Sec. 5.2. These experiments showed that beamrecrystallized polysilicon films have electronic properties that closely approximate those of single crystal material. This result ultimately led to successful attempts to fabricate MOS transistors and integrated circuits directly in beam recrystallized polysilicon 15.1-5.41. A substantial interest now exists in the potential of beam recrystallized silicon-on-insulators (SOI) as a substrate for integrated circuit fabrication. Our development of this subject will begin with a discussion in Sec. 5.1 of the mechanisms by which thin films of polycrystalline Si are recrystallized by a scanning circular cw laser beam. In Sec. 5.2 we will consider the basic majority carrier electrical properties of the recrystallized material and establish the limits *Present address: AT&T Bell Laboratories. Holmdel, New Jersey 07733 ?Present address: TS Associates, San Jose, California 95128. SEMICONDUCTORS AND SEMIMETALS, VOL. 17
227
Copynght 0 1984 by Academic Press, Inc. All rights of reproduction in any form reserved. ISBN 0-12-752117-8
K . F. LEE ET AL. In on r e s i s t i v i t y r e d u c t i o n t h a t are p o s s i b l e w i t h t h i s p r o c e s s . Sec. 5.3 w e w i l l d e s c r i b e t h e b a s i c e l e c t r o n i c p r o p e r t i e s of rec r y s t a l l i z e d f i l m s i n t e n d e d f o r u s e as MOSFET s u b s t r a t e s , l e a d i n g t o a d i s c u s s i o n of t h e e l e c t r i c a l c h a r a c t e r i s t i c s of d e v i c e s fabWe conclude t h e r i c a t e d d i r e c t l y i n t h e s e f i l m s i n Sec. 5.4. c h a p t e r w i t h a d i s c u s s i o n of a v a r i e t y of improvements i n t h e r e c r y s t a l l i z a t i o n p r o c e s s t h a t have been developed t o eliminate v a r i o u s d e f e c t s i n d e v i c e and f i l m p r o p e r t i e s t h a t are o b t a i n e d when a c i r c u l a r beam i s used f o r t h e r e c r y s t a l l i z a t i o n p r o c e s s .
5.1
RECRYSTALLIZATION SCANNING CW LASER
OF
THIN
POLYCRYSTALLINE
FILMS
WITH A
I n t h i s s e c t i o n w e d e s c r i b e f i r s t t h e a n n e a l i n g of t h i n polys i l i c o n f i l m s i r r a d i a t e d by a c w argon l a s e r , w i t h emphasis on t h e r e l a t i o n between l a s e r power and t h e s u r f a c e morphology, c r y s t a l s t r u c t u r e , dopant d i s t r i b u t i o n and m a j o r i t y c a r r i e r e l e c t r i c a l p r o p e r t i e s of t h e annealed f i l m s . Three laser power l e v e l s (low, medium, and h i g h ) were i d e n t i f i e d i n e a r l y s t u d i e s d i r e c t e d toward u n d e r s t a n d i n g t h e r e c r y s t a l l i z a t i o n mechanisms i n p o l y s i l i c o n f i l m s . The g e n e r a l e f f e c t s It can be s e e n from t h i s f i g u r e t h a t a r e d e s c r i b e d i n Fig. 5.1. complete dopant a c t i v a t i o n o c c u r s a t low power l e v e l s , s i m i l a r t o t h o s e l e a d i n g t o complete dopant a c t i v a t i o n i n ion-implanted s i n CONDITIONS
OPTICAL
TEM
SIMS
LOW POWER
I
NO CHANGE IN SURFACE MORPHOLOGY
MEDIUM POWER
I I
HIGH POWER
INCREASE IN SURFACE ROUGHNESS
SURFACE SMOOTHING
I I / I NO INCREASE IN GRAIN
SLIGHT INCREASE IN GRAIN SIZE AT THE
DRASTIC INCREASE IN
REDISTRIBUTION NO DOPANT
DOPANT REDISTRIBUTION AT T H E SURFACE
UNIFORM DOPANT REDISTRlBUTlON
COMPLETE DOPANT LOW MOBILITY
I
COMPLETE DOPANT ACTIVATION WITH INCREASED MOBILITY
I
COMPLETE DOPANT ACTIVATION WITH HIGH MOBILITY
FIGURE 5.1. Summary of t h e e f f e c t s of l a s e r power on t h e ann e a l i n g of p o l y s i l i c o n and t h e corresponding measurement methods.
5.
BEAM RECRYSTALLIZED POLYCRYSTALLINE SILICON
229
gle-crystal silicon. A moderate increase in grain size occurs with dopant redistribution at somewhat higher irradiation intensities. In the medium power range, partial melting causes a growth of surface crystallites and significant dopant redistribution in the molten region. Annealing at the highest power levels results in complete melting of the (thin) polysilicon layer, with an attendant drastic increase in the grain size and uniformity of the doping distribution across the film. A sharp reduction in the electrical resistivity of the film is observed in this high power recrystallization stage. In what follows we provide experimental details on this latter process and also describe features that are observed where a single line scan is compared with an overlapped scan frame. 5.1.1
Sample Preparation
The polysilicon films that are typically used for these studies are 0.5-0.6 pm thick layers deposited by low pressure chemical vapor deposition (LPCVD) on a layer of thermally grown silicon dioxide In some cases the silicon dioxide layer is replaced with a silicon nitride layer. After deposition, the wafers may be doped by either ion implantation or thermal diffusion and then subjected to the beam processing operations. For the laser processing experiments to be discussed below, a cw argon laser was operated in the multiline mode with a laser beam width on the sample surface of 40 pm. This circular spot was scanned across the surface at a rate of 12.5 cm/sec in all cases. The variables employed were substrate temperature, laser power, and the amount of overlap between adjacent scan lines. For reasons that are suggested in Chapter 2 it is advantageous so that relato use subtantial substrate thermal bias (350'C) tively small increments in laser power will produce substantial increments in the surface temperature of the irradiated film. For the experiments reported below, laser power was increased in 1 W increments from 5 W to 13 W, corresponding to a range of surface temperatures from 600' to 1500'C.
It should be emphasized that, while these conditions can produce very desirable changes in the properties of the films, they are certainly not the only set of conditions that would achieve these results; a different laser spot size, scan rate, laser power and substrate temperature can be chosen to yield very similar results. 5.1.2
Growth Mechanisms [5.5]
To study the mechanisms of crystallite growth, Gerzberg et al. L5.53 implanted arsenic and boron ions into separate samples to serve both as dopants and markers which could be used to investigate the recrystallization process. SEM was used to monitor
230
K. F. LEE ETAL.
changes in surface texture as a function of laser power (Fig. 5.2). For the arsenic implanted samples annealed at low and medium laser power levels (5 and 9 W respectively), no significant change in surface morphologywas observed, as shown in Figs. 5.2(a) and ( b ) . Above 10 W, however, distinct lines were apparent in the nonoverlapped sample with the microstructure inside the line pointing in the direction of the scan. Figure 5.2(c) is a SEM micrograph of an annealed line obtained in this high power region, showing the directionality in the line, the smoothing of the original fine surface roughness, and the formation of large surface structures. The grain size obtained at the high power level was analyzed through transmission electron microscopy ('JXM) and diffraction analysis. Samples were jet thinned from the back surface before presents T E M micrographs of boronexamination. Figure 5.3 implanted polysilicon samples irradiated with both single and overlapping lines. The long crystallites in the single line scan are characteristic of the high power region. These crystallites are found to be continuous, columnar structures extending from the Si02 interface to the surface of the film i5.61. The difference between the single and overlapped scanning conditions will be explained below in terms of thermal conductivity variations.
CONTROL AND 5 w
*P
-
9w
12
w
-
2P FIGURE 5.2. SEM photomicrographs of laser processed polysilicon at power levels of 5 W, 9 W and 1 2 W.
5.
BEAM RECRYSTALLIZED POLYCRYSTALLINE SILICON
231
FIGURE 5.3. TEM photomicrographs of single and overlapping lines on boron-implanted polysilicon.
Figure 5.4 is a series of TEM photomicrographs of arsenicimplanted polysilicon, laser annealed at low-, medium-, and highpower levels. The as-deposited samples display fine grain structure with an average grain size of -550 A Laser processing at 7 and 9 W produces grain growth but no significant change in surface morphology compared to the as-deposited samples. However, a sharp change is observed at power levels of 10 W and above, as evidenced by the 15 to 25 pm grain size in the sample annealed with 11 W laser power. The mechanism for this dramatic growth in crystallite size and its implications are discussed in Sec. 5.1.5.
.
232
K. F. LEE ETAL.
FIGURE 5.4. TEM photomicrographs of single and overlapping lines on arsenic-implanted polysilicon.
A global description of the recrystallization process can be gained from plots of the average grain size vs. laser power for the boron- and arsenic-implanted samples shown in Figs. 5.5 and 5.6. In Fig. 5. 5, the grain size of the boron-implanted samples is seen to increase monotonically in the medium power range ( 6 to 8 W). The transition to the high power level occurs in the single line scan when the power is greater than 9 W. In a narrow power range from 8 to 9 W, grain size increases from 1 pm to 10 vm and, at a laser power level of 11 W, the grain size (length) increases t o 20 urn.
Figure 5.6 is a plot of grain size for the arsenic-implanted polysilicon films after laser processing. A s was the case for the boron-implanted samples, grain size increases monotonically with laser power from 6 to 9 W. The transition to the high power level occurs between 9 and 10 W, above which a 20 vm grain size is commonly observed.
5.
BEAM RECRYSTALLIZED POLYCRYSTALLINE SILICON
ipm)
I
I
I
1
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1
233 I
B IMPLANTED POLY AND L.A. n OVERLAPPING LINES NON-OVERLAPPING SCAN w
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OVERLAPPING LINES
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w
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Average grain size of arsenic-implanted laserFIGURE 5.6. annealed polysilicon as a function of annealing power levels for single and overlapping lines.
234 5.1.3
K. F. LEE E T A L .
Dopant Distribution
Secondary ion-mass spectrometry (SIMS) was used to obtain dopant-profile measurements corresponding to the various power levels described in the previous section. Oxygen primary ionbombardment and positive secondary ion-mass spectrometry (B’ and A s + ) were employed, using the CAMECA IMS-3f ion microanalyzer. The boron profiles were determined through standard depth profiling; however, sensitive measurements of the A s profiles required secondar ion initial kinetic discrimination to reduce the amount of (30Si$9 Si. 1 60)+ detected at the same nominal mass as 75As+. This analysis was sufficient to provide an A s detection limit of 3 to 5 x10l8 atoms/cm3. A s can be seen in Fig. 5.7, the high oxygen content of Si02 and minor surface-charging effects produce an anornolously high mass-75 intensity that misrepresent the true arsenic content in the oxide region. This problem, however, did not hinder the measurements in the polysilicon layer.
9w
11w
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-
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lozo
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FIGURE 5.7. SIMS profile of arsenic-implanted polysilicon after laser annealing at different power levels. Inserts are the corresponding TEMS for 9 and 11 W.
5.
BEAM RECRYSTALLIZED POLYCRYSTALLINE SILICON
235
A s can be seen in Fig. 5.7, the dopants did not move from their as-implanted positions at a laser power of 6 W. The profile is approximately Gaussian and can be adequately predicted from Lindhard-Scharff-Schiott (LSS) range calculations [5.7]. Increased laser power results in deeper impurity penetration into the polysilicon.
When the power level was raised to 11 W, the dopant profiles became flat throughout the polysilicon layer. The inserts in Fig. 5.7 are the TEM micrographs measured at 9 and 11 W. It can be seen that, although partial melting occurred at 9 W, the grains are still small (- 3 pm). Their growth is considerable (- 20 pm) only when the melt front reaches the underlying oxide layer. Figure 5.8 shows the boron concentration in the polysilicon layer after laser irradiation. This boron redistribution further substantiates the theory of surface melting. (The artificial boron intensity observed in the oxide could be the result of a change in the ionization yield of the boron atoms implanted in the oxide.) The boron distribution for a 5 W laser power level is identical to the as-implanted distribution and, again, it can be calculated from the LSS theory. A s the power is increased, the region wherein the boron profile is flat also increases. At 8.5 W, substantial redistribution of the boron over 0.2 um is observed but below this depth the boron distribution remains essentially unchanged. lOOKeV 11 1 1
5.1014 Bf
z
0
F 1019
a K !-
z w 0
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0
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z
Ig:1017 .2
.4 DEPTH ( 1 1
.6
FIGURE 5.8. SIMS profile of boron-implanted polysilicon after laser annealing with different power levels.
236 5.1.4
K. F. LEE ETAL. Electrical Properties
Sheet resistivity, carrier concentration, and Hall mobility were measured as a function of laser power by means of a Van der Pauw configuration. The sheet resistance, mobility, and activecarrier concentration (N,) obtained from these measurements for the arsenic-implanted laser-annealed polysilicon are shown in Fig. 5.9. The most important observation is that the carrier concentration is fixed at an approximate level of 3.5 x lOI5 cmq2 (the cm-*), independent of laser power. implantation dose was 5 x This figure also shows that, even a t the lower power level (5 to 6 W), for which the SIMS data indicate no diffusion, the dopants are almost all active. This is in agreement with work on cw laser and electron beam annealing of ion-implanted single crystal silicon, wherein the implanted dopant can be activated with no diffusion by a solid phase recrystallization process. At the medium-power level (6 to 9 W), sheet resistance decreases because of greater
20a
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7 8 9101112 LASER POWER ( W i
6
-u
FIGURE 5.9. Sheet resistance, Hall mobility, and total active carriers of 5 x l d 5 As+ implanted in polysilicon with 170 keV and cw laser annealed as a function of laser power.
5.
BEAM RECRYSTALLIZED POLYCRYSTALLINE SILICON
237
mobility, with no change in the number of active carriers. This increase in mobility drops at 9 W and, as the laser power is raised to 10 W, a sharp transition to the high-power region is accompanied by a rapid rise in mobility. In this region, as power is raised to 10 to 12 W, the trend toward greater mobility decreases, but with no effect on N,.
In the initial experiments with boron-doped samples, the laser power could not be raised above 8.5 W without peeling the polysilicon from the oxide. As a result, the high-power region, where the mobility approaches that of single-crystal silicon, is missing in Fig. 5.10. It is important to note that, in an earlier experiment [5.6] where the polysilicon was deposited over Si3N4, implanted with the same boron dose and recrystallized at 11 W, a grain size of approximately 25 vm, resistivity of 269 Q/o , and mobility of 44.5 cm2/V-sec were achieved. This means that, when polysilicon is deposited over nitride, it is possible to reach the high-power level without causing the film to peel. After this experiment was performed, the problem of peeling of the polysilicon from the oxide was solved by encapsulating the polysilicon with a thin film of Si02 or Si3N4. NS i B/cm2' 7.1014
i
5.1014
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3.1014 1.1014
PH 20
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5
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3
FIGURE 5.10. Sheet resistance, Hall mobility, and total active carriers of 5 x 1014 B+ implanted in polysilicon with 100 keV and cw laser annealed as a function of laser power.
238
K. F. LEE E T A L .
These o b s e r v a t i o n s suggest t h e importance of capping t h e p o l y s i l i c o n f i l m p r i o r t o r e c r y s t a l l i z a t i o n . F o r t h e experiments r e p o r t e d h e r e , i t was found t h a t a t h i n Si-jNq-related cap could be grown by c a r r y i n g o u t a 1100°C N2 anneal f o r 1 hour immediately p r i o r t o t h e laser a n n e a l i n g . The laser power l e v e l could then be i n c r e a s e d without t h e occurrence of p e e l i n g [5.21. D e t a i l e d studi e s of capping procedures and t h e i r e f f e c t s on f i l m i n t e g r i t y are d e s c r i b e d i n Sec. 5.1.8.
5.1.5
Discussion of t h e Basic R e s u l t s
The r e s u l t s o b t a i n e d from t h e g r a i n s i z e a n a l y s i s , dopant red i s t r i b u t i o n and e l e c t r i c a l p r o p e r t i e s provide t h e b a s i s f o r a proposed e x p l a n a t i o n f o r t h e e f f e c t of l a s e r power on p o l y s i l i c o n a n n e a l i n g and r e c r y s t a l l i z a t i o n . To summarize, t h r e e l a s e r power l e v e l s have been i d e n t i f i e d : This range i s c h a r a c t e r i z e d by no change i n ( a ) Low Power. s u r f a c e morphology, no i n c r e a s e i n g r a i n s i z e o r dopant r e d i s t r i b u t i o n and complete dopant a c t i v a t i o n w i t h low c a r r i e r m o b i l i t y . Laser h e a t i n g w i t h o u t m e l t i n g i s r e s p o n s i b l e f o r t h e s e r e s u l t s , under c o n d i t i o n s s i m i l a r t o CW l a s e r a n n e a l i n g of ion-implanted s i n g l e c r y s t a l s i l i c o n . A power l e v e l of 5 W i s c h a r a c t e r i s t i c of t h i s region. ( b ) Medium Power. When t h e l a s e r power i s r a i s e d from 5 t o 9 W, t h e e l e c t r i c a l p- r o-p e r t i e s i n t h e t o p l a y e r a r e changed. Surface roughness i n c r e a s e s a s does g r a i n s i z e , and dopant d i f This region i s f u s i o n occurs i n the top 1000 t o 3000 8 . marked by g r e a t e r m o b i l i t y , reduced r e s i s t i v i t y , and a c t i v e implanted dopants. A l l of t h e d a t a s u g g e s t t h a t t h e p o l y s i l i c o n Because t h e l a y e r begins t o melt from t h e s u r f a c e downward. a c t u a l a n n e a l i n g t i m e is on t h e o r d e r of 1 msec, i t is d i f f i c u l t t o e x p l a i n such phenomena, p a r t i c u l a r l y t h e i m p u r i t y r e d i s t r i b u t i o n , w i t h simple s o l i d - s t a t e d i f f u s i o n .
It i s important t o n o t e t h a t , a l t h o u g h m e l t i n g o c c u r s , no s i g n i f i c a n t change i n g r a i n s i z e is observed. This may be exp l a i n e d by a c r y s t a l l i z a t i o n p r o c e s s t h a t begins i n t h e f i n e grained p o l y s i l i c o n s u b s t r a t e t h a t serves a s the nucleating layer. ( c ) High Power. A l l e l e c t r i c a l p r o p e r t i e s change dramatiIn t h i s narrow c a l l y i n the power range between 9 and 10 W. r e g i o n , t h e r e is a s i g n i f i c a n t change i n s u r f a c e s t r u c t u r e , a very l a r g e i n c r e a s e i n g r a i n s i z e (from 1 t o 20 vm), a s t e e p rise i n carrier m o b i l i t y , and a uniform dopant d i s t r i b u t i o n throughout t h e l a y e r . Large g r a i n s develop when t h e l a y e r i s completely molten a l l t h e way t o the u n d e r l y i n g s u b s t r a t e .
5.
BEAM RECRYSTALLIZED POLYCRYSTALLINE SILICON
239
Complete melting appears to be essential for the development of such grains. If samples are lightly doped and complete melting cannot occur without damage to the samples, then large grains will not develop. 5.1.6
Crystal Orientation of the Beam Recrystallized Films [5.8]
In this section we describe the structural changes produced by the laser recrystallization as they are measured by X-ray diffraction; and we further discuss the effect of the heat treatment and encapsulation techniques that are often required before laser irradiation can be employed. Data are presented for polysilicon films deposited on both silicon dioxide and silicon nitride since these are the most common films employed in present technology. For these experimezts films of low pressure CVD polysilicon Samples were deposited at 625 C to a thickness of -5500 A. of each type were then thermally annealed at 1100°C in flowing nitrogen. The samples were laser processed with the cw scanning argon laser under the melting conditions just described so as t o form large grain structures. A conventional X-ray diffractometer was used to analyze these samples. The results of the analysis are collected in Table 5.1. Measurements were taken with the samples mounted both parallel and perpendicular to the laser scan direction, but no anisotropy was seen. The results reported are the average of the two measurements. The intensities of the signals from the (111)-, (220)(311)-, (400)-, and (331) peaks were measured and normalized to indicate the relative amounts of crystallites with (111)-, (110)-, (311)-, (loo)-, and (33l)orientation in each film. The measured signals were normalized to account for the signal strengths expected in a thick, randomly oriented sample, and also for the finite film thickness [5.91.
While the surface structure observed in polysilicon by optical microscopy does not always correlate with the grain structure, some general observations can be made from Table 5.1 and Fig. 5.11. The films which were not laser processed had a fine-grain appearance. Those laser processed, but not thermally annealed, (14-1, 16-1, 16-3) appeared to have a feather-like, long-grain structure, while those which were also thermally annealed (13-3, 15-1) indicated a large-grained appearance on the surface. One sample (13-1) indicated a less regular appearance, probably related to a mixed structure obtained because of less complete recrystallization. The normalized X-ray results are also shown in Table 5.1. A s expected [5.9], the (110) orientation is dominant in all the films
240
K. E LEE ET AL.
TABLE 5.1. Crystallographic Orientations in Laser Recrystallized Polysilicon. Sample
I---
164
Insulating Film
SigN4
16-1 16-3
I-
sio2
FIGURE. 5.11. Dark-field photomicrographs of laser-recrystallized polysilicon on Si3N4 [(a) and (b)] and on Si02 [(c) and (d)]; ( b ) and (d) show samples thermally annealed before laser recrystallization.
5.
BEAM RECRYSTALLIZED POLYCRYSTALLINE SILICON
241
immediately after deposition, with the normalized intensity being about ten times larger than that from tny other orientation. Thermal annealing of the samples at 1100 C decreases the preference for (110) texture, although this orientation is still dominant. The (lll), (311), and (331) orientations all increase significantly, with little difference seen between the films deposited on silicon nitride and those deposited on silicon dioxide. Laser recrystallization of the polysilicon films deposited on silicon nitride decreases the amount of (110) texture significantly; the crystallites with (111) and (311) orientation increase especially, with the (111) being generally dominant. Similar trends were seen both on samples which were thermally annealed and those which were not, with the thermally annealed samples showing an especially marked decrease in the (110) texture but little further increase in the (331) texture. The polysilicon films deposited on silicon dioxide and not thermally annealed showed the same trends as those deposited on silicon nitride, with, perhaps, somewhat less pronounced changes; there was less decrease in the (110) texture and less increase in the (311) texture. This difference may be partially related to the different laser recrystallization conditions, which were optimized for each type of insulator. The thermally annealed sample which was laser recrystallized to produce a large-grained structure (13-3) also showed similar trends to those found on the silicon-nitride-containing sample. The other sample (13-1) was treated at a lower laser power so that the large-grained structure was not obtained; this film probably contained a mixture of large and small grains. The increase in the amount of (110) texture is probably related to some growth of the structure initially found in the as-deposited films. This mixed case is, of course, of less interest than the fully recrystallized films. A significant increase in (100) texture is seen in many cases, but in no cases does (100) become the dominant orientation. In general, the laser-recrystallized samples exhibit less-strong preferred orientation than do the as-deposited films. The influence of this mixed orientation on the oxidation rate and interface charges of structures containing laser-recrystallized polysilicon will be discussed in Sec. 5.4.
5.1.7
Grain Boundary Diffusion
A conclusion drawn in Sec. 5.1.5 from the dopant distribution experiment for the arsenic implanted sample was that at 10 W, when the dopant profile became essentially flat throughout the film, the melt depth has extended all the way to the dielectric interface. Considering the diffusivity of impurities in the melt
242
K. E LEE E T A L .
(about cm2/sec I5.101) and the duration of the melt (which should be approximately the dwell time of the laser, or about 0.3 msec), the diffusion length of the dopant is about 2 pm. Hence the dopants would be quite uniformly distributed within the molten portion of the polysilicon film. While this is observed for the high power levels (10 W and 11 W), it is seen that the impurity profile for arsenic at the intermediate power levels (8 W, 9 W) exhibits a tail that is inconsistent with the above considerations. The diffusivity associated with this tail is on the order of lo-' cm2/sec. For comparison, the solid phase diffusivity of impurities in single crystal silicon is about cm2/sec near the melting temperature [5.11]. Hence the diffusivity in the tail cannot be explained by either a simple melting or solid phase diffusion within the single crystal grains. Furthermore, 