fabrication and characterization of hydrogenated amorphous

Retrospective Theses and Dissertations Iowa State University Capstones, Theses andDissertations

1994

Fabrication and characterization of hydrogenatedamorphous silicon films, CVD diamond films, andtheir devicesHao JiaIowa State University

Follow this and additional works at: https://lib.dr.iastate.edu/rtd

Part of the Condensed Matter Physics Commons

This Dissertation is brought to you for free and open access by the Iowa State University Capstones, Theses and Dissertations at Iowa State UniversityDigital Repository. It has been accepted for inclusion in Retrospective Theses and Dissertations by an authorized administrator of Iowa State UniversityDigital Repository. For more information, please contact [email protected].

Recommended CitationJia, Hao, "Fabrication and characterization of hydrogenated amorphous silicon films, CVD diamond films, and their devices " (1994).Retrospective Theses and Dissertations. 10485.https://lib.dr.iastate.edu/rtd/10485

http://lib.dr.iastate.edu/?utm_source=lib.dr.iastate.edu%2Frtd%2F10485&utm_medium=PDF&utm_campaign=PDFCoverPages

http://lib.dr.iastate.edu/?utm_source=lib.dr.iastate.edu%2Frtd%2F10485&utm_medium=PDF&utm_campaign=PDFCoverPages

https://lib.dr.iastate.edu/rtd?utm_source=lib.dr.iastate.edu%2Frtd%2F10485&utm_medium=PDF&utm_campaign=PDFCoverPages

https://lib.dr.iastate.edu/theses?utm_source=lib.dr.iastate.edu%2Frtd%2F10485&utm_medium=PDF&utm_campaign=PDFCoverPages

https://lib.dr.iastate.edu/theses?utm_source=lib.dr.iastate.edu%2Frtd%2F10485&utm_medium=PDF&utm_campaign=PDFCoverPages

https://lib.dr.iastate.edu/rtd?utm_source=lib.dr.iastate.edu%2Frtd%2F10485&utm_medium=PDF&utm_campaign=PDFCoverPages

http://network.bepress.com/hgg/discipline/197?utm_source=lib.dr.iastate.edu%2Frtd%2F10485&utm_medium=PDF&utm_campaign=PDFCoverPages

https://lib.dr.iastate.edu/rtd/10485?utm_source=lib.dr.iastate.edu%2Frtd%2F10485&utm_medium=PDF&utm_campaign=PDFCoverPages

mailto:[email protected]

INFORMATION TO USERS

This manuscript has been reproduced from the microfilm master. UMI

thesis and dissertation copies are in typewriter face, while others may

be from any type of computer printer.

The quality of this reproduction is dependent upon the quality of the

copy submitted. Broken or indistinct print, colored or poor quality illustrations and photographs, print bleedthrough, substandard margins, and improper alignment can adversely affect reproduction.

In the unlikely event that the author did not send UMI a complete

manuscript and there are missing pages, these will be noted. Also, if

unauthorized copyright material had to be removed, a note will indicate the deletion.

Oversize materials (e.g., maps, drawings, charts) are reproduced by

sectioning the original, beginning at the upper left-hand comer and

continuing from left to right in equal sections with small overlaps. Each

original is also photographed in one exposure and is included in reduced form at the back of the book.

xerographically in this copy. Higher quality 6" x 9" black and white

photographic prints are available for any photographs or illustrations

appearing in this copy for an additional charge. Contact UMI directly to order.

Jauixi

i \ I * 0 A Beii & Howe!! Information Company

300 f^Qrth 266b Rosd. .Ar^n .Arbor. Mi 43105-1345 USA

313/761-4700 800/521-0600

Order Number 9503570

Fabrication and characterization of hydrogenated amorphous silicon films, CVD diamond films, and their devices

Jia. Hao. Ph.D.

Iowa State University, 1994

1\ fl I X_/ JL T JL JL.

300 N. Zeeb Rd. Ann Arbor, IvII 48106

Fabrication and characterization of hydrogenated amorphous

silicon films, CVD diamond films, and their devices

by

Hao Jia

A Dissertation Subinitted to the

Graduate Faculty in Partial Fulfillment of the

Requirements for the Degree of

DOCTOR OF PHILOSOPHY

Department: Physics and Astronomy Major: Condensed Matter Physics

Approved:

±r n,arge. or lYiajor worK

For the Major Department

For the Graduate College

Iowa State University rviucS , j-Owca

1994

Signature was redacted for privacy.



ii

TABLE OF CONTENTS

GENERAL INTRODUCTION 1

PART ONE. FABRICATION AND STUDIES OF a-Si:H 2

I. INTRODUCTION 3

A. Historical Background 3

B. Basic Properties of a-Si:H and related materials 4

C. Hydrogenated a-Si 8

TN c-*- — TT 3 —. 1/ uj. J. u. ciiiiu vui. v.;^ cii xu

E. Scope of This Work 16

II. SAMPLE PREPARATION 18

III. SAMPLE CHARACTERIZATION 24

A. Thickness Measurements 24

B. Tauc Gap Measurements 24

C. Infrared Absorption 31

D. Electron Spin Resonance Measurements 37

E. Photo and Dark Conductivity Measurements 3 9

?. Secondary Ion Mass Spectrometry 3 9

G. Small Angle X-ray Scattering 42

T\7 -p-rrcTTTrpc Avn nx coTTccT/^v: crn

A. Electronic Stability 50

B. Microvoids and H Content 59

C. Anomalous Hydrogen Diffusion 65

REFERENCES 7 6

iii

PART TWO. FABRICATION AND CHARACTERIZATION OF CVD DIAMOND FILMS AND DEVICES 82

I. INTRODUCTION 83

A. Historical Background 83

B. CVD of Diamond Films 87

C. Growth Mechanism of CVD Diamond 94

D. Scope of This Work 97

II. SAMPLE FABRICATION 99

A. Deposition System 99

B. Intrinsic and Doped Sample Preparation 99

C. Light Emitting Diode (LED) Structure 103

III. SAMPLE CHARACTERIZATION 105

A. Thickness Measurements 105

B. Scanning Electron Microscopy 105

C. X-ray Diffraction Measurements 108

D. Electron Spin Resonance (ESR) Measurements 110

E. Electroluminescence Measurements 113

T7 TV \Tr> C"O OV*" r» x V » K i l o u x j x o L / - L O V - . V J O O _ l - L O

A. Electroluminescence 118

B. The Native Defects of CVD Diamond Films 121

REFERENCES 13 3

GENERAL CONCLUSIONS 136

1

GENERAL INTRODUCTION

Semiconductor materials of hydrogenated amorphous silicon

(a-Si:H) and diamond films were fabricated by radio frequency

sputtering technique and chemical vapor deposition technique

respectively. As the nature of the two materials are not

closely related to each other, this thesis is arranged into

two parts.

The first part of this thesis described the studies of a-

Si:H films by infrared absorption, small angle x-ray

scattering, electron spin resonance, and secondary ion mass

spectrometry. The optoelectronic properties of a-Si:H films

were improved dramatically by annealing at 250-330°C; strong

correlation between hydrogen content and microvoid was found

in a-Si:H film; and contradictory evidences to the widely

invoked "multiple trapping" mode], for hydrogen diffusion in a-

Si:H were also found for the first time and discussed in this

work •

As the second part of this work, thin diamond films and

their related devices were fabricated and studied. The deep

ciodss V7s2rs siicccssfviilv

fabricated. A new defect center in the CVD diamond was also

identified for the first time through the hyperfine splitting

in the electron spin resonance spectra.

2

PART ONE.

FABRICATION AND STUDIES OF a-Si:H

3

I. INTRODUCTION

Unlike crystalline silicon (c-Si) which is well

understood and highly developed, amorphous silicon (a-Si) was

believed to be undopable and hence useless, until Spear and

LeComber^ successfully doped it both n- and p-type in 1975.

Since then, a-Si has been intensively studied and become

increasingly important in many advanced technologies such as

solar cells, thin film transistors, image scanners,

electrophotography, optical recording and gas sensors.

The early work on a-Si was carried out by Chittick,

Alexander, and Sterling^ in 1969. Their films were deposited

by glow discharge of silane (SiH4). However they failed to

produce either n- or p-type doped film due to the large number

of dangling bond states which effectively pinned the Fermi

>~ 7 4— "1 /J T T f •«-. J— . J — — ^ -A.V « a. nCiO i, 1W JL^/'« UliO. L- J_)CWX&

and his collaborators^ discovered the role of hydrogen in

hydrogenated amorphous Germanium (a-Ge:H). In 1976, Paul et

al^ also demonstrated that in rf sputtered a-Si:H hydrogen

passivates almost all of the unsaturated dangling bonds and

thus dramatically lowers the mid-gap density of states. The

Fermi energy can then be moved in either direction by doping.

4

B. Basic Properties of a-SirH and related materials

The periodic arrangement of atoms in crystalline

materials lead to Bloch's theorem and the relationship between

the allowed energies of electrons and their wave vectors.

Crystalline semiconductors are characterized by gaps of

forbidden energies separating the allowed energy bands. In

amorphous semiconductors, which are highly distorted systems,

the long range periodic order is erased. However, short range

order (SRO) still exist, as demonstrated by, e.g., Temkin and

coworkers^ and shown in Figure 1.1. The radial distribution

function (RDF) J(r) =47rr2p(r) is defined as the average number

of atoms in a thin spherical shell of radius r and thickness

Ar with the center of the shell on an arbitrary atom. p(r) is

the atomic density. By comparing the RDF with crystalline

germanium (c-Ge) in Figure 1.1, it is clear that the nearest

neighbor atom distance in amorphous germanium (a-Ge) is as

well defined as in c-Ge. However, the peaks become

increasingly sm.eared as the distance r increases since there

are no long range atomic correlation. The RDF also yields the

bond-bond angle which is given by

where r^ and X2 are the distances of nearest and next nearest

neighbor atoms respectively.

'•d p.

J { i ) (ilomj A*')

w a 'Tl !/)

o i-ii

0 1 o (D

(U ;3 ii.

\r

iu I o (D

l-t) ()

!•( (0 t-h

crj

tr o

hj

>•• — •w o

J(rl (ilorns A"')

UH O o

:>

O

00

O

6

One of the most important consequences of the lack of

long range order in amorphous semiconductors is that the wave

vector k becomes meaningless and hence the Bloch theorem is no

longer valid. However, energy bands and band gaps still exist,

and a density of states N(E) can still be defined to describe

them. Due to the distortions of the network, the sharp

features in the density of states are smeared out (Figure 1.2)

by localized tail states generated from the bond strengths

variation. These tail states are critical in controlling

electronic transport, as almost all earners rapidly

thermalize to tail states at the band edges. The deep

localized states in the gap arise from coordination defects,

predominantly Si dangling bonds, which determine many

optoelectronic properties by trapping and recombination.

A dangling bond arises, for example, when a silicon atom

bonds to only three other silicon atoms. This may happen at

surfaces where another Si atom required for bonding is simply

missing. Or, it may occur in the bulk as a result of the

distorted Si network. In either case, this leaves a Si sp^

orbital unpaired to another sp^ orbital. Thus, a neutral

dangling bond results in an unoccupied near mid gap state

(Figure 1.2).

7

Conduction bond extended states

iij

>* CP

c LiJ

T3 dangling bonds

Valence band

Localized conduction band tail states

T3 dangling bonds

Localized valence band ^^tnil states

O \^AlOltUO^ OIVJIOO

Density of States, N(E)

Fig. 1.2 Davis-Mott model of the electronic density of

cKrM.T-ir^rr W.- Vm M A * A A O N-I' A k f A • s»/ WAA V2.

band, conduction band, tail states, and dangling

bond states near i?.id gap

8

C. Hydrogenated a-Si

Due to the large amount of dangling bonds (up to

lO^O/cm^) , the gap states largely mask the gap in a-Si thus

pinning the Fermi energy level at the center, making doping

impossible. However, the incorporation of hydrogen in a-Si, to

yield hydrogenated a-Si (a-Si:H) greatly reduces the density

of dangling bonds and sharpens the band tail edges (Figure

1.3), enable doping and dramatically improving the

optoelectronic properties.

The reduction of the density of dangling bonds is a

result of hydrogen passivating the dangling bonds, and can be

seen from the electron spin resonance (ESR) measurements.

Typically, spin density is about lO^^-lO^O cm"^ in a-Si and as

low as 10^^ cm~2 in device quality a-Si:H.

Addition of hydrogen also increases the energy gap of a-

Si and produces sharper band edges. The energy gap or Tauc

gap (Eg), which can be measured from optical absorption,

increases continuously from 1.4 ev to 2.0 ev with increasing H

content and is attributed to the silicon-hydrogen alloying

effect. The tail states are mainly formed fromi the weak Si-Si

bonds attributed to the variations of the local bond lengths

and angles. The addition of hydrogen breaks these weak Si-Si

bonds to form stronger relaxed Si-H bonds, thereby reducing

the width of the band tails.

9

22^

Amorphous Si containing no hydrogen

U

O O

containing hvdros,en

Valencc I j Conduction

band Forbidden band

Fig. 1.3 Effect of hydrogen addition in amorphous silicon

10

D. Electronic Stability and Hydrogen Motion

(i) The Staebler-Wronski Effect

One of the major applications of a-Si:K is for solar

cells, which are inexpensive in comparison with c-Si cells and

can be fabricated on a large scale. The problem limiting the

application of a-Si:H solar cells is the so called Staebler-

Wronski effect (SWE). In 1977, Staebler and Wronski^/^

discovered a reversible light induced degradation process in

THo p^otcconciuctivity cc^ciu.ctl^ V^ "ty C'q

of glow discharge (GD) deposited a-Si:H were reduced

considerably upon illumination with intense light. While

was reduced by factor of eight, decreased by four orders of

magnitude (Figure 1.4). Upon annealing at 150°C for four

hours, both ap^ and recovered their original values and the

entire process could be repeated.

Xt IS generally agreed that the illumination itself does

not create deflects in tne micigap regionj ra*ch3r some eiec"cron~

V> I VN T O O T 1 -N 4- <-s "v— -> * 2. AO

jTSCoiulDins.ion/ s.n,v5. "tlis irsj.Ga.SGc3. Gnsir^y cir£a.u.GS tliG ciGfGC't.s

which degrade the optical and electronic properties. Stutzir.ann

et al.^*^ carried out a detailed study on the SWE and reached

the following conclusions: (a) The SVJE effect is basically an

intrinsic bulk effect, independent of the concentration of

major impurities like oxygen and nitrogen, (b) The generated

11

in-3

10 -4

10 -5

I E o

10 -5

^ Illumination

> 10 — I s I

O 10 -s

1 A-s[_

e \

r — 1 0 •lU • L

I

ir_ 0 50 100 150 200 250

Time (min)

Fig. 1.4 Staebler-Wronski effect

12

defects have an ESR signature identical to that of dangling

bonds, (c) The generation rate of the metastable defects does

not depend on the incident photon energy in the range between

1.2 and 2.1 eV. This suggests that the defects are created

after the thermalization of the optically excited carriers

into the deep band tails. However, considerable controversy

still remains as to the nature of these defects and their

generation mechanism.

(i.X) Hvur'ocrcii mouxon

The incorporation of hydrogen in a-Si leads to many

beneficial effects as mentioned earlier. Therefore, hydrogen

motion in a-Si:H could be crucial in determining much of the

electronic properties of the material. The weak Si-Si bonds

broken by the released energy from the same nonradiative

recombination events as in the SWE might require a hop of an H

atom from a neighboring site to stabilize this newly formed

defect and this may constitute the basic diffusion step for

hydrogen diffusion.Kence the SWE may be viewed as "light

enhanced hydrogen diffusion". This potential role played by

hydrogen in the origin of the metastability has provided an

impetus to study the hydrogen diffusion in these materials.

The ability of hydrogen to move into, out of, and within

a-Si:H has both beneficial and undesired results. The low

defect density is a beneficial result from H termination of

13

dangling bonds. Hydrogen is, however, also apparently

responsible for the instability of a-Si:H at elevated

temperature due to the increasing rate of creation of dangling

bond by hydrogen evolution.

We initially assume that H motion obeys Pick's second law

of diffusion

dC = D—Tr 1,2

dt dx 2

where C(x,t) is the hydrogen concentration and D is the

diffusion constant.

The closest form of boundary conditions for the

multilayer structure is that of semi-infinite slabs in contact

at the interfaces (Figure 1.5). The boundary conditions for

such a structure are: C(x,0)=Co for x<0, C(x,0)=0 for x>0, and

/I / 0- \ A nTU 4 i_T a_ - v.. -1 i- J — V/. X no OWJLC1U.JLW11 UllC «^XJLJL«-1AXWA1 CTMU-dUXWll UW

these boundary conditions has the form of a complementary

error function^^

C(x,t) = C o / 2 erfc[x/(4Dt)1 / 2 ] 1 . 3

where CQ IS the deuterium concentration at boundary (x=G); t

is the annealing time. The complementary error function is

defined as

14

C

I X

Fig. 1.5 At t=0, C=Cq to the left of x=0. After diffusion

has proceeded for a period of time (as shown by

solid line), a finite amount of D has diffused into

15

c» _ 2 erfc{x) = \ e ^ dy 1-4

;c

For disperive diffusion where D is time dependent, Dt is

replaced by

r

0(i) = jD(T)dT 1.5

therefor, the solution can be expressed as

C(x,t) = Co/2 erfc[x/(40)1/2] 1.6

despite the time dependence of the diffusion constant, the

profile generally is still an error function, as verified by

SIMS.

The time dependent diffusion constant of hydrogen in a-

Si:H was found to obey a power-law given by^-"!^

" V - /

where Dqo is a prefactor, co is the attempt frequency and a is

the temperature dependent dispersion parameter which is

generally greater than zero since diffusion normally slows

down as time increases. The "hydrogen glass" model, proposed

16

by the Xerox group,suggested that a depends on the

temperature T as

a = l - P = l - T / T q 1 . 8

where IcTq is the width (decay constant) of the exponential

distribution of multiple trapping (MT) sites.

E. Scope of This Work

In this work we studied the electronic and optical

properties of a-Si:H. The samples were prepared by rf

sputtering at varying rf power and H content. Thus the

microstructure of the samples could be controlled and the

relation between hydrogen content, microstructure, electronic

and optical properties could be studied. The major techniques

used in this study were infrared (IR) absorption, secondary

^ ^ ^ ^ -.IT ^ ^ I-*-.---i.v AA WA. jr f j\. l\Q. y O wcx u. a

(SAXS); and electron spin resonance (ESR). The IR yielded

information on the total hydrogen content and the nature of

hydrogen bonding; the SAXS provided insight on the microvoid

structure and content; the ESR yielded the dangling bond

density, and deuterium SIMS profiles of a-Si:H/a-Si: (H,D)

^ i ^ - v — / - > • > - » V > ^ f r O v - / — o V a. ^ =.LA wjL. AiiQ Wii L4.X. Ci i jL J. X wi i •

a-Si:H and a-Si:H/(H,D) multilayers of varying thickness

were prepared at different rf powers ranging from SOW to 600W.

17

The hydrogen partial pressure and the sputtering power

determined the total hydrogen content which varied from 2 to

3 0 at. %. The target to substrate distance for most of the

samples was kept at 1.25" or 2". The substrate temperature was

controlled from room temperature to 250°C. The annealing

temperatures were 150-430°C, and all annealing processes were

carried out in evacuated pyrex tubes.

18

II. SAMPLE PREPARATION

As mentioned above, all of the samples used in this work

were prepared by reactive rf sputtering at a frequency of

13.56 MHz. A schematic diagram of the deposition system is

shown in Figure 2.1. Sputtering is one of the most common

techniques for fabrication of various thin films, rf sputter

deposition of a-Si:H is accomplished by bombarding a Si target

with energetic (up to several kev) ions of an inert gas (Argon

in this work). The inert gas is ionized and accelerated

towards the target by the self-biased electric field. Those

ions then strike the target surface and knock out atoms from

the target. These atoms then travel across the plasma and

deposit on the substrate.

The target used in this work was a 6" diameter 9 9.99%

pure polycrystalline silicon target, mounted on a water cooled

stainless steel backplate located at the top of the plasma.

Several types of substrates were used in this work for various

purposes. Double side polished single crystalline silicon

wafers were used for IR absorption and SIMS m.easurements,

Corning 7095 glass slides for UV-Vis-NIR absorption and ESR

measurements, and thin aluminum foils for SAXS and nuclear

magnetic resonance (NMR) measurements. The substrates were

placed on a heated plate directly underneath the target and

the distance between the target and the substrate holder

19

Generaiion

Matching

Target

Ti-Ba!l Subiimlnator

Ggs Mixing Tank Sputter Etch

Butterfly Valve Pedestal

Argon

- n Hydrogen i-CiO-" •

HDopant Gas

^2)— Nitrogen

npxp Gate Valve

Turbomolecuiar Pnmn " r"

Fore Pump

Fig. 2.1 Schematic diagram of the rf sputtering system used

for deposition of a-Si:H

20

(heater) could be easily adjusted from 1" to 2". The

background vacuum pressure prior to film deposition was pumped

down to about (3±l)xl0~'^ torr. During deposition the gate

valve connecting the plasma chamber to the turbo pump was

opened three turns to slow down the pump rate. Gases were

introduced at a preset partial pressure; the flow rates were

controlled by micrometer valves. The rotatable pedestal in the

chamber enabled cleaning of the target by presputtering.

After the desired gas flow rates and partial pressures

are achieved,, rf power is applied to the target while the

pedestal, on which the substrates were held, is grounded. The

strong rf field then ionizes the gas atoms and forms a plasma.

When the target is at a negative potential the positive ions

in the plasma are attracted to the target. Because the

potential of the target alternates at a very high frequency

(13.56 MHz), the positive ions are not able to reduce the

negative porenriai of the target to zero before the rf power

2.2. to 2. positivs VJlisr* is

s. positive pot^3.sttZTwCts tlis siscwircns in

plasma. Since electrons are much lighter and hence faster than

ions, they are much more effective in reducing the positive

bias of the target, and the average (DC) potential of the

target is reduced from zero to some negative value. Figure 2.2

shows the rf voltage supplied by the power source, V-, and the

potential at the target, Vj^. This naturally occurring self-

21

Time (arb. units)

Fig. 2.2 Voltage vs time characterization for source voltage

V^, and the target voltage V-^, in a rf sputtering

system (From ref. 19)

22

bias of the target causes the positive ions in the plasma to

bombard the target surface and knock out atoms of the target

which are then deposited on the substrate.

Keating the substrate during the deposition process

generally has a beneficial effect. It can enhances the lateral

mobility of the surface atoms thereby reducing an incipient

columnar morphology and microvoid content of the films. At

-250Oc the di- and tri-hydride configuration on the growing

surface tend to break and form molecular H2 which is removed

from the film. Therefore, heating the substrate will reduce

the density of those bonds which normally create voids or

surface-like defects. However, when heating above -400°C will

break the mono-hydrogen bonds and increase the density of

dangling bonds.

Multilayer films of a-Si:H were also prepared by

sandwiching an a-Si:(H,D) layer between two a-Si:H layers.

Figure 2.3 shows the structure of a multilayer film, and an

ideal as-deposited SIMS profile for such a structure. By

sandwiching an a-Si:(H,D) layer in a-Si:H, we can monitor

hydrogen-deuterium interdiffusion through SIMS measurements.

As the major difference between the diffusion of deuterium and

hydrogen is presumably due to their different masses, their

diffusion mechanisms are expected to be identical.

23

a-Si:(H,D)

a-Si:H j — ^1

Si substrato

o q) in \ yj 4-1 C 3 o r i

cc .C cn

•r-i yj

cn s r-t cn

>SI:H a-Sl:{H.D) a-Sl:H

Depth

Fig. 2.3 The multylayer structure and an ideal SIMS profile

of a-Si:H film

24

III. SAMPLE CHARACTERIZATION

Various techniques were used to characterize the

structure, optical, and electronic properties of a-Si:K films.

We used a Dektak profilometer to determine the thickness of

the films, optical absorption to determine the optical band

gap (or Tauc gap), IR absorption to determine the bonding

structure, SIMS for diffusion measurements, ESR for dangling

bond spin count, and SAXS for microvoid characterization.

A. Thickness Measurements

A Sloan Dektak profilometer was used to measure the

sample thickness. A stylus mechanically determines the

thickness of the films with an accuracy of ± 100 nm. During

the deposition, one piece of Corning 7059 glass substrate was

partially masked to create a step for the thickness

measurement. Figure 3.1 gives a schematic illustration of this

technique.

B. Tauc Gap Measurements

The optical band gap, or Tauc gap Eq, was determined by

using Cary 14A spectrophotometer. A schematic of the dual beam

instrument is shown in Figure 3.2. It can measure the optical

density (OD) from 300 to 2000 nm. The incident beam of

intensity IQ passes through the reference compartment while

fvietai Mqsk Z

- J. J L / /

t TJL F

Glass Substrate

Rofn i \j . . ro qni litprlriri i^iv^ v-yi^umciuivj i ^

1

i

y

Dektak Stylus

= -}Film Thickness

i >

U WZZZZZlr

•a-Si :H

Gloss Substrate

After Sputtering

Schematic representation of thickness measuremen

26

the transmitted light of intensity I exits through the sample

compartment. The width of the entrance slits of these

compartments is automatically adjusted to balance the

intensities of the two exiting beams. The optical density is

defined as

OD = logio(Io/I) 3.1

When light passes through a uniform homogeneous film of

thickness d, its intensity will be attenuated

I = Iq exp(-ad) 3.2

Since the film is deposited on Corning 7059 glass for

mechanical support, the reflection at the interfaces of

air/film and film/glass must be taken into account. The

-r> O rp T I T VN ^ /-> V^v-N-V- T" T" O

as 19

r \

- ad) cos{A mi.-d I ?.) ^ ^

where n a=l, n f , and n g = 1.53 are the indices of refraction of

air, film, and Corning 7059 glass, respectively. For a-Si:H

film, the accepted value of nf is -3.50, but it varies with H

content and wavelength. In Ecj. (3.3) , TQ is the transmission

27

Tung«t»n Lomo

u I nmn 1^9 1 • i! 1

9--

Photo Tul>«

A Ht Lamp

B,C>i Lef»«s D Bntronce Slit

r mkrofs H intwrmdJot# Slit

L Exit Slit

P 60 cpft Rototing

^tnmr

Fig. 3.2 Optical system of Gary 14A spectrophotometer

28

coefficient of Corning 7059 glass, which has a value of 0.97

in the wavelength range of interest.

Figure 3.3 shows a typical absorption spectrum. The

interference fringes in the wavelength region where a is small

are due to the cos(Aunfd/?L) term in the denominator of Eq.

(3.3). However, when a become larger at short wavelength, both

terms in the denominator of Eq. (3.3) become so small that we

can neglect them, and a is then given by

1 , . 10"°° a = —mi 1 3. a

The optical gap Eg can then be determined by using the

relation suggested by Tauc et. al.^C

{ccnjhco)^'' = B{jico — E^) 3 . £

where B is a constant which depends on the density of band

tail states. A Tauc plot is shown in Figure 3.4. The optical

gap Eg is obtained by extrapolating the region which obeys Eq.

(3.5) to a = 0; the intercept of the line with the fico axis

yields Eg.

In amorphous materials the optical properties are

characterized by three different absorption regions: At high

incident photon energies, the interband transitions are

similar to those which occur in single crystals. Below this

29

/ uu • , o ( 0\J

r~i r\

ouu /—\ R- Y—V gou yuu

Wavelength (nm)

3 A typical Gary 14 output, optical density vs.

wavelength

30

200 n—i—i—i—r

y'

a o

150 y -v

y

0)

cv

*

100

1 .• i .d I . ( i.O

Photon Energy (eV)

Fig. 3.4 Typical Tauc plot for optical gap

31

region, absorption is due to transitions between the band tail

states which are absent from crystalline materials; these were

first observed by Urbach^l in ionic crystals and are known as

"Urbach edge transitions". Below the Urbach edge, absorption

is due to transitions involving the defect states in the gap.

C. Infrared Absorption

Infrared (IR) absorption is an important technique that

is widely used in characterizing a-Si:H. It can provide

information on the total H content, its bonding

configurations, and oxygen or nitrogen contamination in the

films. The IR measurements were made on samples deposited on

double-side polished single crystal Si wafers by using a

single beam IBM model IR98 Fourier Transform infrared (FTIR)

spectrometer. The absorbance spectrum of a blank substrate was

subtracted from the absorbance spectrum of the substrate-film

sample to give the absorbance spectrum of the film alone.

A typical IR spectrum of a-Si:H is shown in Figure 3.5.

Three main Si-H vibration modes were first identified by

Brodsky and coworkers.22 These are (1) the Si-K stretching

mode at 2000 - 2100 cm~^, (2) the Si-H2 and Si-H3 scissors or

bond bending mode at 840 - 890 cm.~^, and (3) the Si-H wagging

mode at -640 cm~^. The corresponding bonding configurations

are depicted in Figure 3.6. The Si-bonded H content and Si-H2

32

configuration density can be calculated from the 64 0 cm~^

wagging mode by using the relation:

Nff = Aj a{cL>) Id) da ^ ^

where A is a proportionality constant which depends on the

oscillator strength and a(o) is the absorption coefficient at

frequency co. The values of A for different vibration modes

were determined by Shanks et al.23 Cardona ' / 25

given in Table 3.1.

The 640 cm~l band, which shows no discernible structure,

is attributed to the wagging mode of Si-H bonds in any bonding

configuration (see Fig. 3.6). The prefactor A of this mode has

been found to be nearly independent of both deposition

conditions and total H content calculated from this

band is thus considered to be reliable. From Eq. (3.6) and the

experimentally determined A value (Table 3.1), the Si-bonded H

content can be expressed as

H.X.(at, %) = 1.125 As4o/d 3.7

where Ac/n is the area (in cm~^) under the absorption peak of

640 cm~^ wagging mode and d is the film thickness in microns.

Similarly, the concentration of bonds corresponding to the 840

33

2SEB vsyes'jxsess ck-:

Fig. 3.5 Typical IR absorption spectruxn of a-Si;H

34

Table 3.1 The various vibration modes of hydrogen in a-Si:H,

their bonding configurations, and the corresponding

proportionality constant A (from Ref. 23; see Eq.

(3.7))

Uavenumber 540 8A0-890 2000 -2100 2100 ( cn

Mode Wag Sc i sso rs S t re t ch S t re t ch S t re t ch

Bond ing S i -H S i -H2 S i -H S i -H S i -H2

S i -H2 S i -H2 ( I so la ted ) (C lus te r ) SI -Kt Si -H3 (m ic rovo ids ) ^ (m ic rovo ids ) ^

A ( cm-2 ) 1 .6x10^9 2x10^0 2 .2x101^ l .Tx lO^O 9 .1x l0 l9

^Compressed m ic rovo ids

' ^H i c rovo ids v i t h rad ius ^ 2A

See t ex t f o r de ta i l s .

35

•Si OH

t Vxvx?

i 9 e « a ' SYM SYM ASYM

• ASYM, u cu cj 3

1 2 -3

2000cm"' 2090cm*' 2i20cm'i

9r> q''9> a

A f T SYM ASYM SYM

I r J

(11° 3

390cm-' SSOcm-i

'r>

v- r i / \ / i \ n

• ® » cj}^ , yi w

O < I I I J

w cj"

64 0cm"' 590cm''

Fig. 3.6 Illustration of vibration modes of stretching

(top), bending (middle), wagging and rocking

(bottom)

35

cm~l Si-H2 and Si-H3 scissors mode vibration can be expressed

as

n(j(at. %) — 10.44 agaq/'*^ 3.s

The stretch mode at 2000 cm~^ is attributed to isolated

Si-H bonds in the bulk. The peak, at 2100 cm~^ has been

attributed to silicon associated with di- and tri-H bonds or

to mono Si-H at internal surface of larger microvoids.^2

However,, it is not possible to distinguish between silicon

bonded to two or three hydrogen atoms from either the scissors

or stretch peak. It has been pointed out that the stretch

frequency of mono Si-H bond shifts from 2000 to 2100 cm~^ as

the radius of the microvoid around the hydrogen increases to -

2 A.25-27

Presence of oxygen at levels above -0.5 at. % in the

films can also oe easily aetected from rne IR specrra. Tne 900

c~;—— p3=k is related to bulk Si~0 bonds, while ths 1100 cm~~

psH}c is stinri-biitsd. to tlis Si~0 tond. intsirnBl su^rfscss- Tins

later is often observed in films exhibiting columnar

morphology, in which case it grows with time upon exposure to

air.

37

D. Electron Spin Resonance Measurements

Electron spin resonance (ESR) measurements were used to

measure the density of unpaired dangling bonds in a-Si:H

films. The measurements were performed on a Bruker ER2 2 0 DSR

X-band ESR spectrometer. The samples were deposited on Corning

7059 glass cut into pieces of half inch by three millimeters.

The spin density and the g factor were obtained by comparing

the ESR spectrum with a reference DPPH of known spin density

and g factor.

An electron will interact with a magnetic field due to

its magnetic dipole moment The interaction will cause an

energy level splitting or Zeeman shift

the dipole moment is related to the electron spin S by

where g is the g-factor and P = eTi/2mQ is the Bohr magneton.

Since the eigenvalues of S along the magnetic field direction

E = -}l-B 3 . 9

\l = -gps//z 3 .10

^>*0 / O +*"H^ *7 O CiTTi Y-V C"V>1 — — • - / — r — - - — w — ' ^ f r

ae = gpb 3 .11

38

Resonant transitions between these levels can now occur if an

electromagnetic field of frequency co given by

no)= gpB 3.12

is applied normal to B. For B = 3340 Gauss and g = 2, the

resonant frequency is o = 9.3 GHz (X-band).

The ESR measurements were performed by inserting the

sample in a microwave cavity placed in a magnetic field

(Figure 3.7). The microwave frequency was set at 3.3 GKz and

the magnetic field was swept from 3330 to 3350 Gauss. A small

modulation field of 4 Gauss was applied parallel to B and

referenced a phase sensitive lockin detector which then

yielded the resonance signal detected by a microwave bridge

attached to the cavity. This method yields the first

derivative of the absorption spectrum, as shown in Figure 3.8.

The total number of spins in the sample can be calculated

by comparing with a reference's

where the superscripts r and s denote the reference and sample

respectively. In Eq. (3.13), Ah^Q is the derivative peak-to-

peak width of the resonance spectrum, Y is its amplitude,

is the amplitude of the modulation field, is the microwave

39

power, A is the line shape factor, G is the lockin gain, and

Np is the number of scans.^8

E. Photo and Dark Conductivity Measurements

Conductivity measurements were performed by using a home

made two-probe contact setup which is similar to the four-

probe method^^ and had been calibrated with a known sample.

The photoconductivity was measured with a heat-filtered

tungsten-halogen lamp which provided an intensity of 200

mW/cm^ on the sample surface. The photo- to dark- conductivity

ratio was then calculated and compared among samples to

characterize the optical performance.

F. Secondary Ion Mass Spectrometry

Information on long range hydrogen (or deuterium) motion

can be monitored by using secondary ion mass spectrometry

(SIMS).^'^ This technique is often used to profile the near

surface region of a solid as a function of depth. It uses a

primary ion beam to bombard the sample surface and sputter off

atoms and ions from the sample. Secondary ions can then be

identified by a mass spectrometer.

A schematic of the Perkin Elmer SIMS model PHI 5300 SIMS

used in this work is shown in Figure 3.9. A primary beam of

positive ions (Cs''") with diameter -50 S-im and energies -5 keV

wss s 2.^ 30^ ^ s

40

Magnet

Microwave

Generator

Sample

Fig 3.7 Schematic of ESR system

41

D <

b>

'A ul

3.0

2.0 -

1.0 -

0.0

-1.0

-zd

3.29 3.31 3.33 3.35 ilhousanda!

Magnetic Field iKGau&s)

Fig 3.8 Typical ESR spectrum of a-Si:H

42

beam was used to raster a 500x500 jim^ area on the sample

surface. The secondary ions from the sample surface were then

analyzed by the SIMS analyzer which consisted of an energy

filter, a quadrupole mass filter, and a signal amplifier.

Figure 3.10 shows a typical deuterium depth profile of an

a-Si:H/a-Si:(H,D) multilayer film. The x axis corresponds to

the sputtering time, which is proportional to the depth. The y

axis is the secondary ion current intensity in units of

counts/second. The sharp drop in the concentration is used to

define the interface. It also determines the depth resolution

of the profiles, which is typically - 150 a at the interface

of a - 1 um thick film.

Besides hydrogen and deuterium, other atomic species such

as oxygen, carbon, silicon, etc., can also be monitored. This

is achieved by simultaneously monitoring their corresponding

m/e values. Since the quadrupole mass analyzer does not employ

magnets, this peak switching process can be done rapidly

without hysteresis effects.

G= Small Angle X-ray Scattering

Microvoids in a-Si:K and its alloys have been the subject

of extensive research in recent years. A number of methods^

have been used to infer the existence of microvoids and

attempt to quantify them, but small-angle X-ray scattering

(SAXS) remains the most direct, as it can provide information

43

ION GUN

\

\ SPUTTiiR

"^^angli: SIMS

ANALYZER PRIMiARY

IONS SECONDARY

IONS ELECTRON

MULTIPLIER

VACUUM

T C T) T • ".*

Fig 3.9 Schematic of SIMS system

44

as deuos i t ed

i i i 1 i i i i i i i i i i i i i 1 i 1 i 0 .0 ' 2 ,8 5 .6 8 ,4 1 1 . 2 14 ,8 16 .8 19 -6 ' ^2 .4 25 .2 28 .0

SPUTTER TiliE, lilH.

i i 1 i 1 i i i i i i i i i i i i i i i

0 Ui to

M

t

t-" d 0 u

y

0 0 J

723 hrs @300OC

"A*--V

i i . i i I I ! I

e .9 i f i 6 .n 8 14 ,s i ? . s

SAUTTER TiliE, iliH. '•/a 0 0 i.'J t -J V I V

i I I 0 .ri n

1. 'v* I V i ; 0

Fig 3.10 Typical SIMS depth profile of a-Si:H/a-Si:(H,D)

45

on the microvoid size, shape and number density. Small-angle

electron scattering (SAES)'^° and SAXS'^^"'^'^ were applied

several years ago in studies of unhydrogenated a-Si and a-Ge,

where observations of density deficits were explained by the

presence of internal -5 nm sized voids. Subsequent studies of

a-Si:H using SAXS'^^"'^'^, SAES^^"^*^, and small-angle neutron

scattering (SANS)^^"^^ all demonstrated the existence of

microvoids in non-device-quality material.

The a-Si:H films studied in this work were deposited on a

10 thick iron-free aluminum foil, which was then cut into

eight strips after deposition and stacked to produce

sufficient thickness for the SAXS measurements. These

measurements were carried out with a Kratky Compact Small-

Angle System (Anton Paar, Graz, Austria) attached to a Rigaku

rotating anode (Cu) X-ray generator. Some measurements were

made with the sample normal tilted at an angle to the beam.

To estimiate the microvoid volume fraction, Vf, the

foliov;ing integral expression, which is valid for a simple

two-phase system and for the line collimation geometry, is

used^^

\hi{h)dh = 3.14

where 1(h) is the SAXS intensity. h=27c(20)/>. is the magnitude

of the scattering vector, and Ao is the difference in the

46

electron densities of the two phases. The quantity K depends

on various geometrical factors, the electron scattering cross-

section, the X-ray wavelength, and the absorption coefficient

of the filn. Because of the uncertainties in the geometrical

factors, the theoretical value of K was ignored in favor of

the determined quantity. To apply Eq. (3.14), first an

experimental value of K is obtained for an a-Si:H sample by

assuming that the SAXS and the density deficit determined by

the flotation method are caused entirely by microvoids. This K

is then used to predict Vf for all other samples.

Figure 3.11 displays the SAXS data (I*h vs. h) of samples

#1-5 as deposited. The radius of gyration, Rg, which

characterizes the size of the electron density fluctuations,

can be obtained from the linear behavior of Guinier plots^®

lQi) = I txv{-Rlh- n) 3.15

The departure from linearity indicates that the films contain

more than one size of microvoids.The strong scattering of

sample #1 indicates that irs Vf is larger than others. Sample

#2 had much stronger scattering that the intensity has been

scaled down by factor of 3 in Figure 3.11. Figure 3.12 shows

the tilting effect of sample #1 at 45° and 60°. The tilt-

dependence changes suggests that the microvoids in this film

are oriented, rod-like voids^'^. From previous SAXS studies on

47

it appears that its microstructure is columnar. The

other three samples (#3-#5), which contain -10 at. % Si-bonded

H and -0.5 vol. % microvoids, showed no anisotropy upon

tilting, suggesting either spherical microvoids or a random

orientation of non-spherical microvoids.

43

Fig 3.11 Plots of I*h vs h for the as-deposited a-Si:H

films. Note the intensity of sample #2 has been

scaled down by factor of 3

49

1.50

non-tilt in s

45° - t i l ted

0.<30 0

h (nm

•ig 3.12 0° and 45° tilted SAXS of sample #1, showing the

strong tilt angle dependence of Vf (see Eq. (3.15))

50

IV. RESULTS AND DISCUSSION

A. Electronic Stability

Fifteen samples were deposited on nominally unheated

substrates for the systematic study of electronic and optical

stability- The rf power Pj-f and H2 partial pressure Ph2 varied

from 200 to 600 W and 0.2 to 1.0 mT, respectively. The samples

were then characterized by optical absorption to determine

their optical gap Eg, IR absorption to determine the Si-bonded

K content and. its confiyuration, and. the photo— & dark.—

conductivity to characterize the optoelectronic performance.

Table 4.1 summarizes the properties of the samples.

generally increased with decreased with increasing Pj-f

(Figure 4-1). The optical gap Eg increased with the hydrogen

content. The photo- to dark-conductivity ratio for all

as-deposited samples was very poor, and the ESR measurements

confirmed the high initial defect density of -10^^/cm^- These

defects are recombination centers which decrease the

photoconductivity and pin the dark Fermi level near mid gap.

However, after annealing at 3 00°C for 9 hrs,

dramatically increased to -10^ (Figure 4.2). Further annealing

at 3 00°C showed slight additional improvement. But the crp>,/ajj

of some samples began to decrease following annealing at

51

Table 4.1 Characterization of the as-deposited samples

#

PH2

(m-c)

Power

(W)

Thick

(M-)

%

(ev)

CH

(%) (ncm)-l r

I r

1 200 1.15 1.78 22.2 2.1x10-10 146

2 1 300 0.81 1.69 21.7 3.0x10-10 • 5.8

3 1 400 0.85 1.67 20.3 9.4x10-10 5.5 j

4 1 500 1.00 1.55 11.7 4.0x10-9 3.2

5 1 600 0.95 1.64 10.7 2.1x10-8 1.5

6 0.5 200 1.10 1.62 14.2 6.5x10-9 1.3

7 0.5 300 1.10 1.58 11.2 1.2x10-8 1.2

8 0.5 400 1.10 1.55 7.1 !. 5.4x10-8 1 1 j t 1

j 9 j 0.5 j 500 j 0.94 1.47 V.5 j 9.3x10-8 i 1.1

10 j 0.5 j 600 1 1.00 1.53 6.56 3.6x10-7 L2

11 1 0.2 200 1.25 1.41 1 8.0 2.0x10-6 1.1 i 1

12 I 0.2 i

300 1.15 1 1.39 •i i

6.67 5.0x10-6 j 1-0 ! 1 > * ^ j ^ i i 1 ,

13 I 0.2 i 400 I 1.05 I 1.37 j 8.8 j 2.9x10-5 i 1.0 |

14 I 0.2 j 500 I 0.77 j 1.34 j 4.8 j 7.8xlO-S j 1.0 j

15 i 0.2 I 600 i 0.60 I 1.36 I 6.4 ! i ixlO-4 i 1.0 ! • • i ! ! } X. .L A. J. V,' , 1

52

20 n I i i i ' I i i i ' i I i I i I i > i i i I '

15 a o c o

10 d o w! o u

^ 5

ph2 = 0-5 mt

i 1 i 1 i i i \ \ i i t i i i i i i i i i i i

200 300 400 500 600

RF Power (W)

V-i o

25

20

1 I I I I I I > I j t 1 i I I 1 I 1 i I I i 1 i j I 1 1 ]

d 15 o o

^ id o -to >>

200 W

600 W

0 h i ^ i i i i i i i i i i i i i i i i i i i i i i i i i i i i ii t

0 • 0.2 0.4 0.6 0.8 1 1.2.

H2 Partial Pressure (mr)

Fig. 4.1 Correlation between deposition parameters and the

resulted hydrogen contents

53

200 W 200 w 200 W 300 W 300 \V 300 W

i mi \J.D mi vj.Z mi 1 mT 0.5 mX O.z mi

Fig. 4.2 Optical performance characterized by OQh/'Jd

dramaticallv chanced after annealincr

54

3 3 0OC for 10 hours.

Figure 4.2 also showed that higher Cjj and lower Pj-f

resulted in better CTph/c^d' which is understandable because

more hydrogen is available in the sample, and more dangling

bonds will be terminated by hydrogen during its motion,

reducing the defect states in the gap; higher P^-f normally

creates more defects due to the higher energy bombardment

during sample preparation.

The dramatic improvement of during the annealing

processes is believed to be related to hydrogen motion. As

hydrogen diffuses, it will change the bonding configuration of

the silicon network. Figure 4.3 shows IR spectra of sample #3

before and after annealing. The peaks at -2100 cm~l

represented the Si-H2 and Si-H3, or surface-like Si-H (eg. on

void surface), or 0-Si-H bonds had decreased after annealing.

Another beneficial result of hydrogen diffusion is that it

will compensate many in-gap defect states and change the

conducting mechanism from localized hopping among gap states

to extended state transport. Figure 4.4 shows the ESR of

sample #3 before and after annealing. The as-deposited

dangling bond density ng = 2.0x10-^ cm"^ ^as very high. After

annealing at 300°C for 9 hours, it decreased so much that the

ESR was undetectable. However the upper limit was estimated to

be at -1.0x10-'-° cm"^. In other words, the density decreased by

more than two orders of magnitude after annealing.

55

as deposited

3SB0 3S00 2500 !59S l o a a 5SB

\Vavenumber (cm')

Fig. 4.3 IR spectra showed the decreased intensity at -2010

cm"^ after annealing

56

5.0

4.0 -

10

0.0

<

H.O

A /

as deposit ns=2.0xl0l8 /cm3

^

\l . . . .

3.23 3.31 3.33

0.3

0.2 -I

0.0

-0.1

-0.2

I

-0.3 -

-0.4 -

-0.5 —1

-0.5 -

-0.7

3.35 3.37 3.33

- 'A jt, 1

"- \Af H r, / \

300 oc 9 hrs

ng < 1.0x10^6 /cm3

\j y \ u 1 / I

1 ' 1 1 ' ' ' , lU J I r 1

\ / \ f \ . A ! '•i .1 k.

im 1' 1.'

\ ,1

3.23 3.31 3.33 3.35 3.37

Magnetic Field (K Gauss) 3.33

Fig. 4.4 ESR raeasurements showed the dramatically decreased

dangling bonds after annealing

57

Unlike a-Si:H deposited by glow discharge (GD)

decomposition of silane, which normally has a strong SWE, the

samples prepared by rf sputtering and treated by annealing

displayed enhanced light soaking stability. The samples with

high Ch and low Pj-f showed a weak S'WE (see Figure 4.5 and

compare with Figure 1.4), while the samples with low Cjj and

high P>-f showed no SWE. It is not clear what aspect of the

deposition process and the annealing is responsible for the

superior stability against light-induced degradation in these

samples. Small angle x-ray scattering studies^S found that the

microvoid content increased with increasing hydrogen content

in rf sputtered a-Si:H films, and the SWE is generally

believed to result from light induced defects which reduce

both photo- and dark-conductivity. In that sense the enhanced

light soaking stability might be explained as the high initial

defect density indicated by which made the light

induced defect density relatively small. Generally, device

quality GD a-Si:H with similar high hydrogen content (-10 at.

%) show strong SWE and ~ 10^.

In conclusion, the electronic and optical properties of

rf sputtered s-Si:H sample are strongly improved after

annealing at 250 - 300°C; can be as high as lO'^, which

is close to GD a-Si:H. As expected, the SWE increased with

increasing cr^h/cJj. The sample with the highest showed

58

10 -3 I I I r I I "i—r

«

o

t3! c o o

10 -4

a 10 o

-5

^ 10"^

10 - 7

10

10

- 8

h 1 n -9 i \ j

- 1 0 L 1 1 1 ]

Illumination

j l

I

=3 —t q 1 1

~i

I I I I ! ' I ] n 9 4 jl

Time (iiours)

a

Fig. 4.5 Weaker Staebler-Wronski effect (coir.pare with Fig.

1 a •\

59

much weaker SWE, while low cTpj^/a,^ samples deposited at high rf

power and low H partial pressure showed no SWE at all.

B. Microvoids and K Content

Five a-Si:H were prepared for SAXS studies by rf

sputtering of a Si target onto a nominally unheated grounded

substrate located 1.5" below the target. The Ar and H2 partial

pressures were 10 and 0.5 mT respectively, and the rf power

varied from 200 to 600 W. The effective temperature of the

film during deposition was estimated to be -150°C.^5 For each

sample, three different substrates were used. The first was a

standard crystalline silicon wafer, which was used for IR

characterization, the second was Corning 7059 glass for

optical gap and photoconductivity characterizations, and the

third was a 10 jim thick iron-free aluminum foil for SAXS

measurements. Samples were characterized by IR and SAXS after

each annealing step, which was carried out in evacuated pyrex

tubes and lasted 6 hours.

The properties of the as-deposited samples are presented

in Table 4.2. As clearly seen, sample #1 and #2 had the higher

hydrogen content and larger microvoid volume fraction Vf;

in samples #3-#5 Cu was -10% and Vf was -0.5 vol.%. These

values are close to device-quality GD films. All samples

were about 2 um thick. After the first annealing step at

250 c cv of sstp.'oxs —!l <5sc2r05ssd, 22.% 2.3%

60

Table 4.2 As-deposited properties of the rf sputter-deposited

a-Si:H films, including H content Cjj and Si-H2 bond

density 0^2? SAXS determined microvoid content Vf,

and the mass density p determined by flotation

measurements

Sample # rf power

(W)

thicknes

s

(lim)

(at.%)

^H2

(at.%)

Vf

(vol.%)

P

(g/cm3)

1 200 2 . 4 5 2 0 . 8 3 . 4 1 . 7 * 2 . 1 8

2 3 0 0 2 . 4 0 1 9 . 0 -10* 2 . 0 4

3 4 0 0 2 . 7 0 1 1 . 2 - 0 0 . 5 4 2 . 2 2

4 500 2 . 0 0 1 0 . 2 - 0 0 . 5 1 2 . 2 5

5 600 1 . 9 5 9 . 5 - 0 0 . 4 7 2 . 2 4

* The actual void fraction is less, because the tilt-

dependence of the scattering intensity causes an overestimate

of Vf .

61

while Vf increased from 2 to 2.25 vol.%. The other samples

showed similar slight changes except sample #2 which exhibited

strong post-depositional oxidation after 75 days in air (as

deterinined by IR absorption) and was not annealed. Further

annealing from 250°C to 310°C produced basically no changes in

either Cjj or Vf (see Figure 4.6). Other samples prepared under

similar conditions and annealed in this temperature range

showed sharply reduced dangling bond densities and C7ph/C7(j

ratios of -10^.^^ After annealed at 350, 380, 400, and 430°C,

however, Cjj decreased significantly. Further annealing at

43 0°C for 3 0 more hours further decreased and sharply

increased Vf. The drop of hydrogen content when annealed at

above 330°C was consistent with others studies^^ that Si-H

bonds begin to break and the hydrogen evolves from the film

above 3 3 0°C.

If microvoids are simply approximate to be spherical,

then the SAXS also yields the void size distribution as

described in Eq. (3.15). Figure 4.7 shows the size

distributions of sample #l and #4 before and after annealing

process. They yield several features which should be

addressed: (i) They ail indicate a significant content of

relatively large voids, regardless of the overall morphology

of the films, (ii) While most of the voids in unannealed

sample #4 are -1 nm across, most of those of unannealed sample

#1 are 1 - 3 nm in diameter. This observation strongly

62

25

S o m p l e

20

o

c

o c

25

o 2 0

CJ •a o £

o

o

Z i

o o Scmp

o

150 200 250 300 350 400 ' .0 70

Annecl lempcrcture ( C) Time (h)

Fig. 4.6 The integrated SAXS intensity or the microvoids,

and the hydrogen content after each annealing step

63

Fig. 4.7

Sample 1

as —deposifed ^ 430°C X 35 h ^ 40

Sample 4

as —deposited 4-30 C X 35 h

10^

D i a m e t e r (n m )

Tne distribution of void sizes in sample #i and #4

before and after annealing

64

suggests that the high overall inicrovoid content of sample #1

is clearly due to the large size of microvoids rather than to

more numerous small voids, (iii) After annealing, the

distribution of void sizes are wider in both samples, but the

large-scale density fluctuations are clearly larger in sample

#1 than in #4.

Three major conclusions resulting from the measurements

described above are clear: (a) They are consistent with the

well-known correlation between Cm, 0^2r and Vf: The initial

Si-bonded H content of Sample #1 and #2 is high (-20 at. %),

and it indeed contains a significant dihydride concentration.

The tilt angle-dependence of the SAXS clearly indicates that

the microstructures are columnar-like. The SiH2, SiH3, and

microvoid density of all of the other samples, which contain

9-11 at. % Si-bonded H, is very low. Both SAXS intensity and

Cu for the noncolumnar samples (#3-#5) did not change during

prolonged annealing (beyond 6 hrs) at 430°C. This behavior is

believed to result from migration of m.ost of the Si-bonded H

to the internal surface of that film, recombination to

molecular H2, and its escape through the largely

interconnected voids. Most of the remaining hydrogen of that

film was then probably bonded to the void surface in isolated

sites which are not adjacent to a neighboring H atom, (b)

Annealing at 350°C and above sharply increased the integrated

SAXS intensity,, or increased the average void size, and

65

decreased the Si-bonded H content C^. This behavior was

apparently due to hydrogen migration from the Si network to

microvoid surface, recombination of adjacent H atoms at the

surfaces, and accumulation of molecular K2 in isolated voids.

The decrease in Si-bonded H content then increased the density

of the Si network. Its reconfiguration and probable high

pressure of the trapped molecular hydrogen then increased the

average void size. It should be noted that decreasing density

of molecular H2 through evolution may also results in the drop

of SAXS detected Vf, however, the magnitude of the changes in

the SAXS may exclude such a possibility, (c) Annealing at

430°C resulted in the formation of a sufficient volume

fraction of microcrystalline silicon (|j.c-Si) domains to be

detected by Raman and XRD. Although the amorphous phase was

still dominant after 12 hrs at 430°C, it was reduced to -50%

after 3 6 hrs at the same temperature, and the columnar

structure was almost eliminated. This behavior is not

surprising in view of the large-scale crystallization

processes occurring when annealing at such a temperature

range.

C. Anomalous Hydrogen Diffusion

The time dependent diffusion constant of hydrogen in a-

Si:H was found to obey a power-law given

66

D(t) = Doo(®t)-a 4.1

where a is the dispersion parameter and co and DQQ are

constants. In the multiple-trapping (MT) model which has been

frequently invoked to describe hydrogen motion in a-Si:Hl2,14

it is assumed that the density of hydrogen sites is

exponentially distributed in trapping energy Ef This

exponential distribution of trapping energies leads to

a = 1 - T/Tq 4.2

where T is the annealing temperature and RTq is the

distribution width of trapping energies In this

multiple trapping picture of hydrogen motion, the hydrogen

will be trapped into the deeper and deeper sites as it

migrates. The longer it diffuses, the deeper it will be

parameter a should always be greater than zero. Indeed, all

values of a reported to date were between 0 and 1. However,

the dependence of a on T varied widely among different

samples. In doped GD films, the results were consistent with

Eq. (4.2) but not sufficiently detailed to confirm In

undoped GD films deposited at high rf frequency a actually

strongly increased with T above 350°C.®2 In some rf sputtered

67

15.0 300° C

12.5

185 hrs 10.0

7.5

5.0 724 hrs. 50 hr^l 85 hrs.

2.5 30

330° C z)

<

CO 20

co 30 hrs> 72 hrs>\146 hrs. 293 hrs. _

25 363° C

20

-c r \

24 hrs> • i"v

48 hrs.

0 0.05 0.1 0.15 0.2 0.25

Diffusion D!st3.nc6

Fig. 4.8 Fitting of sample A between SIMS profiles (dots)

and complementary error function [Eq. (1.6)]

68

i i i i i i i i i i

3 <

CO cz 3 o o cd

00

- 1 I I I I I I 1 I

325° C

M

2, 4, 10. 20 hrs

343° C

~ \V\ 1.2, 2.4, 4.7 hrs

362° C

.A3.0 hrK7.2 hrs. 15.7 hrs. i i rs ! i i i

1.0 hrs;

0 0.05 0.1 0.15 0.2

uiTfusion Distance (ji)

0.25

Fig. 4.9 Fitting of sample B between SIMS profiles (dots)


69

384° C

1.2 hrs

20 -0.8 -0.4

8.8 hrs.-Z ) <

1.2 hrs. 4.7 hrs.

c

0 25 CD

1 20

405° C

C O

1.2, 4.8 hrs.

0.2 0.25 u.uo

Diffusion Distance (u)

Fig. 4.10 Fitting of sample B between SIMS profiles (dots)


70

samples used in this work, we found and reported^"^ a strong

deviation from the multiple trapping model, in which a was

negative i.e., the D(t) increases with t.

The films were deposited by rf sputtering at 550 W,

annealed, and monitored by IR and SIMS as described earlier.

Sample A was a 2.4 iim thick a-Si:H/a-Si: (H,D)/a-Si:H trilayer,

and each layer was deposited over 3 0 minutes- Sample B was a

1.65 jim thick a-Si: (H,D)/a-Si:H bilayer, grown at 30 minutes

per layer. The IR stretch band of both samples before and

after each annealing peaked at 2000 cm~^, indicating almost

exclusive bulklike Si-H bonding. Cj^, as determined from the

640 cm~l wagging mode vibration, generally decreased from -5

to -3 at. % over the annealing periods shown in Figures 4.8 -

4.10.

9(t) defined by Eq. (1.5) was obtained by fitting the

deuterium SIMS concentration profile c{x,t) to complimentary

error function as given by Eq. (1.6). As clearly seen from

Figures 4.8-4.10, the agreement between the SIMS profiles and

the complimentary error function was always reasonable, and

usually excellent.

Substituting Eq. (1.7) into Eq. (1.5), we get

o - Oq-C^ ^ 4.3

71

X 300°C, a=-0.47

V 330°C, a = -0.34 o 362°C, a = -0.04

r io

£ u

a:>

SAMPLE A

7 s

Annealing lime (sec)

Fig. 4.11 logiQ[0(t)] vs. log]_o[t] of sample A. Note that a

slope larger than 1 implies a < 0 [see Eq. (4.3)]

72

10 -10

- D 325°C, a = 0.31 -

x 343°C,a= 0.15 ~ + 36Z°C,a=-0.79 —

: o 384°C, a = -0.51 -

- V 405°C,a = 0.43 -

10"!'

10-'2

!q-'2 u_

i O -

SAMPLE B

i i i i i i i

!0"

"ieclip.g Time (sec)

Fig. 4.12 logi_o[0(t)] vs. logi_n[t] of sample B. Note that a

slope larger than 1 implies a < 0 [see Eq. (4.3)]

73

Therefore, the slope of logio[0(t)] vs log3^Q[t] is 1-a. Figure

4.11 shows such a plot for sample A annealed at 300, 330, and

362°C. Remarkably, the value of a is -0.47 at 300°C but

increases to -0 at 362°C. Figure 4.12 exhibits a similar plot

of sample B in the temperature range of 325 - 405°C. From 325

to 362°C a decreases from 0.31 to -0.79, but then increases to

0.43 at 405OC.

Both samples A and B exhibited (i) negative values of a

and (ii) an a which increased with T at high annealing

temperatures. All of these results are contradictory to the MT

model, and attributed to the complex relaxation processes in

which the distribution of H site energies is probably

modified. If the shallower band narrows during annealing at

300 < T < 350°C, and the deep band rises toward the transport

state (see Figure 4.13), then Dj^(t) should increase with time

and a should be negative. Such a behavior would be expected if

the shallow band is due to interstitial sites similar to those

of H in crystalline Silicon, since the activation energy of

atomic H diffusion in crystalline silicon is -0.5 eV.^^

Although such processes may entail a rise in the free energy

Fjj of the H atoms, the total free energy of the H and Si atoms

may decrease. Roorda et al.^® indeed found that Si amorphized

by ion implantation undergoes exothermic relaxation beginning

at temperatures as low as 110°C, v/hen heated as rapidly as

4Q°C/ir.in. The structural relaxation occurring during several

74

Transport level

>

LiT

U)

o c LU U) c •q. q. cc

h-

x

l^ciioiiy vj 5 vjllcto

Fig. 4.13 The suggested distribution of H trapping energies

in a-Si:H. Hpj is the chemical potential of H atoms.

See Ref. 64

75

hours at T > 300°C should therefore be considerable and may-

affect the shallow H sites and the position of the deep H band

in the aforementioned manner.

As mentioned earlier, the SAXS study of a-Si:H^2

indicated that microvoid content Vf steadily increases when

annealing at 250 < T < 430°C. As the deepest H-trapping sites

are believed to be mono Si-H bonds on internal microvoid

surfaces, we conclude that the strong increase of a with T is

due to an increasing microvoid content. The concentration of

hydrogen in shallow sites thus decreases as an increasing

number becomes deeply trapped on the microvoid surfaces.

76

references

1. W. E. Spear and P. G. LeComber, Solid State Commun. 17,

1193 (1975).

2. G. C. Chittick, J. H- Alexander, and H. F. Sterling, J.

Electrochem. Sco. 116, 77 (1969)

3. A. J. Lewis, G. A. N. Connell, W. Paul, J. R. Pawlik, and

R. J. Temkin in Tetrahedrallv Bonded Amorphous

Semiconductors. edited by M. H. Broidsky, S. Kirkpatrick,

and D. weaire, AIP Conference Proceedings, 20, 27 (1976).

4. W. Paul, A. J. Lewis, G. A. N. Connell, T. D. Moustakas,

Solid State Commun. 20, 969 (1976).

5. R. J. Temkin, W. Paul, and G. A. N. Connell, Adv. Phys. 22,

581 (1973).

6. A. Matsuda, M. Matsumura, K. Nagiakgaqwa, T. Imura, H.

Yamanoto, S. Yamusaki, H. Okushi, S. lizima, and K.

Tananaka, in Tetrahedrallv Bonded Amorphous Semiconductors,

edited by R. A. Street, D. K. Biegelson, and J. C. Knights,

American Institute of Physics, New York, 1981, p. 192.

7. S. Zhao and S. Kunklinger, Solar Energy Mater. 6, 22 3

(1982).

8. D. L. Staebler and C. R, Wronski, Appl. Phys. Lett. 31, 292

(1977).

9. D. L. Staebler and C. R. Wronski, J. Appl. Phys. 51, 3262

(1980).

77

10. M. Stutzmann, W. B. Jackson, and C. C. Tsai, Phys. Rev. B

32, 6062 (1985).

11. R. A. Street, C. C. Tsai, J. Kakalios, and W. B. Jackson,

Phil. Mag. B 56, 305 (1987)

12. J. Crank, The Mathematics of Diffusion. Clarendon, Oxford,

England, 1975, Chap. 2.

13. J. Kakalios, R. A. Street, and W. B. Jackson, Phys. Rev.

Lett. 59, 1037 (1987).

14. W- B. Jackson, Phys. Rev. B 38, 3595 (1988).

15. n . B. Jackson and J. Kakalios, Phys. Rev. B 37, 1020

(1988).

16. J. Shinar, R. Shinar, S. Mitra, and J. Y. Kim, Phys. Rev.

Lett. 62 2001 (1989).

17. W. B. Jackson, Phys. Rev. B 41, 1059 (1990).

18. W. B. Jackson and M. D. Moyer, Phys. Rev. B 36, 6217

(1987).

19. M. Albers, Ph.D. Thesis, Iowa State University, (1987)

unpublished.

20. J. Tauc, R. Grigorovicci, and A. Vancu, in Amorphous and

Liquid Semiconductors, edited by J. Tauc, Plenum, New York,

1974.

21. F. Urbach, Phys. Rev. 92, 1324 (1953).

22. M. H. Brodsky, M. Cardona, and J. J. Cuomo, Phys. Rev. B

16, 3556 (1977).

78

23. H. R. Shanks, C. J. Fang, L. Ley, M. Cardona, F. J. Demond,

and S. Kablitzer, Phys. Status Solidi (b), 100, 43 (1980).

24. C. J. Fang, K. J. Gruntz, L. Ley, M. Cardona, F. J. Demond,

G. Muller, and S. Kalbitzer, J- Non-Cryst. Solids 35 & 36,

255 (1980).

25. M. Cardona, Phys. Status Solidi (b), 118, 463 (1983).

26. H. R. Shanks, F. R. Jeffrey, and M. B. Lowry, J. de

Physique 42, C4-773 (1981).

27. H. Wagner and W. Beyer, Solid State Commun. 48, 585 (1983).

2S. C. P. Poole, Jr., in Electron Spin Resonance. 2nd Ed.,

Wiley, New Yo-^k, 1983, p. 409.

29. F. M. Smits, Bell Sys. Tech. J. 710 (1958).

30. The SIMS measurements were performed by Dr. Ruth Shinar at

Microelectronics Research Center, Iowa State University,

Ames, Iowa 50011.

31. J. Baum, K. K. Gleason, A. Pines, A. N. Garroway and J. A.

Reimer, Phys. Rev. Lett. 56, 1377 (1986) .

32. Y. J. Chabal and C. K. N. Patel, Rev. Mod. Phys. 59, 835

(1987).

33. H. Wagner and W. Beyer, Solid State Commun. 48, 585 (1983).

34. A. K. Mahan, P. Raboisson, D. L. Williamson and R. Tsu,

Sol. Cells 21, 117 (1987).

35. A. H. Mahan, P. Raboisson and R. Tsu, Appl. Phys. Lett. 50,

335 (1987).

79

36. A. H. Mahan, P. Menna and R. Tsu, Appl. Phys. Lett. 51,

1167 (1987).

37. W. Beyer, H. Wagner and H. Mell, MRS Symp. 49, 189 (1985).

38. V. G. Bhide, R. O. Dusane, S. V. Rajarshi, A. D. Shaligrain

and S. K. David, J. Appl. Phys. 62, 108 (1987).

39. Y. J. He, M. Hasegawa, R. Lee, S. Berko, D. Adler and A.-L.

Jung, Phys. Rev. B 33, 5924 (1986).

40. S. C. Moss and J. F. Graczyk, Phys. Rev. Lett. 23, 1167

(1969) .

XT T T.T -r Oxiov J./V O.Xi^ ri* X'OU.X/ U * O OOX

(1972) .

42. G. S. Cargill, Phys. Rev. Lett. 28, 1372 (1972).

43. R. J. Temkin, W. Paul and G. A. N. Connell, Adv. Phys. 22,

581 (1973).

44. N. J. Shevchik and W. Paul, J. Non-Cryst. Solids 16, 55

(1974) .

45. P. D'Antonio and J. K. Konnert, Phys, Rev. Lett. 43, 1161

/ r \ r s \ v x 3 / .

46. A. <J. Leadbetter, A. A. M. Rashid, R. I»I. Richardson, A. F.

Wright and J. C. Knights, Solid State Commun. 33, 97 3

(1980) .

47. P. D'Antonio and J. H. Konnert, AIP Conf. Proc. 73, 117

(1981).

48. J. F. Graczyk, Phys. State Sol. (a) 55, 231 (1979).

80

49. R. Vanderhaghen, B, Chaurand and B. Drevillon, in P. Pinard

and S. Kalbitzer (eds.), Polv-microcrvst.alline and

Amorphous Semiconductors. Editions de Phys., Les Ulis,

France, 1934, p. 661.

50. R. Vanderhaghen, B. Chaurand and B. Drevillon, Thin Solid

Films 124, 293 (1985).

51. T. A. Postol, C. M. Falco, R. T. Kampwirth, I. K. Schuller

and W. B. Yelon, Phys. Rev. Lett. 45, 648 (1980).

52. A. J. Leadbetter, A. A. M. rashid, N. Colenutt, A. F.

Wright and J. C. Knights, Solid State Commun., 38, 957

(1981).

53. R. Bellissent, A. Chenevas-Paule and M. Roth, J. Non-Cryst.

Solids 59/60, 229 (1983),

54. SAXS measurements were performed by Mr. Yen Chen at

Department of Physics, Colorado School of Mines, Golden, CO

80401.

55. D. L. Williamson, A. H. Mahan, B. P. Nelson and R. S.

Crandall, Appl. Phys. Lett. 55, 783 (1989).

56. A- Guinier and G. Fournet, Small-Angle Scattering of X-

ravs. John Wiley, New York, 1955, p. 30.

57. G. Porod, in Small Angle X-rav Scattering, edited by O.

Glatter and O. Kratky, Academic Press, New York, 1982, p.

17.

58. H. Jia, J. Shinar, Y, Chen, and D. L. Williamson, Mat. Res.

Soc. Proc= 258, 281 (1992).

81

59. M. L. Albers, J. Shinar, and H. R. Shanks, J. Appl. Phys.

64, 1859 (1988).

60. A. H. Mahan, D. L. Williamson, B. P. Nelson and R. S.

Crandall, Sol. Cells 27, 455 (1989).

61. H. Jia, J. Shinar, and H. S. Wu, Bull. Am. Phys. Soc. 37,

128 (1992) .

62. D. K. Biegelsen, R. A. Street, C. C. Tsai, and J. C.

Knights, Phys. Rev. B 20, 4839 (1979).

63. X.-M. Tang, J. Weber, Y. Bear, and F. Finger, Phys. Rev. B

I ^ 4 t \ Zf \J ) •

64. R. Shinar, J. Shinar, H. Jia, and X.-L. Wu, Phys. Rev. B

47, 9361 (1993).

65. S. J. Pearton, J. W. Corbett, and T. S. Shi, Appl. Phys. A

43, 153 (1987)

66. S. Roorda, S. Doom, W. C. Sinke, P. M. L. O. Scholte, and

E. van Loenen, Phys. Rev. Lett. 62, 1880 (1989).

82

farr xwo.

fabrication and characterization

cvd diamond films and devices

83

I, INTRODUCTION

A. Historical Background

Diamond is one of the most attractive materials, as it

has a unique combination of excellent mechanical, physical,

and chemical properties.^ The well known properties of diamond

are summarized and compared with silicon in Table 1.1. Its

4000^0 melting point and its resistance to all chemical acids

make it an ideal material for harsh environments. The wide

band gap of diamond can be used in far-ultraviolet detectors

to eliminate the visible and infrared background which

normally appear in silicon based detectors. It is a highly

electrically insulating material but can be doped both p and n

type, and its superb thermal conductivity render it an

excellent heat sink which, if applied in integrated circuits,

will enable a sharp increase in the density of integrated

circuits.

The industrial demand for diamond was inhibited by its

prohibitive cost or unavailability until the l950s2/2 when

General Electric successfully developed a high-pressure high-

temperature (HPKT) method to synthesize it. Today 30 - 40 tons

of HPKT synthetic diamond are produced worldwide each year,

and all of it is used for cutting, grinding, and polishing

tools. However many potential applications of diamond require

84

Table 1.1 Comparison of semiconductor properties between

diamond and silicon (from Ref. 1)

Properties Diamond Silicon

Lattice constant (nmi (A)) 0.3567 (3.557) 0.5430 (5.430) Thermal expansion (xlO'V''C)1.1 2.6 Density (g/cm^) 3.515 2.328 Melting point CC) 4000' 1420 Band gap (eV) 5.45 1.1 Saturated electron velocity

(x 10' crn/s) 2.7 1.0 Carrier mobility (cm'7(V • s))

Electron 2200 1500 Hole 1600 600

Breakdown (xio^ V/cm) 100 3 Dielectric constant 5.5 11.8 > - , . » . , » c

«-( OQ 1 Qin/I7> ' f\ } . . . w w w J I J

i n*3 1 \J

-< r\3 i \J

Thermal conductivity (W/(cm • K)) 20 1.5

Absorption edge (/xm) v.d. 1.4 Refractive index 2.42 3.5 Hardness (kg/mm') 10000 1000 Johnson figure of m.erit

( x lO^^ (W-0 ) / s2 ) 73 856 Q O ^

Keyes ligure oi merit ( x lO -VV / ( cm-s - "C ) ) 444 1 « » W . V-/

85

thin films of coating which can not be produced from either

natural or HPHT synthetic diamond.

The phase diagram shown in Figure 1.1 clearly indicates

that diamond is stable only in the high pressure high

temperature region as obtained in the GE process. Yet that is

not the only way that diamond can be synthesized. If the

activation energy between stable and metastable states is high

enough, it serves as a barrier to interconversion, and

metastable material can be formed under kinetically controlled

conditions similar to numerous other reactions. Hence a

parallel effort directed toward the growth of diamond at low

pressures where it is the metastable state was explored

several decades ago.

The first successful attempt was reported by W. G.

Eversole^ who initiated the work in 1949 and achieved growth

on diam.ond seed at the end of 1952 by using decomposition of a

carbon-containing gas. In Eversole's work, the diamond and

graphite were first codeposited in a seed crystal, then the

substrate and deposit were transferred into a separate

autoclave. The graphite was removed by hydrogen etching at

temperatures T > 1000°C. The process was repeated many times

to achieve the desired thickness. Angus and co-workers

confirmed Eversole's results.They also studied the rate of

diamond and graphite growth in a CH4 - H2 gas mixture and v/ere

the first to report on the preferential etching of graphite

86

Pressure, kbar 400

?nn . /\j\j t

200

100

0

Shock wave synthesis.

Liquid carbon

a ' 1 -

Diamond & metastable graphite

Diamond \ Catalytic high-pressure, high-temperature synthesis

High-pressure, high-temperature synthesis

Chemicai vapor u e pO S1 liO n

! !

Graphite & g metastable M diamond ^

u 1000 2000 3000

Temperature, °C

conn ou vu

Fig. 1.1 Carbon phase diagram showing the stable diamond

phase at high pressure and temperature and

metastable region at low pressure

87

vs. diamond by atomic hydrogen and to demonstrate the

deposition of boron-doped semiconducting diamond.

The worldwide interest in low pressure chemical vapor

deposition (CVB) of diamond started in 1982 after N. Setaka,

Y. Sato, M. Kamo, and S. Matsumoto in Japan published a series

of papers describing in detail the techniques for synthesizing

diamond at rates of several |im per hour from microwave or DC

discharges, or from gases decomposed by a hot filament.

The processes produced individual faceted crystals without

using diamond seed crystals. Since then the CVD of diamond

from hydrocarbon precursors has been extensively studied and

become the most successful technique for metastable diamond

growth.

B. CVD of Diamond Films

Thin diamond films have been successfully grown at low

pressure by a large variety of energetically assisted CVD

processes. These can be basically divided into two groups:

thermally assisted CVD, and plasma assisted CVD (PACVD). They

all share a common feature: the carbon-containing gas is

decomposed into atom.ic species either thermally or by a

plasma.

Figure 1.2 shows a hot filament assisted CVD system,

which is typical of the thermally assisted process. The

filament, which inay be W. Ta.. Mo_. or Re, is heated un to

88

AC power

SUDDIV

t

Gas inlet

Heater

CH Gas exit

Substrate

Fig. 1.2 Schematic of hot filament assisted CVD system

89

2500°C. The CH4 gas is heavily diluted with H2 at a ratio of

-1/100; -5% of the mixture is then dissociated into atoms or

radical fragments when it passes through the hot filament. The

substrate which is about 1 cm under the filament is kept at

70D-1000°C. The advantage of hot wire CVD is its ease of

operation. The disadvantage is that the film may be

contaminated by emission of metal atoms from the hot filament.

Figure 1.3 shows a DC plasma assisted CVD. A stabilized

DC plasma produced growth rates up to -20 jam per hour at the

anode. However, only graphite deposited on the cathode. The

typical applied voltage and current density are 1 kV and 4

A/cm^ respectively.

The DC plasma jet technique (Figure 1.4) focuses the

diamond coating on a relatively small (5x5 mm^) area, but the

plasma can be scanned over a large area to produce a uniform

deposition. This technique uses a high temperature plasma in

which nearly all molecules are completely dissociated. Growth

rates of more 500 jim per hour have been achieved by this

method.

Microwave plasma CVD offers another way to deposit

diamond films. Figure 1.5 shows a bell jar reactor for this

technique. The microwave radiation is guided to the top of the

reactor by a rectangular metal wave guide and picked up by an

sntGnns tlist p^rotiruciss into ths ci^rculsr*. Tiis iTiixtiiirs cz

m o 3 o p >-i •» c- ^ • c- ^ •— W - » • W . • W. • • jj' S_> C >>.A. C% * « .4- Alt Ck. O A 1 ±. CX

90

Water Gas exit

Substrati

holder

DC

I I — r r

Gas

inlet CH

Water

Fig. 1.3 Schematic of a DC plasm.a assisted CVD system

91

dc source

i 1 { ! Argon & hydrogen

Cooling

Methane

Plasma jet

• vw X caic

U*«-A.>S.,JUWK.A *—Al*«^

! 1 •

i T » • aako/^ w'wv./mti^ riuici

Fig. 1.4 Schematic of a DC plasma jet system

92

AnlGnns

Microwaves

Rectangular metai waveguide

Circular, transparent metal waveguide

Silica bell jar

Heater

Ball-shaped plasma -

Subslrale

MCIMSHB & hydrogen

i

m

\ r ^ \

S T o D u m p

r—1 m 3

\\\\\\\\\\\X\\\\\\\\\\\\\\\\\ ^

Fig. 1.5 Schematic of a bell jar microwave reactor system

93

FLOWMETERS TORCH

o2

c2h2

SUBSTRATE

PYROMETER

f1 I

WATER

X Cu MOUN

INNER CONb

ACETYLENE PEA 1HtH

OUTER t-LAMt

Fig. 1.6 Schematic of an cxy-acetylene torch CVD sysrsm

94

plasma above the substrate. This technique can deposit diamond

films on substrates up to 8 cm in diameter.

Another noteworthy technique is the combustion flame

shown in Figure 1.6, which uses an oxygen-acetylene torch and

a water cooled substrate. It can deposit high quality diamond

films in an ambient atmosphere, and the growth rate can be as

high as 180 [im per hour.l^"^^

C. Growth Mechanism of CVD Diamond

As discussed earlier and shown in the phase diagram

(Figure 1.1), graphite is the stable form of carbon under the

conditions used in CVD of diamond. Why is it then possible to

grow diamond at less than atmospheric pressure at temperatures

750° < T < 1100°C?

Although an established mechanistic answer to this

question is not yet available and is still the subject of

current research, one model consistent with experimental

results was developed by Spear and coworkers. 0 ? 21

mechanism is based on the fact that growth occurs at the gas-

solid interface of the carbon-hydrogen system. The vapor-

growth process does not involve elemental carbon only, but

also hydrogen. A diamond surface saturated v.'ith sp^ C-K bonds

is more stable than a diamond surface free of hydrogen. Once a

surface carbon is covered by another diamond growth layer.

95

DIAMOND

GRAPHITE

structure of diamond and graphite. The hydrogen

atoms depict their role in stabilizing the diamond

surface

96

then that covered carbon possessing four sp^ C-C bonds is

metastable with respect to graphite.

Lander and Morrison^^ were the first to hypothesize that

hydrogen can stabilize a diamond surface by forming sp^ C-K

bonds with surface carbon. Without the hydrogen maintaining

the sp^ character of these surface carbons, it is easy to see

that the diamond surface will collapse into the more stable

planar graphite structure (Figure 1.7) during the growth

process.

Hydrogen is also responsible for other important

mechanisms during diamond growth. As proposed by Frenklach and

Spear23^ the surface bonded H is first removed by atomic H

from the gas phase to form an activated carbon radical and H2

c c \ \

H- + C-H J- H2 + C- 1-1

/ / c c

then this surface-activated carbon radical acts as a site for

adding more carbons to the structure by reacting with

acetylene or other carbon-hydrogen species in the gas or

plasma, e.g.,

C C H H \ \ i ' N \ I / C- + K-C=C-K > C-C=C* --2

/ / c c

97

the C=C graphite-like double bond can be reconstructed into

sp2 bonds by H2•

C C C C

\ / \ / C=C + H-H > H-C-C-H ^-3

/ \ / \ c c c c

Deryagin and Fedoseev^^ proposed that a "super

equilibrium" concentration of atomic hydrogen at the growth

surface is responsible for the major reduction in graphite

codeposition. They argued that atomic hydrogen behaves like a

"solvent" for graphite. In fact, numerous studies24-26 show

that at T > 8 00°C atomic hydrogen etches graphite at rate

almost 500 time faster than it etches diamond. That results in

the net growth of diamond under metastable conditions.

D. Scope or Tnis worx

'T'1 0 vork began with the constructiori of a new hot

filament assisted CVD system for diam.ond deposition by using

minimum, resources of equipment that already existed in the

laboratory. After the successful deposition of films, the goal

was expanded into (i) studies of native defects in thin

diamond films, and (ii) fabrication of light emitting diodes

(LEDs) from both intrinsic and doped films. The major

techniques used in this study were x-ray diffraction (XRD) to

98

characterize the diamond structure, scanning electron

microscope (SEM) to characterize the morphology of the films,

and ESR spectrometry to study their native defects.

99

II. SAMPLE FABRICATION

A. Deposition System

All samples used in this work were prepared by using a

home built hot filament assisted CVD system. A schematic of

the deposition system is shown in Figure 2.1. Using a

mechanical pump, the deposition chamber was routinely

evacuated to a base pressure of -3xlO~2 Torr prior to each

deposition. The 0.5 mm diameter tungsten filament was heated

to about 2000^C by an AC power supply; the temperature of the

filament was monitored with an optical pyrometer. The

substrate was about 1 cm underneath the filament and kept at

about 800-900°C by radiation from the filament. The 0.5%/99.5%

CH4/H2 gas mixture was introduced into the chamber at a flow

rate of -200 standard cubic centimeters per minute (seem). The

deposition pressure of 50 Torr was obtained by closing the

main valve connecting the chamber and the mechanical pump, and

adjusting the micrometer valve in the bypass line.

B. Intrinsic and Doped Sample Preparation

The diamond films were deposited on 1.5x3 cm^ n or p type

(100) Si wafers. It is well known that diamond nucleation is

greatly enhanced and uniform polycrystalline diamond films are

obtained if the Si substrate is scratched with diamond powder

before deposition.The Si wafers were thus usually

100

/

TUHGSTEN FIIiMEKT

SUBSTRATE

ttrttt^ir/*v^/nrrn^ r * iniii\rika^uuci.r. —p a c?3

TO PUMP f-

\

I i I !

GLASS JAR

< chd &: k2

Fig. 2.1 Scneraatic of hot filaraent assisted CVD system

101

pretreated before deposition by either polishing or

ultrasonically bathing the wafers with 0.25 |iiti diamond paste

for about 30 minutes, followed by ultrasonic cleaning in

deionized water, acetone, and methanol. Figure 2.2 shows an

SEM image of the scratched surface.

Table 2.1 Typical parameters for CVD of diamond films

Parameter Range

CH4 (%) 0.5 - 2 . 0

h2 (%) 98.0 - 99 . 5

Total flow rate (seem) 180 - 220

Deposition pressure (Torr) 40 - 60

Filament temperature (°C) 1900 - 2200

Substrate temperature (°C) 800 -• 900

r>i ^ ^ n j-'u ^ ^ ^ z i. i ^ — . - _ ^ ^ 1 * A. ^ X J. w i. w j_ wi ic; V.A ^ o u. a. V-/1X w tz:

noted that the tungsten filament became carbonized. As a

result, the resistance of the filament increased and the power

supplied to filament had to be increased accordingly to keep

the temperature of the filament at about 2000°C. The whole

deposition process normally lasted more than 3 hours depending

on the desired thickness of the film. The typical growth rate

of the film was between 0.5 - 1.5 }j.m/hr depending on the

CH4/H2 ratio.

102

Fig. 2.2 SEM images of the scratched surfaces of the Si

wafer after polishing with 0.25 )im diamond paste

103

The extreme refractive properties of diamond inhibit

doping by means such as diffusion, etc. It has become a clear

advantage that CVD diamond can be doped during deposition

process by introducing the dopant in the reactant gas.27

Attempts to deposit n type films were also made in this work,

to improve the light emitting efficiency of diamond LEDs. The

dopant gas was N2 heavily diluted with H2.

C. Light Emitting Diode (LED) Structure

The technology for GaAs-based yellow-orange LEDs has long

been established. Many efforts have also been directed at

development of blue LEDs. Diamond is one of the candidates for

these devices, since it has a wide band gap of 5.4 eV, and a

high breakdown voltage of a few MV/cm. Although 5.4 eV

photons are clearly well within the UV range, an intragap

donor-acceptor transition yields blue electroluminescence

(EL). A LED structure as shown in Figure 2.3 was also

fabricated from diamond films to study the EL.

The diode fabrication process began first with the

deposition of a 3-10 [im thick diamond film, on silicon wafer.

The silicon substrate was then dissolved by HF to yield a

free-standing film. A thin metal film (A1 or Ga) of 0.2 ,um,.

and indium tin oxide (ITO), which is a transparent conductor,

were then deposited on each side of the free standing diamond

film by electron beam evaporation.

104

ITO

Fig. 2.3 Schematic structure of a thin diamond film LED

105

III. SAMPLE CHARACTERIZATION

A. Thickness Measurements

The thickness and surface roughness of the films was

measured using a Sloan Dektak profilometer as described in

Part One. SEM provided an additional method to determine the

thickness and morphology of the films by using a broken

section of the sample, as shown in Figure 3.1 in which the

thickness of the film was determined to be -3 0 jam.

B. Scanning Electron Microscopy

Samples were usually studied with an optical microscope

(Olympus Vanox) to monitor surface uniformity and morphology

on an optical scale. When deemed necessary, SEM images were o

also obtained. The SEM has a resolution of -250 A, and a

depth of field 300 times more than that of the optical

microscope, all of which result in photographs of dramatic

three-dimensional quality.

Prior to each SEM measuremenr, a thin layer of gold was

DC sputtered on the diamond film to prevent the charge

accumulation which can interfere with the scanning electrons

and cause smear output image. Figure 3.2 shows typical SEM

pictures of diamond films taken with a Cambridge Stereoscan

200 SEM. In general, high quality diamond films exhibit well

106

Fig. 3.1 SEM photograph of CVD diamond film,

of the film is determined to be -30

picture

the thickness

um from the

107

108801

Fig. 3.2 Typical SEM images of CVD diaip.ond filir.s

108

faceted crystal structure with (001) or (111) planes. Figure

3.3 illustrates the change from cubic to octahedral diamond.28

The diamond films were also characterized by powder x-ray

diffraction (XRD). They were directly mounted on a microscope

slide. The XRD spectra were obtained by using a microcomputer-

controlled Rigaku diffractometer equipped with a cooper target

and a diffracted beam graphite monochromator for Cu Ka

radiation with a step scan rate of 0.01 degree/second over the

20 angular range of 20-95 degrees.

The basis of XRD is Bragg's law^^ which describes the

condition for constructive interference for x-rays scattered

from atomic planes of a crystal:

Here d is the interplanar spacing of the crystal, 20 is the

angle of the diffracted beam, and X is the wavelength of the

incident x-rays as shown in Figure 3.4. The distance between

the (hkl) planes of a crystal is given by

C. X-ray Diffraction Measurements

2dsin0 = X 3.1

3 . 2

109

Cubo-octohedrons Cube

Octohedron

Fig. 3.3 Polyhedrals of diamond

110

where a = 3.56 A is the lattice constant of diamond. For

face-centered cubic (fee) structures such as diamond, the

indices (hkl) have to be either all even or all odd in order

to yield a diffracted beam. In this work the incident beam was o

the Cu Ka line of wavelength ?l=1.5418 A. The expected XRD

pattern from diamond, as determined by Eq (3.1) and (3.2), are

shown in Table 3.1.

Figure 3.5 shows a typical XRD spectrum of a diamond

film. Three sharp peaks are identified as the diffraction from

(111) , (220) , and (311) planes. A broad peak at 20 = 70*^ is

due to the (400) Si substrate.

D. Electron Spin Resonance (ESR) Measurements

ESR was discussed earlier in Part One, except for the

case of a hyperfine interaction between the electron and a

nuclear spin.

In this case, the spin Hamiltonian is given by^*^

H = cSB•S + AS•I 3 .:

Where the first term represents the normal Zeeman splitting,

the second term represents the hyperfine coupling between the

electron spin S and nuclear spin I, and A is the hyperfine

coupling constant. Assuming gp3*3 is much larger than AS"I,

Ill

Table 3.1 Expected XRD patterns of polycrystalline diamond

films

(hkl) (111) (220) (311) (400)

20(°) 44.1 75 . 5 91.8 120. 0

Incident Beam

Detector

Fig. 3.4 Schematic diagram of XRD measurement, the detector

can be scanned at rate of 0.01 degree/second over

range of 20-95 degree

112

1250

1000

^ 750

ijo C Q> 500

250

I I I i i 1 i I i i i i i i i i i i i I i 1 i

(111)

u • J l .

I rrr m i—i—:—i—:—:—l—1_

(220)

11

J 11 i U 1 L

_1

J

(311)

I ' ' ' W1 i 1 I I

40 50 60 70 80

Diffract angle- 29 (

or* 1 ao c5 v-/

Fig. 3.5 Typical XRD spectrum of CVD diamond- The broad

pattern at 20 = 7 0° is due to the Si substrate

113

the energy levels obtained from first order

perturbation theory:

where M and m are the magnetic quantum numbers of the electron

and nucleus, respectively. The level structure and allowed

electron spin transitions between them are illustrated in

Figure 3.5 for the case S=l/2, 1=1/2.

The allowed ESR transitions are AM=±1, Am=0. Therefore,

the energy absorbed by the spin system when a microwave photon

is absorbed and the state changes to iM,m> is given

by

hv+ = g.BB + Airv 3 . 4

Here iTi " il/2 • This gives irise: co cwo resonance lines

separated by A from, each other, or by A/2 from B = hVQ/g!3.

Two satellite peaks in the ESR of CVD diamonds are indeed

observed as shown in Figure 3.7. The origin of the two

satellites will be identified and discussed in the next

section.

E. Electroliininescsiice Measurements

To investigate the electroluminescence (EL) from diamond

LEDs, a simple device was fabricated as shown in Figure 2.3.

114

M

+1/2

hv.

-1/2

/

hva.

J

hv_

m

-J./

-1/2

-1/2

•1/2

Fig. 3.6 Energy-level diagram for case S=l/2, 1=1/2. The

left side shows the Zeeman splitting only, the

right side shows the further hyperfine splitting.

The allowed ESR transitions are also indicated

115

< 1 1 i 1 1 1 1 1 i i i 1 i 1 1 : 1 • 1 ; i i j ' i —I—

Ns = 3.7E+17 gm~^

= 2.0026 -

CH.(0.5%) 4- H2(99.5%) ~

1 _

I 1 i 1 . 1 1 1 1 1 r 1 1 1 1

1 1

1 1

r 1

1 1

1 1

1

3.3 3.32 3.34 3.36 3.38 Magnetic Field (kG)

Fig. 3.7 ESR spectrum of diainond film (sample B) . The tv;o

satellites indicates hyperfine interaction between

dangling bond and nuclear spin

116

The EL spectrum and its relative intensity were measured using

an experimental set up as shown in Figure 3.8. The diamond LED

was driven by a DC power supply with voltage range 0-400

volts, and the emitted light from the LED passed through a

chopper and a monochromator (M) to a photomultiplier tube

(PMT). The signal was then sent to a lock-in amplifier, and

the spectrum was obtained by scanning the monochromator from

300 to 650 nm.

117

i

chopper

n PMT I 1

i m i ! I

• f

i

l -1

Lock-in amnjifif^r

Fig. 3.8 ExpejTiTusnta 1 set up for EL measuireinent

118

IV. RESULTS AND DISCUSSION

A. Electrolziminescence

EL was observed on both undoped and nitrogen doped

diamond films. However, the strongest emission was from an

undoped film deposited under regular conditions as described

earlier. Figure 4.1 shows a typical EL spectrum of an LED

fabricated from a 12 lam thick undoped film deposited over 24

hours- The broad blue-green emission band peaking at -450 nm.

has been observed by several groups and is called "band A"

31,32 « •

The LED characteristics varied widely among the devices.

The lifetimes varied from a few seconds to 10 minutes and

became longer as the electric field decreased. As the applied

voltage was increased, the luminance was enhanced and showed a

rapid increase above 2 00 volts. The dependence of the

luminance of a diode on the applied DC voltage is shown in

Figure 4.2. It is actually not clear whether the emission is

from the bulk of the diamond layer or originates from other

regions such as the interface between diamond and the metal or

the ITO. However, the peak position and the shape of the

spectra are very close to these of natural diamond-based LEDs

obtained by Prior and Champion.According to their

observations, blue light is emitted from the interior of the

119

1 i i 1 I i r 1 r

6 *—

4 —

2 —

l»

400 500 600

Wavelength (nm)

Fig. 4.1 Electroluminescence spectrum of diamond LED

120

iOO 200 300 400 500 T "J t Voltao'e o ~ \ • /

Fig. 4.2 The EL intensity vs applied DC voltage

characteristics of the LED

121

region. Nitrogen and some other defects in the films are

believed to serve as emission centers. Even though the films

were not intentionally doped, minute amounts of nitrogen were

probably incorporated into the films because the base vacuum

prior to deposition was higher than 10"^ torr.

B. The Native Defects of CVD Diamond Films

The nature of the defects in CVD diamond is very

important for the electronic and optical applications of these

films. In this work we identified for the first time and

reported^"^ a new defect center in CVD diamond manifest in the

hyperfine-split satellites in the ESR spectrum.

As mentioned earlier, we always observed two partially

resolved satellites in the ESR spectra, as shown for sample B

in Figure 3.7. The possible candidates responsible for the

hyperfine splitting are either (1.1% natural abundance) or

^H. They each have nuclear spin 1=1/2 and both participate in

the deposition processes. To identify which one is responsible

for the splitting, we first fit the ESR spectrum with a broad

Lorentzian, a narrow Lorentzian, and two off-center

Lorentzians as shown in Figure 4.3. The agreement is excellent

and indicates that the satellite intensity is -4% of the total

intensity. Since has 1.1% natural abundance and the

satellite intensity also varies from sample to sample, we can

rule out the oossibility of ^^c. However, we also prepared a

122

3.38 3.3 3.34 3.36 3.32 Magnetic Fieid (kG)

4.3 Fitting of the ESR spectrum by four Lorentzians

(solid lines). The dotted line is the sum of the

four Lorentzians

123

V = 9.53 GHz Ns = 3.6E-i-17 gm ^ g, = 2.0026 J ^^CK.(0.5%) ^ H2(99.5%) j

f*

3 o cj

_9

-4

3.25 3.3 3.35 3.4 Magnetic Field (kG)

Fig. 4.4 The ESR spectrum of 100% enriched diamond film.

Note the x scale is 200 G as coinpared to 100 G in

the other spectra

124

100% enriched sample (sample E) using 0.5% 12CH4 diluted

with 99.5% H2 as the reactant gases. The broad ESR spectrum

(Figure 4.4) is clearly due to the hyperfine splitting between

the dangling bonds and their neighboring nuclei.

To prove that H is responsible for the two satellites, we

replaced the H by deuterium which has much weaker nuclear spin

moment (1=1) than Indeed, the two satellites completely

disappeared from the ESR spectrum of sample C (Figure 4.5)

prepared by using 0.5% CD4 diluted with 99.5% D2 as the

reactant gases. The satellites also disappeared from sample D

prepared by using 0.5% CH4 diluted with 99.5% D2 (Figure 4.6).

Previous NMR studies^^"^"^ have indicated that most of

the hydrogen in thin diamond films resides on the crystallite

surfaces, and is removed after about 2 hours annealing at

850°C or above. However, the ESR line shape and the spin

density of sample B did not change at all even after annealing

for 1 hour from 740 to 1100°C at steps of 30°C (Figure 4.7).

Finally, the sample was annealed for 20 minutes in an Ar flow

by an oxyacetylene rorch at T > 12 00°C. The resulting ESR line

shape rem.ained essentially unchanged but the total intensity

dropped by factor of -4.5 (Figure 4.8). The satellites' strong

resistance to the high temperature annealing suggests that the

carbon dangling bond-H centers responsible for these

satellites are deeply trapped either on internal surfaces, or

125

Ns = l.lE-rl8 gm '

I,-. = 2.0028

c

o o

-5

3.38 3.36 3.32 o.o

Msignstic Fisld (kG)

Fig. 4.5 ESR spectrum of sample C. prepared from 0.5% CD4 +

99.5% D2- Note that the two satellites have

completely disappeared

126

-• 1 • • ' • 1 ' • • • 1 ' '

A

, , ^

"1 p = 9.53 GHz / Ns = 1.4E-fl8 gm~'

" = 2.0028

/ CH4(0.5%) 4- 02(99.5%)

— / -^Q'rr I , , • , I . , , , I , , , , ! , , . . I


Fig. 4.6 ESR spectru-i of sample D, prepared from 0.5% CH4 +

99.5% D2- Note that the two satellites have

completely disappeared

127

- 9.53 GHz Ns = 2.7E-rl7 gm

Annealed at llOO^^C x Ihr

-I 1

3.3 3.32 3.34 3.36 3.38 Magnetic Field (kC-)

4.7 ESR spectrum of sample Bl, which was annealed for 1

hour at steps of 30°C from 740 to 1100°C. Note that

the line shape did not change

128

• • I • ' • ' I ' ' • ' i ' • • • t ' ' '

Ns = 5.2E-rl6 gm"'

Annealed at 1500°Cx20mir


4.7 esr spectrum of sample b2, which was annealed for

20 minutes in an Ar flow by an oxyacetylene torch

at t > 1200°c. Note that the line shape was

essentially unchanged but the total intensity

drooped bv factor of -4.5

129

even in the bulk tetrahedral diamond network. Indeed such

centers may be expected since H is needed during the growth of

each layer to stabilize the surface, and hence some H atoms

may be "left behind" during this process.

All of the ESR spectra except the enriched one are

fitted into either two or four Lorentzains. The fitting

parameters and ESR results are summarized in Table 4.1. The

central narrow and broad Lorentzains, invariably observed in

all films, have been attributed to exchange-narrowed clusters

and isolated carbon dangling bonds at spinless nuclei,

respectively.^^ The two deuterated samples (C and D) showed

increasing central narrow Lorenztians may due to the

unresolved hyperfine split dangle bonds.

The hyperfine coupling constant of the H-dangling bond

complexes is equal to the field splitting between the two

satellites, i.e., about 14.5 G. Their intensity indicates a

density of about 0.03-0.2 ppm (Table 4.1), i.e., they are a

minute fraction of the total H content, which is typical 0.1-1

at. % in CVD diamond.

As mentioned earlier, the H centers could be located

either on internal surfaces or in the tetrahedral diamond

network. The SEM images showed that the average crystallite is

several m.icrons across at the top of the films. Yet if the

hydrogen

130

Table 4.1 The dangling bond density of each sample Ng, the

fraction of each Lorentzian component I, and its

derivative peak-to-peak linewidth PP. Sample B is a

regular diamond film, C and D are deuterated films,

and E is the enriched sample.

Sample

Ns

loiVi

I ( % ) PP(G)

Sample

Ns

loiVi broad narrow sat broad narrow sat

B 2.7 83 13 4 7.8 2.6 2.9

B1 2.7 78 15 7 7.5 2.6 3.0

B2 0.62 81 10 9 5.8 1.9 2.7

C 11 56 44 8.7 2.8

D 14 65 35 10.1 2.9

E 3.6

131

is essentially confined to the crystallite surface, the NMR

spin count indicates that the average crystallite is less than

O.Siim across.^® This discrepancy may be resolved by the fact

that the crystallite size rapidly increases during film

growth. It therefore does not allude to the existence of

internal surface within the observed crystallites. However,

such internal surface and/or microvoids are to be expected

within the kinetic model described by Anthony.In this

model, the hydrogenated surface is constantly etched and

repassivated by H atoms, but occasionally a methyl radical is

attached to an exposed carbon dangling bond, providing the

fundamental diamond network growth step. Such dynamical

processes may yield considerable internal defect structures.

Fanciulli and Moustakas'^'^ have indeed recently invoked

multivacancies and (111) stacking faults to explain the

properties of the ESR of their diamond films. However,

hydrogen that resides on the surfaces of such voids which are

connected to the crystallite surface will be released within a

few hours at -S50°C. In addition, a dangling bond on a largely

hydrogen-decorated surface should generally be adjacent to

more than one H atom. We therefore turn to the model of

dangling bond-H centers which are embedded in the tetrahedral

network.

In the simplest approach, substitution of a C by an H

stotp. slioiild 2705122. 1 stct. ens ijond.

132

and an interacting system of three other dangling bonds. The

ground state of this system should probably be an unpaired sp^

electron slightly overlapping the reconstructed paired

wavefunction of other two. However, this model should be

verified by a calculation of the defect wavefunction which

should yield an hyperfine splitting in agreement with the

observed value.

133

REFERENCES

1. J. E. Field, The Properties of Diamond. Academic Press,

London, 1979.

2. F. P. Bundy, H. T. Hail, H. M. Strong, and R. J. Wentorf,

Nature 176, 51 (1955).

3. P. W. Bridgman, Sci. Am. 193, 42 (1955).

4. W- G. Eversole, US Patents 3,030,187 and 3,030,188 (1962).

5. J. C. Angus, H. A. Will, and W. S. Stanko, J. Appl. Phys.

39, 2915 (1968).

6. J. C. Angus and N. C. Gardner, US Patent 3,661,526 (1969).

7. J. C. Angus, Sint. Almazy 3, 38 (1971).

8. S. Matsumoto, Y. Sato, M. Kamo, and N. Setaka, Jpn. J.

Appl. Phys. 21, 183 (1982).

9. S. Matsumoto, Y. Sato, M. Tsutsumi, and N. Setaka, J.

Mater. Sci. 17, 3106 (1982).

10. M. Kamo, Y. Sato, S. Matsumoto, and N. Setaka, J. Cryst.

Growth 62, 642 (1983).

11. Y. Matsui, S. Matsumoto, and N. Setaka, J. Mater. Sci.

Lett. 2, 532 (1983).

12. S. Matsumoto, ibid. 4, 600 (1985).

13. S. Matsumoto, M. Hino, and T. Kobayashi, Appl. Phys. Lett.

51, 737 (1987).

14. K. Suzuki et al., Appl. Phys. Lett. 52, 728 (1987).

P. snd, Msssis^r, 2— ^2.989)

134

16. Y. Hirose and M. Mitsuizumi, New Diamond 4, 34 (1988).

17. Y. Hirose, Proc. of 1st intl. Conf. on New Diamond Sci. and

Tech.. Tokyo, Japan, Oct. 24-26, 1988.

18. L. M. Hansen, W. A. Carrington, J. E. Butler, and K. A.

Snail, Mat. Lett. 7, 289 (1988).

19. Y. Tzeng et al., Appl. Phys. Lett. 56, 134 (1990) .

20. K. E. Spear, Earth Miner. Sci. 56, 53 (1987).

21. K. E. Spear, M. Frenklach, A. Badzian, T. Badzian, and R.

Messier, Ceram. Eng. Sci. Proc. 9, 1095 (1988).

22. J. J. Lander and J. Morrison, Surf. Sci. 4, 241 (1966).

23. M. Frenklach and K. E. Spear, J. Matt. Res. 3, 133 (1988).

24. B. V. Deryagin and D. V. Fedoseev, Growth of Diamond and

graphite from Gas Phase. Ch. 4 Izd. Nauka. Moscow, USSR,

1977.

25. B. V. Deryagin and D. V. Fedoseev, Sci. Am. 233, 102

(1975) .

26. N. Setaka, Proc. of 10th intl. conf. on CVD. Edited by G.

w. Cullen and J. Bloche. Electrochemical Society,

Pennington, NJ 1987.

27. N. Fujimori, T. Imai, and A. Doi, Vacuum 36, 99 (1986).

28. K. Kobashi et al., Phys. Rev. B 38, 4067 (1988).

29. B. D. Cullity, Elements of X-ray Diffraction, 2nd ed.

(Addison Wesley, Reading 1978) .

30. J. Loubser and J. V. Wyk, Rep. Prog. Phys. 44, 1201 (1978).

31. Y. Taniguchi et al., Jpn. J. Appl. Phys. 23, L1S48 (1989).

135

32. M. Kadono et al.. Mat. Res. Symp. Peoc. 162, 358 (1990).

33. J. R. Prior and F. C. Champion, Proc. Phys. See. 80, 849

(1962) .

3 4. H. Jia, J. Shinar, D. P. Lang, and M. Pruski, Phys. Rev. B

48, 17593 (1993).

35. J. M. Redwing, T. W. Root, and W. F. Banholzer, J. Appl.

Phys. (submitted).

3 6. K. M. McNamara, D. H. Levy, K. K. Gleason, and C. J.

Robinson, Appl. Phys. Lett, 60, 580 (1992).

37. S. Mitra and K. K. Gleason, Diamond and Related Materials

2, 126 (1993).

38. M. Pruski, D. P. Lang, S. J. Hwang, H. Jia, and J- Shinar,

Phys. Rev. B 49, 10635 (1994).

39. T. R. Anthony, Mat. Res. Soc. Pro. 162, 61 (1990).

40. M. Fanciulli and T. D. Moustakas, Phys. Rev. B 48, 14982

(1993) .

136

GENERAL CONCLUSIONS

Semiconductor materials of a-Si:H, CVD diamond films and

the related device have been fabricated and studied in this

work. The of a-Si:H increased by order of four after

annealing at -3 00°C, and the film showed weak light induced

degradation. The SAXS determined microvoid content was found

to be strongly correlated with the IR determined H content,

and high K content films showed columnar microstructure while

low H content films showed spherical or randomly oriented

voids. In the studies of H diffusion in a-Si:H, we found for

the first time the negative dispersion parameter a, which is

contradictory to the widely invoked "multiple trapping" model.

The CVD diamond based LEDs were successfully fabricated,

and a blue-green electroluminesence peaked at wavelength of

-450 nm was observed from such a device. A new H center which

may be deeply trapped either on internal surface or in the

bulk of diamond was identified for the first time through the

hyperfine features in the ESR signal. Further calculation

should be able to show the exact configuration of this new H

center.

137

ACKNOWLEDGMENTS

I would like to thank my research advisor Dr. Joseph

Shinar for his guidance, encouragement, and support throughout

this work.

My sincere appreciation goes to Dr. Ruth Shinar for her

collaborative work on SIMS, Drs. Don Williamson and Yen Chen

for their SAXS measurements, Dr. Scott Chumbley and Fran Laabs

for allowing me to use the SEM, and Drs. Davis Johnston and

Lance Miller for allowing me to use the x-ray diffractometer.

I am also very grateful for the help and counsel of my

graduate committee: Drs. Bing-Lin Young, Kai-Ming Ho, Alan

Goldman, Hsung-Cheng Hsieh, Michael Tringides, and Che Ting

Chan.

I also thank the physics department and Ames laboratory

for making the facilities available. Ames laboratory is

operated for the US Department of Energy by Iowa State

University under contract no. W-74 05-eng-8 2.

Finally, I would like to thank my wife Li Li to whom this

dissertation is dedicated. Her love, patience, and

understanding has made this research possible.

fabrication and characterization of hydrogenated amorphous

Documents