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N BS SPECIAL PUBLICATION 400-23

U.S. DEPARTMENT OF COMMERCE / National Bureau of Standards

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ARPA/NBS Workshop IV.

Surface Analysis for

Silicon Devices
NATIONAL BUREAU OF STANDARDS

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Semiconductor Measurement Technology:

ARPA/NBS Workshop IV.

Surface Analysis for Silicon Devices

A. George Lieberman

Electronic Technology Division

Institute for Applied Technology

National Bureau of Standards

Washington, D.C. 20234

This activity was supported by

The Defen.-e Advanced Research Projects Agency

under ARPA Order 2397, Program Code 5D10

and The National Bureau of Standards Washington, D.C. 20231

U.S. DEPARTMENT OF COMMERCE, Elliot L. Richardson, Secretary

James A. Baker, III, Under Secretary

Dr. Betsy Ancker-Johnson, Assistant Secretary for Science and Technology

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Acting Director

Issued March 1976
Library of Congress Cataloging in Publication Data Main entry under title: ARPA/NBS workshop IV. (Semiconductor measurement technology) (NBS Special

publication ; 400-23)

"Contains the proceedings of the ARPA/NBS workshop IV, Surface analysis for silicon devices, held at the National Bureau of Standards on April 23-24, 1975.'' Supt. of Docs. no.. 0 13.10:400-23.

1. Semiconductors—Testing —Congresses. -. Silicon

Testing—Congresses 3. Surfaces (Technology)—Congresses.

4. Spectrum analysis—Congresses. I. Lieberman, Alfred George, 1937- II. United States. National Bureau of Standards. III. Series.

IV. Series: United States. National Bureau of Standards.Special

publication ; 400-23. QC100.U57 no. 400-23 [TK7871.85] 602 .1s [621.3815'2]

76-608043

National Bureau of Standards Special Publication 100-2

Nat. Bur. Stand. (U.S.), Spec. Publ. 400,23, 238 pages (Mar. 1476) CODEN: XN-BSAV
PREFACE

The ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices} was held at the

National Bureau of Standards on April 23-24, 1975 under the cosponsorship of the National

Bureau of Standards and the Defense Advanced Research Projects Agency. It was the first

meeting, on a national level, to address the applicability of modern surface measurement

techniques to the analysis of semiconductor devices. Representatives from industrial,

governmental and academic organizations concerned with device manufacture, analysis and

instrument design participated in this workshop. Speakers were selected from among the

finest and most active researchers in the field. The outcome was a large and enthusiastic

workshop attendance. It is hoped that this report will convey the spirit of the workshop

to those who could not attend.

A. George Lieberman

Editor

iii
ABSTRACT

This report contains the proceedings of the ARPA/NBS Workshop IV, Surface Analysis for

Silicon Devices, held at the National Bureau of Standards on April 23-24, 1975.

The Workshop, as part of an NBS program to develop measurement technology for the

field of semiconductor devices, was held to discuss the present capabilities and future

prospects of modern analytical beam techniques as applied to silicon, and associated

insulator films and device structures. Of particular interest were the determination of

impurity profiles, surface contamination, and interface characteristics. Techniques

utilizing impinging electron, ion, neutral or photon beams were considered. The Workshop

was directed at the analysts, the semiconductor manufacturers who use the analysts' re- sults, and the instrument people who design and manufacture the analytical equipment.

Transcripts of the discussions following each paper are also included within these proceedings .

Key words: Auger spectroscopy; depth profiles; electron beam induced imaging; ESCA;

insulator films; interface characteristics; internal photoemission ; ion scattering spectroscopy; nuclear resonance profiling; photodepopulation ; photovoltaic imaging; Rutherford

backscattering; SCANIIR; secondary ion mass spectroscopy; semiconductor devices; silicon;

surface analysis; X-ray photoelectron spectroscopy.
TABLE OF CONTENTS

PAGE

WELCOMING REMARKS 1 Judson C. French, Chief, Electronic Technology Division

National Bureau of Standards, Washington, DC 20234

INTRODUCTORY CONCEPTS FOR SILICON SURFACE ANALYSIS 3

A. George Lieberman, Electronic Technology Division,

National Bureau of Standards, Washington, DC 20234

IDENTIFICATION OF INTEGRATED CIRCUIT PROCESS AREAS AMENABLE TO DIAGNOSIS AND CONTROL

BY ANALYTICAL BEAM TECHNIQUES 7 Bruce E. Deal, Integrated Circuits R&D

Fairchild Camera and Instrument Corporation, Palo Alto, CA 94304

ION AND NEUTRAL BEAMS - TECHNIQUES AND APPLICATIONS

LOW ENERGY ION SCATTERING SPECTROMETRY STUDIES OF Si, Si0 2 AND RELATED MATERIALS . . 21 William L. Harrington, David Sarnoff Research Center

RCA Laboratories, Princeton, NJ 08540

SURFACE ANALYSIS BY SECONDARY ION MASS SPECTROSCOPY TECHNIQUES 31 Robert D. Dobrott, Materials Characterization Laboratory

Texas Instruments Incorporated, Dallas, TX 75222

SOME EFFECTS LIMITING SIMS DEPTH PROFILE ANALYSIS AND METHODS FOR IMPROVEMENT .... 45

Robert K. Lewis, Applications Laboratory

Cameca Instruments, Incorporated, Elmsford, NY 10523

QUALITATIVE ASSESSMENT OF ION EROSION DAMAGE BY MEANS OF ELECTRON CHANNELING

PATTERNS 61 Dale E. Newbury, Analytical Chemistry Division

National Bureau of Standards, Washington, DC 20234

THE EFFECT OF SPECIMEN COOLING ON THE MIGRATION OF SODIUM IN THIN FILM Si0 2 65 Bradway F. Phillips, Naval Weapons Support Center, Crane, IN 47522

Alfred E. Austin, Battelle Columbus Laboratories, Columbus, OH 43201

Harold L. Hughes, Naval Research Laboratory, Washington, DC 20375

SILICON-ON-SAPPHIRE IMPURITY ANALYSIS 73

D. Howard Phillips, Electronics Research Division

Rockwell International Corporation, Anaheim, CA 92803

SURFACE COMPOSITION BY ANALYSIS OF IMPACT RADIATION 81

Clark W. White,

Bell Laboratories, Murray Hill, NJ 07974

NUCLEAR RESONANCE AND BACKS CATTERING SURFACE ANALYSIS OF SILICON AND RELATED

INSULATORS 95 Kenneth L. Dunning, Radiation Technology Division

Naval Research Laboratory, Washington, DC 20375
TABLE OF CONTENTS

PAGE

ELECTRON AND PHOTON BEAMS - TECHNIQUES AND APPLICATIONS

APPLICATIONS OF SCANNING AUGER SPECTROSCOPY (SAM) TO THE SILICON INTEGRATED CIRCUIT

(SIC) TECHNOLOGY . 105

Joseph M. Morabito,

Bell Laboratories, Allentown, PA 18103

USE OF AUGER ELECTRON SPECTROSCOPY TO DETERMINE THE STRUCTURE OF SILICON OXIDE

FILMS 119

Jan S. Johannessen and William E. Spicer

Stanford Electronics Laboratories, Stanford University, Stanford, CA 94305

Yale E. Strausser, Varian Associates, Palo Alto, CA 94305

AN AUGER ELECTRON SPECTROSCOPY STUDY OF SILICON SPECTRA FROM SILICON MONOXIDE,

SILICON DIOXIDE AND SILICON NITRIDE 125

Yale E. Strausser, Varian Associates, Palo Alto, CA 94305

Jan S. Johannessen, Stanford Electronics Labs., Stanford Univ., Stanford, CA 94305

SURFACE COMPOSITIONAL CHANGES WITH ELECTRON BOMBARDMENT OBSERVED BY AES 139

Simon Thomas , Semiconductor Analytical Laboratory

Motorola Semiconductor Products Division, Phoenix, AZ 85008

COMBINED SCANNING ELECTRON MICROSCOPY - AUGER SPECTROSCOPY FOR MICRO-SPOT SURFACE

AND IN-DEPTH ANALYSIS OF SILICON AND TRANSISTOR METALLIZATIONS 143

A. Christou, W. Weisenberger and H. M. Day

Naval Research Laboratory, Washington, DC 20375

APPLICATIONS OF X-RAY PH0T0ELECTR0N SPECTROSCOPY (ESCA) TO MIS DEVICES 151

Frank J. Grunthaner, Jet Propulsion Laboratory

California Institute of Technology, Pasadena, CA 91103

CHOOSING BETWEEN ESCA AND AUGER FOR SURFACE ANALYSIS .... 175

Gary E. McGuire , Materials Characterization Laboratory

Texas Instruments, Incorporated, Dallas, TX 75222

SILICON DEVICE APPLICATIONS USING A COMBINED ESCA/AES ANALYSIS SYSTEM 183

L. E. Davis and G. E. Riach

Physical Electronics Industries, Incorporated, Eden Prairie, MN 55343

PHOTODEPOPULATION TECHNIQUE FOR THE STUDY OF ELECTRONIC TRAPS IN INSULATORS .... 189

T. H. DiStefano, IBM Watson Research Center, Yorktown Heights, NY 10598

J. M. Franz, IBM System Products Division, Essex Junction, VT 05452

PHOTOEMISSION AND PHOTOVOLTAIC IMAGING OF SEMICONDUCTOR SURFACES 197

Thomas H. DiStefano

IBM Watson Research Center, Yorktown Heights, NY 10598

ELECTRON BEAM INDUCED IMAGING OF SILICON SURFACES .... 211

William R. Bottoms , Department of Electrical Engineering

Princeton University, Princeton, NJ 08540
TABLE OF CONTENTS

PAGE

CONCLUSION

A COMPARISON OF THE TECHNIQUES FOR SILICON SURFACE ANALYSIS 219

Charles A. Evans, Jr., Materials Research Laboratory

University of Illinois, Urbana, IL 61801

APPENDIX

List of Workshop Participants 233

Author Index 239

Certain commercial materials and equipment are identified in these proceedings in order to adequately specify the experimental procedure. In no case does such identification imply

recommendation or endorsement by the National Bureau of Standards, nor does it imply that

the material or equipment identified is necessarily the best available for the purpose.

vii

WELCOMING REMARKS

Judson C. French, Chief

Electronic Technology Division

National Bureau of Standards

Washington, D. C. 20234

Good morning. It is my pleasure this morning to welcome you on behalf of the National

Bureau of Standards and our cosponsor, the

Defense Advanced Research Projects Agency,

to the Workshop on Surface Analysis for

Silicon Devices.

This Workshop is the fourth in a series

dedicated to the furtherance of the measurement technology needed by the semiconductor

device industry in its attempt to provide to its customers products that are based on the

most advanced technology yet have high re- liability and the affordable costs which result from high yields.

These products are essential components in

modern electronics, and they consequently

play a vital role in the social and economic

welfare of our Nation and in its defense.

They are, therefore, a matter of considerable interest both to the Department of

Commerce and to the Department of Defense,

the parent organizations of the sponsors of

these Workshops.

To those of you who are not in the semiconductor business, and in this Workshop in

particular there are many such people, it

may come as a surprise that there is any

need for NBS and ARPA to carry on work in a field which is so technologically sophisticated and so innovative. The reason is that

the sophistication and innovative abilities

of the semiconductor industry have led to the development of new processes' and new

devices much faster than the measurement

techniques for their control and characterization have been developed.

In the fifteen years that our NBS staff has

worked with the semiconductor industry and

its customers we have seen increasing need

for improvements in practical methods of

measurement for analysis, control, and

specifications in this field. And we have

learned that the Bureau can be especially

helpful in this field because of its neutrality in evaluating measurement methods

and associated technology, and because its

charter encourages it to work in the area of

generic measurement for industry-wide use

and market-place application. This is an area where individual companies understandably find less incentive for extensive re- search than in areas leading to new and proprietary processes and designs.

As a result, the NBS Semiconductor Technology

Program has been established, having as its

goal the development and standardization of

improved methods of measurement for use in

specifying materials and devices and in

control of device fabrication processes:

methods that have been well documented and

tested for technical adequacy, are of demonstrated precision of an industrially acceptable level, and are acceptable to both users

and suppliers.

When such methods are used by the electronics

industry, they are expected to provide a more

consistent set of measured results and interpretations and, hence, lead to improved

quality control and yield in the manufacturer's

plant, and to improved reliability and economy

in the customer's applications.

In recent years, ARPA has joined with us as a major sponsor of the program in order to provide a new approach to the solution of DoD

problems in component reliability and availability, system costs, and system maintainability.

Our program now encompasses work on selected

measurements ranging from those needed to characterize process materials; through those

for photolithography, process control using

test structures, bonding and die attachment,

and hermeticity; on to thermal and electrical

properties of finished devices.

Modern devices are dependent for their operation on the properties of extremely thin

layers of silicon and oxides, and their

interfaces. Performance demands, including

those of reliability and radiation hardness,

call for knowledge and manufacturing control

of the chemical and physical makeup of these

structures that pose requirements for sensitivity and spatial resolution far exceeding

those provided by traditional analytical

methods

.

Thus, an exciting new area of interest in our Program is the subject of this Workshop:

the determination of the present qualitative

1
and quantitative capabilities, and the

future prospects, of modern analytical beam

techniques as applied to the analysis of

silicon and associated insulator films and

device structures.

Of particular interest are the determination

of impurity profiles, surface contamination,

and interface characteristics, using electron, ion, neutral particle, or photon

beams

Our initial work in this area has disclosed

both great promise, and numerous problems in

interpretation and quantification of the

results of beam analysis methods. It has

also disclosed that too little communication

exists between many in the semiconductor

industry who have need for the new techniques, the many expert analysts, and manufacturers of the analytical instruments.

It is the purpose of this Workshop to bring

some of these people together to foster

discussions between these three groups for

the mutual benefit of all three.

I share the pleasure of your Chairman,

George Lieberman, at seeing the large

attendance of representatives of these

groups. I hope that the report of the

Workshop, when it is published, will carry

your discussions to the many others who were

unable to attend. The titles of the papers,

and the stature of our speakers, promise

most interesting and valuable presentations,

and I encourage you to take advantage of the

expertise of our speakers and of your colleagues in the audience in the discussions.

I would like to acknowledge the efforts of our Workshop Committee in making the

arrangements for today's meeting:

Dr. Lieberman, its chairman; Kathy Leedy and

Harry Schafft, for coordination and publicity; and Sara Torrence, for the local ar- rangements .

I want to thank all of you, and especially

our speakers, for attending and helping us

to make this a useful Workshop. And now I will return the microphone and the conduct

of the day's program to Dr. Lieberman.

Again, thank you all.

2
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

INTRODUCTORY CONCEPTS FOR SILICON SURFACE ANALYSIS

A. George Lieberman

Electronic Technology Division

National Bureau of Standards

Washington, D. C. 20234

The planar silicon technology for semiconductor device manufacturing was first described in 1960. Since then, planar silicon

technology has become the principal method

for fabricating semiconductor devices and

integrated circuits (ICs) . The technique

has now been developed to the extent that

3,000 to 10,000 MOS components can be manufactured on chip areas which only fifteen

years before could hold no more than a dozen

bipolar components. To meet today's component density requirements, on chips barely

millimeters on a side, design rules have

been met that call for micrometer linewidths

and gate oxides less than 0.1 pm thick. In addition, most of the device physics is found to occur within nanometers from the

silicon surface. The continuing trend to- ward larger scales of integration and microminiaturization has consequently increased

the need for quantitative. measurements in extremely shallow multilayer device structures and accounts for the growing interest

in surface analysis for silicon devices.

Over the years, manufacturing techniques for control of the IC process have not substantially changed. Yet, it has been estimated

that only three percent of the silicon

entering the manufacturing process now ends

in acceptable devices. New techniques are clearly called for to improve production

yield and assure the quality of completed

devices. Modern surface analysis can pro- vide the IC manufacturer with the potential

to identify, count, and locate, with monolayer precision, trace impurity atoms. For

the device scientist, these analytical

methods also allow determinations of chemical bonding, carrier trapping levels, and

densities of states information. Several of the surface analysis techniques are already

in successful use for IC failure analysis.

From among the myriad of surface spectroscopies which have been spawned over the past

decade, a half dozen of the most promising

techniques for the analysis of silicon device structures and associated materials

were selected for workshop discussion. The

determination of impurity profiles, surface

contamination, and interface characteristics

for IC process control were held foremost in mind. The instrumentation had also to be

commercially available or capable of readily

being constructed.

Referring to figure 1 we see that each spectroscopy employs a beam of primary particles

to probe and interact with the impurity atoms

on the device structure. A flux of secondary

particles is detected which may, or may not,

be the same as the primary particles. The

difficulty in detecting neutrals has been reason to exclude this species as a viable

secondary particle spectroscopy. The ion

and neutral excitation spectroscopies on the right half of the diagram form a group distinct from the electron detection spec- troscopies on the left. Of course, there

are other spectroscopies, but they do not appear to meet the objectives of the workshop .

The ion and neutral excitation spectroscopies may be ranked in energy. At the low

end, between 0.5 and 2 keV, is Ion Scattering

Spectroscopy (ISS) , while at the high end,

from 0.5 to 3 MeV is Rutherford Backscattering Spectroscopy (RBS) . Both spectroscopies measure the energy loss of the pri- mary ion after it experiences an elastic

binary collision with an impurity atom. The

energy loss serves to identify the impurity atom through its mass. Because of its low energy, ISS is extremely surface sensitive. A most important aspect of RBS is that it is the only quantitative and essentially non- destructive spectroscopy for depth profiling

available today.

RBS faces problems detecting light elements

in heavier substrates, but here the complementary accelerator technique known as Nuclear Resonance Profiling (NRP) is of

aid. The chief disadvantage of all these

accelerator techniques is the cost and size

of the equipment. However the increasing

usage of ion accelerators for impurity implantation in device structures may eventually offset this.

At intermediate beam energies, from 1 to 20 keV, the impact of inert or reactive ions

is effective in removing and ionizing surface material in a controllable manner. The

method of Secondary Ion Mass Spectroscopy

(SIMS) consists of determining the mass to

3
charge ratio of these sputtered ions. SIMS

is the most sensitive of all the depth profiling methods, but it suffers from orders

of magnitude variation in sensitivity from

one element to the next.

SCANIIR, an acronym for Surface Composition

by Analysis of Neutral and Ion Impact Radiat ion , is the optical analogue of SIMS. SCANIIR utilizes the characteristic radiative relaxations of sputtered metastable

ions to identify and quantify the amount of an element present on an investigated surface. The major advantage of this technique

is its ability to use neutral beams to analyze insulating films under field-free

conditions

Auger Electron Spectroscopy , AES, is the

most popular of the surface analysis methods.

In conjunction with simultaneous sputteretching it can be used to determine compositional depth profiles. Impurity elements

are identified by the characteristic energies of their detected Auger electrons.

Sensitivities vary by less than a factor of ten for the various elements, except for

hydrogen and helium which cannot be detected.

In addition to having superb depth resolution, AES is also capable of excellent

lateral resolution in the plane of the

specimen, e.g., when the focussed beam of

a scanning electron microscope is used as the primary excitation.

The UV and X-ray Photoelectron Spectroscopies (UPS and XPS*) are becoming increasingly important for electronic materials analysis. UPS has been applied to study the effects of contaminants on surface potentials, and to study the trapping

levels in the energy bandgap of insulators.

Since UV photoelectrons originate from

valence band states, or thereabouts, they are not element specific. On the other

hand, the core level binding energies

measured by XPS not only uniquely identify

the parent atoms, but also contain information concerning its chemical environment.

The chemical information contained in the

XPS spectra is what makes this technique

so useful.

Statements comparing sensitivities, detection limits, resolutions, sampling depths,

etc., of the various spectroscopies must be

tied to a particular material, instrument,

measurement procedure, and the skill of the

*Also known as Electron Spectroscopy for

Chemical Analysis (ESCA)

experimentalist. There is also a certain

transmutability of the instrument parameters

themselves, e.g., the ability to trade sen- sitivity for spectral resolution. It may be

well to consider the definitions of sensitivity and detection limit used in evaluating a spectroscopy. These definitions are illustrated in figure 2 for a typical impurity

calibration curve.

The sensitivity of an analytical technique

refers to the slope of a calibration curve,

i.e., the change in the output signal to the

increment in material concentration. It is, so to speak, the gain of the instrument. By

the output signal is meant the amplitude of

the spectral peak less the background level.

The background consists of all unwanted information which is correlated with the

desired signal, and it is different from

noise which is not correlated and, in principle, can be reduced to any desired level

by averaging the signal over an ensemble of

identical specimens or over a sufficently

long period of time. The data acquisition

time is, of course, limited in practice by

the rate at which the surface deteriorates

under the action of the probe or by contamination. Even under ultrahigh vacuum

conditions, at pressures of 10~ 8 pascals

(vlO- ^ Torr) , a surface can become com- pletely covered by deposited contamination

in less than two hours

The detection limit is a concept which is

related to the sensitivity. The detection

limit is defined as the minimum concentration of an impurity required to produce an observable signal with 95% confidence. The

criterion in common use for whether a trace

detection has occurred is that the signal

count N must exceed the background count by three times the square root of the background count (N > N + 3v/ R~) . D D The detection limit of a surface analytical

technique for a specified element is usually

stated as being a fraction of a monolayer . A surface measurement should ideally sample

only the very outermost layer of surface

atoms, a depth of less than one nanometer.

Silicon has a bulk atomic density of

5 x 10 22 cm

-3 and therefore a monolayer of

coverage would contain about 10^ atoms per

square centimeter of surface. Most surface

measurements today sample depths of several

nanometers with the result that impurity

atoms residing below the surface, but within

the sampling volume, are interpreted as belonging to the surface. Consequently, the

detection limit for a surface analysis may

appear to be either better or worse than it

4
actually is, depending upon the technique,

the sampling depth, and the particular impurity distribution. When a surface analytical technique is applied, together with

some form of surface removal, for the purpose of measuring an impurity concentration - depth profile, it becomes appropriate to introduce bulk detection limits and to speak

of volumetric concentrations. It follows

that those surface analytical techniques

employing the smallest sampling depths are capable of yielding concentration profiles

displaying the greatest depth resolution and

possibly the poorest bulk detection limits.

These concepts are not, of course, the only

factors to be considered in evaluating a spectroscopy. One has only to examine a spectrometer manufacturer's specification

sheet to realize how complex the choice of a technique or an instrument can be.

In summary, the evolution of high component

density, planar silicon technology has

forced us to look for new techniques to assist or replace the traditional process

control methods. Microprobes of extreme

sensitivity are being called for which can

identify, quantify and locate impurity atoms

at surfaces, interfaces and within thin

films. Only a few commercially available

surface analysis techniques have demonstrated

their usefulness for studying silicon device

structures and related material properties.

These techniques will be examined at this

workshop in light of their performance for

analyzing silicon, silicon dioxide, and re- lated materials.

1
PRIMARY BEAM (Probe SECONDARY FLUX [Emission]

DEVICE SURFACE V.' '.'/// '/ '-'.v.'//. I

HEAT, WWFIELDS, ETC.

PROBE

EMISSION i- PHOTON

• ELECTRON

ION

1 PHOTON

1 ELECTRON1 ION1 NEUTRAL

SCANIIR SCANIIR

UPS

XPS|ESCA| AES

SIMS, ISS, RBS, NRP

Figure 1. The analysis of a surface by photon,

electron, ion or neutral beams. The surface

may be stressed to enhance the spectral information carried by the secondary emission. The

primary and secondary particles need not be

identical. A description of the techniques

appears in the text.

Figure 2. Definitions of sensitivity, detection limit, minimum detectable impurity level

and background are illustrated for a typical

impurity calibration curve.

MATERIAL CONCENTRATION ,C Ld

L 0

= DETECTION LIMIT Nmin = MINIMUM DETECTABLE IMPURITY COUNT N R

= BACKGROUND COUNT

6
NBS Special Publication 400-23, ARPA/N3S Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg , Maryland, April 23-24, 1975 (Issued March 1976)

IDENTIFICATION OF INTEGRATED CIRCUIT PROCESS AREAS AMENABLE TO DIAGNOSIS AND CONTROL BY ANALYTICAL

BEAM TECHNIQUES

Bruce E. Deal

Fairchild Camera and Instrument Corporation

Palo Alto, California 94304

ABSTRACT

The general process steps required to fabricate silicon integrated

circuits are reviewed, and examples in each area are indicated where

analytical beam techniques can be used for compositional analysis and

process control. In some cases satisfactory methods of analysis have

been established and are in current use, while in others new techniques

need to be developed. Those characterizations related to device reliability and yield improvement are considered to be the most important;

these include metallization, chip protection, and oxide charge control

for reliability effects, and photomasking and cleaning in regards to yield considerations.

INTRODUCTION

Analytical beam techniques have been employed

for a number of years to evaluate and control

processes used to fabricate semiconductor devices and integrated circuits. As the structures of these devices and circuits have

become more and more complex with increasing

component densities, the requirements for more refined analytical techniques have increased accordingly. The purpose of this paper is to review areas of integrated circuit

fabrication where such analytical beam

techniques of analysis are important.

There are two general types of integrated

circuits —MOS and bipolar. Examples of their structures are shown in Fig. 1. While specific process steps used to fabricate these

circuits vary widely from product line to product line, some general process areas are common to all MOS or bipolar structures.

These are as follows

:

Silicon Substrate Preparation

Thermal Oxidation

Photomasking and Cleaning

Diffusion

Dielectric Deposition

Metallization

Chip Protection

In this paper we will discuss each of these

general areas, and will indicate where problems in composition or structure may arise

which can affect device properties or performance. We will give examples which show how

some of these problems are being solved by

the proper use of analytical beam techniques,

and where, in other cases, additional work

needs to be done. Due to space limitations,

only one or two examples will be given for each process area.

Before discussing these possible problem

areas, we should emphasize that many of to- day's bipolar integrated circuit structures

are every bit as complicated as MOS-type circuits, and we can no longer say that MOS process control problems are more severe than

those of bipolar circuits. Also, the formidable requirements for tools used in the

analysis of integrated circuit structures

produced today are demonstrated by the yearly

increase of circuit complexity as shown in

Fig. 2. These data were produced by Gordon Moore [1] and show that in 1975 the semiconductor industry can produce circuits containing more than 10,000 individual components on chips with areas not too much larger than

those of discrete devices produced only fifteen years ago. Thus, the challenge to those

involved with developing and carrying out materials analysis techniques is readily

apparent

.

DISCUSSION OF INTEGRATED CIRCUIT PROCESSING

AREAS

(A) Silicon Substrate Preparation

Characterization and analysis of semiconductor

substrates is an established procedure and includes several areas. The first is the original crystal as produced by various crystal

growing techniques. Such properties as structure, impurities, dopant distribution and

defects such as dislocations must be

7
characterized and kept under suitable

control. Next, the silicon surface after

wafer slicing must be maintained free of con- tamination and particles. In addition, the

surface should have the proper crystal orientation and be relatively free of stress. Finally, subsequent substrate processing

steps such as buried collector preparation

and epitaxial film deposition must be controlled. In the latter case, all considerations applicable to the original crystal bulk

and surface apply, but in addition, analysis

and control of the dopant profile become

increasingly important.

As indicated above, methods of analysis used

to characterize the properties of the silicon

substrate are well known. Of course, as larger slices are used and more complex circuit structures are developed, the analytical

techniques may have to be refined. Since

this area of integrated circuit fabrication

is probably under better control than the

others, no specific examples will be presented here. However, the reader is referred to

a general reference on the subject of surface

characterization [2].

(B) Thermal Oxidation

Thermal silicon dioxide (Si02) is the basis

for today's semiconductor devices and integrated circuits. It provides passivation for

active device regions that is not equaled by

other types of dielectric layers generally

prepared by deposition techniques. The thermal oxidation kinetics and properties have

been characterized in a satisfactory manner

so that devices with predictable properties

may be produced. Included in this characterization are the electrical properties as determined by the four types of oxide

charges . These charges are Qss , the fixed

oxide charge; Nst , fast interface states;

Q0 , mobile ionic charge; and Not , radiationinduced charge. While the empirical dependence of these charges on processing

variables is well known, a number of questions remain as to their exact origin and

physical nature [3]. Indications of how

these charges may be investigated using

analytical beam techniques are as follows.

Figure 3 depicts a cross-section of a thermal

silicon oxide as oxidation is proceeding.

Also included in the figure is the general

relationship for the thermal oxidation process. It can be noted that a thin oxide re- gion at the Si02~Si interface is shown to contain partially ionized silicon. It is believed that the ionized or excess silicon in

this region is the origin of Qss , the fixed

oxide charge. However, investigators to date

have not verified this origin and additional

work needs to be done. An example of the

type of analysis that can be employed is

shown in Fig. 4. Here, an analysis of the

Si/0 ratio in a thermal oxide has been obtained using the ion backscattering technique

[4]. Although it is not clear in the figure,

the data indicate that a thin layer of about

three atomic layers is present which contains

a silicon rich oxide. This is compatible with

the Qss origin model described above. However,

it is clear that additional characterizations

and analyses are required to provide a satisfactory understanding of this phenomenon.

The same questions arise regarding origin of Nst , the fast interface states. These states

or charges may have a similar origin to Qss , but since they are in electrical communication with the silicon, they can respond to changes in surface potential. While much is

known about Qss and Nst , the question of their

origin is still quite thought-provoking and

presents a challenge to those involved with

beam analysis

Similar questions may be asked about the

physical nature of QQ

, mobile impurity ions.

These ions, which include sodium, lithium,

potassium, or hydrogen protons, have long

been a source of instabilities in semiconductor device fabrication. A recent development

has been the incorporation of a chlorine species in the oxidation process. This chlorine

has been reported to "cure all ills" in areas

of ionic instability, dielectric breakdown,

junction leakage, fast state generation, and

others. An example of a verified improvement is shown in Fig. 5 , where MOS device

drift is almost eliminated in two cases of

sodium contamination [5]. Auger analysis is

being used to determine some of the compositional properties of these chlorine-containing

oxides as shown in Fig. 6, as reported by Chou

et al [6] . However, many additional questions

remain concerning ion instabilities, in general, and even effects of the analysis technique

on the impurity profiles themselves have been

noted. Thus, another challenge is presented

to materials analysis scientists.

The identification of the origin and

elimination of Not , the radiation induced

charge is no exception to the discussion presented above. Much time and money have been

expended on this subject over the past ten or more years, and many questions still remain.

In Fig. 7, Gwyn [ 7 ] summarizes some of the

different "cures" for radiation hole trapping

in oxides. These proposed solutions include,

optimizing oxidation conditions, incorporating

species such as aluminum or chromium in the

oxide, or even the use of mixed oxides or
dielectrics. The need for developing

analytical beam techniques for identification

of radiation charge origin and elimination

mechanisms is obvious, if we are to be able

to produce MOS devices that will withstand

radiation environments.

In addition to oxide charges, other areas of

thermal oxidation that need to be controlled

or better characterized are oxide integrity , complex oxide-silicon interfaces and dopant

impurity redistribution . These will be

briefly discussed in later sections.

(C) Photomasking and Cleaning

Of all the areas of integrated circuit

fabrication, none has advanced as rapidly as the photomasking operation. This is demonstrated in Fig. 8, where a single MOS transistor of 1964 is pictured as an insert in a photo of a recently developed calculator

chip. The latter, which contains about

10,000 individual transistors, is still not

as densely packed as more advanced memory

chips. The dimensional scales on the two devices in Fig. 8 are roughly equivalent;

the larger circuit being about 150 x 180

mils

.

The tremendous increase in complexity of

today's circuits results in enormous problems

related to carrying out the photomasking process as well as to the associated cleaning

treatments. The use of analytical beam

techniques for identifying and controlling

contamination left over from these processes

will become increasingly important. Likewise, methods need to be established for

better characterization of the photoresist

materials. In addition to electrical instability problems, the question of oxide

integrity presents itself. In Fig. 9, Kern

has provided an example of how two different

types of resists may affect defect formation

in the underlying oxides [8]. While the incentive for those involved in materials analysis may not be as obvious or the results as satisfying in this area, it is certain that

the effects on device yield and reliability

are more significant than in almost any

other IC processing step.

(D) Diffusion

One of the oldest processes associated with

silicon devices is that of dopant diffusion . Typically in the past, the source of dopants

that are deposited on the silicon surface

prior to diffusion into the silicon has been

a gaseous species transported in a diffusion

tube from solids (P2O5, B2O3), liquids

(POCI3) or gases (BoHg). More recently,

9

other types of sources have been employed,

such as doped vapor deposited dielectrics,

spin-on oxides and ion implantation. In all

cases, the main problem has been to determine

the diffusion profile before, and more important, after the subsequent high temperature

diffusion steps. Also, it has been important

to know what percentage of the doping species

is electrically active. Thus, many analytical

techniques have been used for obtaining these

profiles, as well as to determine anomalies

and structural damage to the silicon. An example showing the importance of being able

to determine dopant profiles is shown in Fig.

10, where the depth of a dopant's penetration

during ion implantation is determined by the

type of dielectric on the silicon surface

[9].

Closely related to the diffusion process and

requiring similar types of profile analysis

are impurity gettering , and redistribution

of dopants during oxidation. In the case

of gettering, unwanted metallic impurities,

such as copper, iron, nickel, and the like are removed by a complexing action at high

temperatures with a species such as phosphosilicate glass. A knowledge of the distribution of these impurities before and after

gettering is important. Likewise, the changing dopant profiles in the silicon due to redistribution during thermal oxidation must

be known in order to end up with the proper

device electrical properties. It has been

difficult to accurately predict and verify

redistribution profiles after just one oxidation treatment. When it is considered that

today's integrated circuit processes employ

many such treatments, the correct prediction

and/or determination of dopant profiles is

almost impossible. In Fig. 11, an example

is shown where attempts are being made to calculate and determine profiles after just two oxidation steps [10]. It is obvious

that much more work, both theoretical and

analytical, needs to be done in this area.

Finally all of the above considerations

discussed thus far must be put together and

then the problems multiplied by some unknown

factor when a three-dimensional aspect is

considered. This is exemplified by the

drawing in Fig. 12, which depicts a device

structure employing an oxide or Isoplanar

type of isolation. The interface between

this oxide and the device silicon now can result in changes in orientation-dependent,

oxide charge densities, as well as dopant

impurity concentrations, all of which can severely degrade device properties. Furthermore, the difficulty in predicting these

interface properties is exceeded only by

the difficulty in analyzing them. So it is
obvious that there will be plenty for the

materials analysis scientist to do in co- operation with the device engineer, when

this complex type of structure is considered.

(E) Dielectric Dep osition

Vapor deposited dielectric films are generally used in conjunction with thermal

oxides to provide additional improved device properties. These properties include

impurity ion masking, increased voltage

breakdown, gettering of impurities, and

others. Dielectrics most commonly used for

these applications are silicon nitride, sili- con oxide, aluminum oxide and phosphosilicate

glass. While the use of these deposited dielectrics provides improved characteristics,

some adverse effects may result. One such

effect is an electrical instability due to charge generation. Four general types of

charge formation instabilities have been

found to be associated with the combination

of deposited dielectric films over thermal

silicon oxides. These are indicated in Fig.

13. The degree of all these effects is dependent on the structural composition of the

deposited dielectrics. Thus, the analysis

and control of the film compositions and

structures are very important. An example

of this effect is shown in Fig. 14. Here,

the electrical conductivity of a silicon

nitride film is inversely proportional to the ratio of nitrogen to silicon in the film,

the latter being a function of the reactant

gas ratios during film formation [11]. The

conductivity of the SiaNij thus determines

the degree of instability (c) in Fig. 13,

The etch rate of the nitride film also is a function of its composition. Many other

examples could be given which show the importance of being able to characterize the

compositional and structural properties of

deposited dielectric films, which can lead

to adverse effects on device properties.

(F) Metallization

The metal interconnection system represents

a key aspect of integrated circuit fabrication. This metal has typically been pure

aluminum, but lately aluminum alloys have

been used, as well as inert metals (Pt, Au)

refractory metals (Mo, W) and, in the case

of MOS circuits, polycrystalline silicon.

In the latter case, the films are doped with

phosphorus or boron to achieve the required

conductivity. As integrated circuits have

become more complex, the analysis and control of the metal interconnection systems

have been accordingly more important. This

has especially been the case with the addition of alloying constituents to the

aluminum for improved electromigration

resistance (Cu) and for reducing silicon

substrate alloying effects (Si). An example

of effects on the metal film properties due

to the addition of these elements is shown

in Fig. 15. Here, it is important to analyze and control the precipitate or cluster

formation as well as nodule growth [12]

Another example where structural analysis of

interconnections is important is that of

polycrystalline silicon deposition, as shown

in Fig. 16. In this case, X-ray analysis

has been used to determine relative amounts

of orientation textures which can affect

both electrical properties as well as dopant

diffusion kinetics [13]. It is felt that

satisfactory techniques exist for charactizing these metal films, but that much more work is needed to monitor changes in film

structure and composition during various

heat treatments and even device operation.

In the latter case, reliability aspects be- come important as chemical reactions occur between the metal films and the environment

in the case of nonhermetic packages (see

next Section)

(G) Chip Protection

Very soon after integrated circuits were

developed in the early 1960's, it was discovered that some sort of layer must be deposited over the metallized chip to protect

the metal interconnections from mechanical

damage during subsequent assembly procedures.

Films used for this purpose were plastics,

sedimented glasses, and vapor deposited oxides. Later it was found that this same type

of chip-protecting film was required to prevent the metal from reacting with moisture

and other gases which penetrated plastic

packages. At the same time, it was difficult

to obtain a film which exhibited all the nec- essary properties—physical, chemical and

electrical—to provide the necessary protection. This is true to a large extent at the

present time. One of the most common causes

of reliability problems in plastic packaged

integrated circuits is metal corrosion.

Many of the above problems result from

improper analysis and control of the chip

protection film, which quite often contains

phosphorus. This phosphorus prevents cracking of the deposited silicon oxide film, but

also can react with water to form phosphoric

acid. This acid in turn is a common etchant

for aluminum. Paulson and Kirk [14] have presented some curves showing optimum phosphorus

concentrations for these films, and one of

these is shown in Fig. 17. This example helps

to point out the requirements for the determination of composition and structure of chip
protection films. An additional difficulty

is that of obtaining a proper depth profile

of the constituents.

One last type of characterization that is

very important in this area involves failure

analysis. When integrated circuits are subjected to life tests or fail in actual use,

the understanding of the failure mode is

very important and may be accomplished by a carefully executed analysis. Such an analysis includes identification of reaction

products, and analytical beam techniques can

be used very effectively for this purpose.

CONCLUSIONS

In summary, we have reviewed the various

processing steps of integrated circuit fabrication. In doing so, we have noted the many

areas where analytical beam techniques can be used to provide a better control of these

processes and the resulting device properties. Many such analyses are being successfully used today, while in some areas, new

or improved methods must be developed. If priorities are to be stated, those areas of

analysis related to device reliability

should be listed first. These include chip

protection, metallization and control of

device electrical properties. Next in importance would be wafer sort yield. In this

case, control and characterization of the

photomasking and cleaning processes would

undoubtedly be the most important factors,

since this process area dominates yields of

semiconductor integrated circuits today.

Finally, it is important that the scientists

who are responsible for developing and carrying out the analyses keep in close communication with the IC processing engineers, so that the problems that need to be solved are

solved, and that our always limited resources are used most efficiently.

ACKNOWLEDGMENTS

The author wishes to thank A. J. Learn and

R. C. McDonald for helpful discussions

concerning various aspects of this paper.

REFERENCES

1. Moore, G. E. , Private Communication.

2. Characterization of Solid Surfaces,

P. F. Kane and G. B. Larrabee, Eds.,

(Plenum Press, New York 1974).

3. Deal, B. E. , J. Electrochem. Soc. 121 , 1986 (1974).

4. Sigmon, T. W. , Chu, W. K. , Lugujjo, E. and Mayer, J. W. , Appl. Phys. Lett. 24 , 105 (1974).

5. Kriegler, R. J. and Devenyi, T. F., Proa. 11th Annual Reliability Physics

Conf., Las Vegas, Nevada, April 3-5, 1973.

6. Chou, N. J., Osburn, C. M. , van der

Meulen, Y. J., and Hammer, R. , Appl.

Phys. Lett. 22, 380 (1973).

7. Gwyn, C. W. , Sandia Laboratories Report

No. SLA-73-0017, Jan. 1973.

8. Kern, W. , RCA Review 34, 655 (1973).

9. Reddi, V. G. K. and Yu, A. Y. C. , Solid

State Tech. , 35, (Oct. 1972).

10. Margalit, S. , Neugroschel, A. and

Bar-lev, IEEE Trans. Elec. Dev. ED-19 , 861 (1972).

11. Gyulai, J., Meyer, 0., Mayer, J. W. and

Rodriguez, V., J. Appl. Phys. 42_, 451

(1971).

12. Learn, A. J., Thin Solid Films 20, 261

(1974)

.

13. Kamins, T. I. and Cass, T. R. , Thin

Solid Films 16, 147 (1973).

14. Paulson, W. M. and Kirk, R. W. , Proc. 12th Annual Reliability Phys. Conf. Las Vegas, Nevada, April 2-4, 1974.
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13
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19
DISCUSSION OF THE PAPER

Lieberman : I am particularly interested in

the high energy backscattering data presented in your fourth slide for the siliconsilicon dioxide interface. You suggested

that the O/Si stoichiometric offset was due

to a silicon-rich transition layer several

atomic layers thick. On the basis of this

data alone, could not the excess silicon

region have extended hundreds of angstroms

into the oxide film.

Deal : I do not think it extends that far.

Most of the change effects that we observe

probably result from a region within 50

angstroms of this interface.

independent of the electronic screening and

so this technique cannot be used to determine the ionization state of the excess

silicon.

DiStefano : We, at IBM, have used field

dependent internal photoemission to probe

that interface down to 4 angstroms from the

silicon and find no evidence of oxide non- stochiometry of large enough extent to perturb the conduction band of the Si025 i.e.,

the conduction band of the Si02 appears to be flat to within 0.02 eV up to 4 angstroms

of the silicon surface. There may be a small amount of excess silicon but we cannot

find any ionized silicon of consequence.

Deal : This is for a thermal oxide?

Lieberman: Another point is that the cross section for Rutherford backscattering is DiStef ano : Yes.

20
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

LOW ENERGY ION SCATTERING SPECTROMETRY STUDIES

OF Si, Si02, AND RELATED MATERIALS

William L. Harrington

RCA Laboratories

Princeton, New Jersey 08540

INTRODUCTION

Although low energy ion scattering spectrometry (ISS) has been described by several

authors as a surface and in-depth analytical

tool 1

' 2 ' 3

, characteristics of the method will

be illustrated in this paper by using analytical examples relating to the topic of this

workshop—Si, Si02, and related materials.

After consideration of the technique in general, some fundamental studies of Si and the

Si/Si02 interface will be discussed in which

ISS is particularly suited.

METHOD

Figure 1 is a schematic diagram of the scat- tering process in which rare gas ions of mass M^ and initial energy Eq are scattered

by surface atoms of mass M2 in elastic, bi- nary collisions. Although this ion-surface

interaction results in the neutralization of about 99.9 percent of the incoming ions, a significant, measurable number survive as scattered ions, and the energy loss suffered

by these ions can be related to the mass of the scattering center by the equation given

at the bottom of the figure. At a lab scattering angle of 90° this equation reduces to the simple relation

Ei/E0

= (M 2 - M!)/M2 + MX ) where E]_ is the energy of the scattered ion

and the other terms have been defined above.

Since Mj and Eo are known quantities and one measures E^ in the experiment, it is a simple matter to deduce M2 . The rare gas ions

used are generally in the 1-3 keV energy

range and will thus penetrate most solids on

o the order of 10-60 A. Survival of these

projectiles as scattered ions from a single

collision, however, can be restricted almost

entirely to the first average atomic layer

and thus defines ISS as a true surface technique. Of course, bombarding a solid with

such energetic ions also causes sputtering,

and for this reason the technique naturally

provides information as a function of depth.

To better put this technique in perspective

with related methods, Figure 2 contrasts the

principles just described with those of high

energy backscattering and secondary ion mass spectroscopy (SIMS) . High energy ISS uses

MeV projectiles which

o penetrate thousands to tens of thousands of A into a solid, and the energy loss suffered by the scattered particle (this particle does not have to remain

ionized to be energy analyzed at these energies) is due not only to the mass of the

scattering center, but its depth as well.

It should be noted that this technique does

not depend on the sputtering process for obtaining information in depth, and can thus

be used very effectively to calibrate difficult, layered samples. SIMS, on the other

hand, detects not the scattered ions, but

those sputtered particles which escape the

solid as ions. The sputtered ions are then

extracted into a mass spectrometer and analyzed according to their mass/charge ratio. A point sometimes forgotten when comparing

surface techniques is that SIMS data come from the sputtered species, which in the

steady state are representative of the bulk

sample composition. Low energy ISS and other

techniques such as Auger Electron Spectroscopy analyze the surface that remains after

sputtering, so that differential sputtering

effects must be considered in the interpretation of data.

The instrumentation which in practice detects

the low energy scattering event is commer- cially available from the 3M Company 1*' 5 and

is schematically shown in Figure 3. Rare

gas ions are formed in a filament ion source

and accelerated toward the sample surface at energies between 500 and 3000 eV with a very

small energy spread. The ions are focussed

to a spot which is variable from about 1-3 mm in diameter (recent enhancements of the

ion gun have reduced the beam size to 100 pm)

This beam can be kept stationary for highest

sensitivity with relatively poor depth reso- lution or rastered in combination with a small area detection gate for optimum reso- lution in depth. Ions that scatter from the sample surface at 90° enter a 127° electric sector for energy analysis and are detected

with a channel electron multiplier operating

in a pulse counting mode. A feature of great importance in this instrumentation is

a charge neutralization filament located

near the sample surface. Insulators, such

as glass or thermal layers of Si02, which

can be a problem for techniques employing

high density ion or electron beams, are

21
flooded with thermal electrons during ion

bombardment with feedback regulation to keep

the sample surface at a fixed potential near

ground. The importance of sputtering with an uncharged surface will be demonstrated in some of the examples to follow, but the crucial point is to eliminate large electric

field gradients across a sample to prevent

migration of mobile ions.

The output from this technique is best described with the help of a pertinent example

as shown in Figure 4. These are typical ISS spectra where the scattered response in

counts/second is plotted vs. E^/Eq, the

ratio of scattered to initial ion energy as defined previously in Figure 1. The sample

in this example is an alkali strontium silicate glass which is certainly a good insulator. With charge neutralization by the flood

filament, however, no charging effects are noted. That is, the peaks fall at the same energy ratio as those obtained from a good

conductor, and there is no evidence that mo- bile ions are affected by the sputtering

beam. The top spectrum in this figure is

that for a normal, untreated glass, and one

can discern peaks for 0, Na, Si, K, and Sr. The bottom spectrum is from the same glass

after the application of a heat and voltage

treatment with non-blocking contacts to move mobile ions into the glasses 7 —only 0 and Si

remain at the surface. A series of such

spectra while continuing to sputter yields a concentration profile in depth as shown in

Figure 5. Note the depletion and subsequent

pile-up of Sr, the largest and least mobile

of the ions, whereas Na, the most mobile

species, is barely detected even at large

distances from the surface. Potassium, having an intermediate mobility, shows an intermediate trend.

The magnitude of the scattered signal from an element in a given matrix is basically a function of three variables—(1) the geometric arrangement of atoms on a surface, i.e.,

does one atom shield another from the incoming beam?, (2) the extent of neutralization

of the incident ions, and (3) the scattering

cross section (the probability that scattering will occur) as a function of atomic number. Geometry, for most analytical problems

concerning amorphous, non-oriented layers,

is not critical; but for situations where

the position of atoms is important, the one

layer resolution of ISS can be used to an advantage. Although neutralization may vary

from one surface to another, differential

neutralization between elements in a given

matrix would be expected to be a rather

small factor. Relative scattering cross sections have been studied extensively, and

certainly as a first approximation, one can correct elemental scattering data for this

variable. Figure 6 is a plot of Bingham's

differential cross sections 8 as a function

of atomic number for 1500 eV4 He +ions . Of course, there is a family of such curves for

each initial energy and for each different

probe ion. The magnitude of the correction

implied by these data over the range of a light element such as 0 to a heavy element

such as Pb is only about a factor of 6 as compared to several orders of magnitude. cor- rection for ion yields in the SIMS technique.

To illustrate this correction, Figure 7 shows

a spectrum of glass used to form a liquid

crystal cell. When a voltage is applied to specific areas of such a cell, the liquid

crystal material sandwiched between the glass

surfaces will align to produce visible contrast for displays. For this application

the glass surface composition has been found

to be important, and in some cases leads to liquid crystal misalignment. The spectrum

shown in this figure was taken in an area of

the cell which operated normally. The Na,

Ca, 0, and Si concentrations were calculated

by subtracting a background for each peak

and dividing the net peak height by the relative scattering cross section. A simple

summing of all corrected signals and divisions to determine individual percentages

produced very good agreement with the bulk

composition as given by the manufacturer.

Figure 8 is a similar spectrum taken in an area of the glass cell just a few mm away

from that in Figure 7. In this case liquid

crystal misalignment was a problem, and the

surface composition of the glass, as determined from the same calculation procedure

described previously, is completely different

with very high concentrations of Na and Ca.

CRYSTALLOGRAPHIC ORIENTATIONS OF SINGLE

CRYSTAL Si

Having established the characteristics of

the method, it is instructive to describe

the results of two ISS studies which focus on Si and the Si/Si02 interface.

The (111) and (100) surfaces of single crystal Si wafers are used extensively in Si devices and offer an excellent system for

studying surface sensitivity and information

depth of an analytical technique. In addition, for those techniques using sputtering

to provide information as a function of

depth, single crystal Si provides a sensitive means of evaluating damage to these surfaces

from ion bombardment.

The detailed preparation of these surfaces

22
has been described previously , but Figure 9 shows spectra of a typical surface as re- ceived—there is the expected small amount

of residual oxide, some carbon, and F from

the final HF etch. In order to obtain accurate ISS response on the pure Si surface, it was necessary to carefully sputter through

this thin contamination layer to reach a clean Si surface, as shown in the right half

of this figure.

If one considers the geometrical packing of

Si atoms on the (111) and (100) surfaces,

the surface atom density of (111) Si is 1.15

times greater than that of (100) Si. ISS responses for these two orientations are

shown in Figure 10 as a function of several

energies for both He"*" and Ne+ . The columns

under the figures show the signal level for

each orientation, and the last column is the

calculated (111)/(100) ratio. For He+ the

observed ratio does not fall within the

range of the true ratio until an energy of

1500-2500 eV is reached. Because the true ratio is observed at these higher energies,

one can deduce that no detectable damage has

been done to the crystal lattice by He bombardment. At low energies the ratio ap- proaches unity because the low current density and very low sputter yield at these en- ergies result in a situation where the surface is not kept clean during the scattering

measurement. With such a contamination layer

on the surfaces, the true Si response cannot

be obtained. If the surface were first

cleaned by sputtering with 1500 eV He+ , the

ratios obtained with Ne+ at 500 and 1500 eV

are correct within the measurement precision.

The slightly lower value at 500 eV could

again be due to contamination since the

higher Ne sputter yield was somewhat offset

by a low current density of about 3 pA/ cm2

. The low ratio at 2500 eV indicates that damage has been done to the crystal lattice.

This damage is shown more clearly in Figure

11, where Ne+ at 500, 1500, and 2500 eV was

used to sputter the two surfaces until the

oxide was removed. Each location sputtered

was then examined with 1500 eV He"*" ions

.

The ratio at 500 eV is the same as that obtained with no Ne+ sputtering, so that again

there was no detectable damage. The ratios

at 1500 and 2500 eV, however, become progressively smaller; and if one examines the individual Si responses, it is the (111) surface that remains stable and the (100) surface that becomes damaged and pushed into a configuration of higher surface density.

Even at these low energies lattice damage

can result —especially in Si lattices of

less than maximum density. Those techniques

which obtain depth profile data by sputtering

with heavy ions (0, Ne, Ar) at much higher

energies are definitely causing damage to the surface being examined, and the possible

effect of this damage on true concentration

profiles must be kept in mind.

STOICHIOMETRY OF SiO? AND THE Si02/Si

INTERFACE

The Si - Si02 system is a vital part of

present semiconductor device technology as evidenced by the fact that this workshop has

been organized on this topic. The thin re- gion making up the Si/Si02 interface dominates many aspects of MOS behavior—surface

states, surface charge, radiation damage,

dielectric breakdown, etc. —and for this reason has received a great deal of attention

by analytical techniques. High energy backscattering, in particular the work of Sigmon,

Chu, Lugujjo, and Mayer 10

, has reported a rapidly decreasing 0/Si ratio as one measures thinner and thinner oxide films. This

work is interpreted as showing excess Si near the interface which has been pulled out

of the substrate lattice into the oxide during the initial stages of oxide growth. However, the limit of depth resolution by this

technique is on the order of 200 A, so that

only integrated values over these thicknesses

can be obtained, and thus the data can only

infer the Si-rich interface.

The low energy scattering technique has such

a well-defined, small information depth as illustrated previously in this presentation

that the question of Si02 stoichiometry and

the character of the Si02/Si interface can now be examined in much greater detail.

Assuming that excess Si does exist at the

interface, it is inconceivable that these

substrate atoms could move large distances

into an oxide. For this reason an oxide

thickness of about 1000 A was chosen as an

ISS standard for Si02- Certainly, in the

outer few hundred A of such a film, as shown

in Figure 12, the 0/Si ratio should be 2:1. In this and all the following figures the

scattered response for 0 was multiplied by

1.72 as an empirical, relative scattering

cross section correction, which for the 904

o . A oxide in Figure 12 produces an 0/Si ratio

of 2. Figure 13 shows a complete profile of

this same film after controllably etching

down to a thickness of 498 A (all these

oxide thicknesses were measured by ellipsometry before ISS profiling) . At this

thickness the 0/Si ratio is still 2 within

the measurement precision, and it might be

noted that at such oxide thicknesses, high

energy ISS already reports a deviation from

the 2:1 ratio. This etching procedure was

23
continued for several steps, and down to 44 A, as shown in Figure 14, the ratio at the sur- face is still 2. It is at this point that

there is a slight indication in the Si re- sponse that might be direct evidence of excess Si at the interface.

Because it is difficult to controllably etch

films much thinner than in Figure 14, the

problem was then approached from the other

direction—clean to bare Si and grow to known, measured thicknesses of 600°C oxide.

At 7 A, in Figure 15, the 0 signal had to be

multiplied by 5 to show up effectively on the same scale as the other plots. At the maxi- mum 0 level the O/Si ratio is only 0.3-0.4,

in good agreement with some of the thinner

oxides measured by high energy ISS. At these

thicknesses the integration effect of the

high energy technique is of little consequence. Figure 16 shows a large change when

growth proceeds from 7 to 17 A*, but in this

experiment the ratio is still between 0.2 and

0.5 in the region of constant oxygen response.

Again at 40 A, as shown in Figure 17, the

0/Si ratio is 2 near the surface and this

profile agrees quite well with the 44 A layer

obtained by thick film growth and etching

(Figure 14)

The results shown in Figure 18 were obtained

after the installation of a beam raster/gated

detector system. Such a modification gives

the much needed improvement in depth resolution that was lacking with a stationary,

Gaussian beam, and almost equally important

allows a fine control over the sputter rate

for obtaining detailed data on these very

thin films. Over much of the film in this

figure the 0/Si ratio is again 2, but at the

interface it is clear that the Si signal

rises to a plateau before the 0 signal falls.

It is this region of 10-20 A that is established experimentally as Si-rich. Using the

Si atom density of 7.9 x 10 11+ at/cm 2 in the

thermal Si02 region as a point of calibration,

one can calculate rather directly that the

Si-rich plateau corresponds to a Si density

of 9.4 x 10 ll+ at/cm2

. Therefore, there is an excess Si density of 1.5 x 10^ 4 at/cm2 per

monolayer which extends out from the interface 10-15 A or about 3-4 molecular layers.

REFERENCES

1. Smith, D. P., Scattering of keV Ions from Solid Surfaces, Proc. 15th Ann.

Conf. on Mass Spectrometry and Allied

Topics, Denver, 1967.

2. Smith, D. P., Scattering of Low Energy

Noble Gas Ions from Metal Surfaces, J. Appl. Phys. 38, 340-347 (1967).

3. Honig, R. E. and Harrington, W. L., Ion

Scattering Spectrometry Below 10 keV,

Thin Solid Films 19, 43-56 (1973)

4. Goff, R. F. , Ion Scattering Spectrometry — A New Technique for Surface

Composition Analysis, Proc. 6th Natl.

Conf. on Electron Probe Analysis,

Pittsburgh, 1971.

5. Goff, R. F. , Ion Scattering Spectrometry, J. Vac. Sci. Technol. W, 355-358

(1973).

6. Carlson, D. E. , Ion Depletion of Glass

at a Blocking Anode: I, Theory and Experimental Results for Alkali Silicate

Glasses, J. Am Ceramic. Soc. 57 (7)

291-294 (1974).

7. Carlson, D. E. , Hang, K. W. and

Stockdale, G. F. , Ion Depletion of

Glass at a Blocking Anode: II, Properties of Ion Depleted Glasses, J. Am. Ceramic Soc. 57(7), 295-300 (1974).

8. Bingham, F. W. , Tabulation of Atomic

Scattering Parameters Calculated Classically from a Screened Coulomb Potential, Sandia Research Rept. SC-RR-66-

506, TID-4500-Physics, 1966.

9. Harrington, W. L. and Honig, R. E. , Low

Energy Ion Scattering Spectrometry of

(111) and (100) Silicon, Proc. 22nd

Ann. Conf. on Mass Spectrometry and Allied Topics, Philadelphia, 1974.

10. Sigmon, T. W. , Chu, W. K. , Lugujjo, E. and Mayer, J. W. , Stoichiometry of Thin

Silicon Oxide Layers of Silicon, Appl.

Phys. Lett. 2A_, 105-107 (1974).

* Recent data on similar films obtained with

the raster/gate system for better depth

resolution indicate an 0/Si ratio of 2 at the surface of films only 15 A thick. This

observation also establishes the extent of

o the Si-rich region as 10-15 A.

24
SCHEMATIC OF ION SCATTERING PROCESS

PRIMARY ION SCATTERED ION

Mi

o o ©40 o o

000000SAMPLE

Figure 1. Schematic diagram and formulae

of the ion scattering process.

OOOOO9

E .'£.-(.-,"y'(-"t, "'"''

,! -'i,!,

]

V! f

E| M2 -M, FOR 90* SCATTERING

E 0 M 2 +M,

Figure 2. Comparison of high energy

backseat tering, low energy ISS and SIMS1

HIGH ENERGY

ISS (MeV) AE

LOW ENERGY

ISS (KeV) SECONDARY IONMASS ANALYSIS

SAMPLE TURRET

Figure 3. Schematic diagram of the 3M ion

scattering spectrometer.

25
Figure 6. Differential scattering cross sections

for 4 He .

21
Figure 7. Relative scattering cross section Figure 8. Calculated ISS composition from a correction applied to ISS data from a "normal" "misaligned" liquid crystal cell,

liquid crystal cell.

4 He + ISOOV^ Si INITIAL SURFACE AFTER SPUTTERING OXIDE

(100)

c/s/nA 500V 1.02 t 7 0%

1500V 10.90 t 0 9%

2500V 14.56 ± 3.3% 500V 9 48 i 3 I %

1500V 11.25 ± 2 7%

2500V 15.74 ± 2.6%

c/s/nA 0 98±3 8%

II 92±0.4%

17 32±4.5%

I0.47±2.6%

I3.28t 1,6%

I6.90i 1.4%

(lll)/(IOO) ATOM DENSITY

RATIO

1.15

0.96±8.0% 09tl.0%

19 t 5.5 %

10 i 4 0%

18 ±3.1 %

07±2 9%

E,/Er E|/E 0

Figure 9. 1500 eV He ISS spectra of a typical Si surface.

Figure 10. ISS Si responses from (111) and

(100) Si surfaces.

27
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28
DISCUSSION OF THE PAPER

Participant : Do you have any measure of the excess silicon versus charge density via

depth?

Harrington : No, I do not know what the

charge density was on these films. This is

very recent work and I am continuing on it.

I am sure we will do measurements like that.

Johanessen : What is the penetration depth

of the helium ions?

Harrington : The penetration depth is on the

order of 10 to 60 Angstroms. I think it is

probably about 20 Angstroms at 1.5 keV, how- ever the scattering occurs only from the

first surface layer.

Johanessen : Is that region distorted by the

measurement ?

Harrington : It is distorted somewhat , but

with helium scattering the damage done by

helium sputtering is very very slight and at this point we are basically observing the

scattering before the damage has been done. In addition, I think the surfaces do recon- struct and certainly this is borne out by

many of the LEED studies. I cannot give you

a good explanation as to why we see so little

damage in comparison to say reconstruction or

this type of thing. The scattering seems to occur before the damage exists. Even if you

are able to penetrate and sputter several

atoms the signal remains constant.

Johanessen : Would you repeat the value you

cited for the extent of the interface?

Harrington : The silicon signal seems to rise over roughly 10 to 20 Angstroms, which

looks like about 3 molecular diameters for

silicon dioxide. This is about as much as I can tell on such a thick film. I am also

going to pursue this more.

DiStef ano : Could you estimate the contribution of multiple scattering events to the

penetration depth? It is possible that this

could lead to a misinterpretation of the extent of the silicon dioxide/silicon interface .

Harrington : Because you are following both

the oxygen signal and the silicon signal in

a situation like this and you see one rise

before the other falls I do not think that

is really a possibility. The probability

for the scattered particles to penetrate

even several stomic thicknesses and emerge

as ions from a single scattering event, and we can tell the difference between a single

scattering event and a multiple scattering

event, is very very small. I do not know

whether to say it is a percent or less, I really do not know. The fact is that you can see density differences on (111) and (110)

silicon surfaces. If the beam penetrated two layers for instance the density would be the same on the two surfaces; one is basically

behind the other.

Participant : How were your samples prepared?

Harrington : The silicon surfaces as prepared were as finely polished and as free of damage

as we could get them. They were finely polished, a thick oxide was grown on them at, I believe, 600° or 900°C and this thick oxide

was removed by etching in order to get rid of

all the damage basically.

Participant : What is the background vacuum

level you are operating in?

Harrington : Backfilling is done to a static

pressure of about 5 x 10~5 of argon or neon

or helium, whichever you are using. The vacuum before backfilling is 10 and

titanium sublimation pumping is in effect

during the whole process.

30
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg , Maryland, April 23-24, 1975 (Issued March 1976)

SURFACE ANALYSIS BY SECONDARY ION MASS SPECTROSCOPY TECHNIQUES

Robert D. Dobrott

Materials Characterization Laboratory

Texas Instruments Incorporated

Dallas, Texas 75222

The secondary ion mass spectrographic technique implies a source of primary ions, a source of secondary ions and a spectroscopic

method of analyzing the latter. Figure 1 is an artist's conception illustrating the pri- mary and secondary ions referred to in this

presentation. The primary ions in the collimated column are normally generated from a gaseous ambient either by dc or rf fields or

in separate sophisticated ion guns such as duoplasmatrons . These primary ions then

collide with the solid surface which results

in the sputter process which in turn gives

rise to a plasma containing positive and

negative ions, neutral atoms and molecules

having the composition of the solid surface.

These sputtered ions are the secondary ions

which are analyzed by mass spectroscopy.

Figure 1 also illustrates the 20 angstrom

possible escape depth of a secondary ion

which defines the minimum surface depth.

The remainder of this presentation will be

restricted to the three types of commercially

available instruments. Figure 2 is a schematic of the ARL QMAS system. Primary ions are generated in the duoplasmatron , accelerated to 10 kV in the gap between the anode

and pickup electrode, and focused onto the

sample by the primary lens. The sputtered

secondary ions are accelerated by the sample

voltage back through the extraction lens

into a spherical electric sector where they are energy separated. The ions are then mass analyzed by a quadrupole mass spectrometer and detected by ion to electron con- version-scintillator-photomultiplier technique. The quadrupole mass analysis scheme

is illustrated in figure 3. A dc field ap- plied to opposite poles is superimposed with

a radio frequency field. This oscillating

field will then start the ions oscillating

between the poles. At any specified fre- quency only ions of a given mass will undergo

stable oscillation which is necessary for them to pass through the length of the field

without being collected by the electrodes.

The mass is selected by choice of both the

dc field and the radio frequency field. Mass spectra are recorded by continuously

varying both fields keeping their ratios

constant. Figure 4 shows a portion of a 1 ppm boron doped silicon spectrum taken with

this instrument. 2 The boron 10 peak shows

that the sensitivity is well into the sub-ppm

region. This spectrum also shows the mass resolution is sufficient for most inorganic

analysis

.

Figure 5 is a schematic of the Cameca ion

microanalyzer . This instrument is unique in

that it uses a magnetic prism for mass analysis which acts as an ion emission microscope

as well as a microanalyzer. The primary ions

are generated in the duoplasmatron and focused

onto the sample. The secondary ions are ac- celerated and picked up by an electrostatic

immersion lens. These ions are then mass or momentum analyzed in the first leg of the magnetic prism which has both radial and tranverse focusing properties. The ions are directed towards the electrostatic mirror

where only ions with an energy below a pre- selected threshold are reflected back through

the aperture into the second leg of the mag- netic prism. This leg of the magnetic prism

reproduces the single mass ion image of the

specimen surface on the ion to electron image converter where the electrons can be used for

fluorescing a screen, exposing a film or activating a scintillator for photomultiplier

detection. A given mass can be preselected

by using the appropriate magnetic field, or a complete mass spectra can be recorded by

sweeping the magnetic field. Figure 6 is a photograph of this instrument. The ion-gun

and sample are to the left of the magnetic

prism, and the ion to electron converter,

fluorescent screen and photomultiplier to the right

.

Figure 7 shows the scheme of the third type

of instrument. This is the block diagram of

the ARL ion microprobe mass analyzer instrument. The primary ions are generated in the duoplasmatron and accelerated to the primary

magnet where they are mass separated. The

selected ions are then focused onto the sample with a condenser-objective electrostatic

lens combination. The secondary ions are accelerated to the pickup electrode, focused

by a retrofocal lens, energy analyzed with a spherical electric sector, and mass separated

by the secondary magnet system. The mass of interest is allowed to pass through the slit

where the ions are detected with the ionelectron conversion-scintillator-photomultiplier combination. Mass spectra are obtained

31
by sweeping the magnetic field. Ion images

of a preselected mass are obtained by ras- tering the primary beam on the sample and

using the photomultiplier current output to modulate the brightness of a CRT by synchronizing the raster with the primary beam.

Figure 8 is a photograph of the ARL-IMMA

instrument. This particular one is the instrument in our laboratory at Texas Instruments. The duoplasmation and primary magnet

are to the right in the main console and the

secondary ion analyzer system is to the left.

The data collection system is interfaced to

a Texas Instruments 960A computer (installed

in the top rack on the left console) . The

960A can drive a digital plotter which has

been used to generate most of the data fig- ures in this presentation.

This introduction to the instruments, which

is by no means complete, is only intended to cover the general principles; let us go on

to their application to surface analysis.

No review of secondary ion mass spectroscopy

would be complete without showing Anderson's

famous figure 9 depicting the ion yield

comparison between the kinetic and chemical

ionization processes. +

The curve on the left

is the ion yield of Al as a function of

sputter time (or depth) . Note that the ion

yield falls off dramatically as the surface

oxide is sputtered away. The curve on the

right is a repeat of the same experiment

using 02+ as the sputter beam. Here the

surface yield is high; then the yield takes

a small dip, and then returns to the high

surface point after about 40 seconds of

sputtering. The conclusion is that in both

cases the surface yield is primarily the re- sult of chemical ionization (oxides) and in

the Ar"*" case the mechanism quickly converts

to kinetic ionization. For the O2 beam

the dip is related to a combination of

kinetic and chemical ionization which becomes completely chemical when the sputtered

depth reaches the original primary ion implant depth. This experiment has lead to the widespread use of oxygen as the primary

ion sputter beam while maintaining a relative high partial pressure of oxygen near

the sample surface.

Figure 10, also from Anderson's 1 work, shows

the relative intensities of secondary ions

obtained from pure metals using an 0~ sputter

beam. From this table Al appears to be the

most intense, and noble metals represented

by Au the least intense. This chart, even

though taken from the pure metals, does cor- relate moderately well with the sensitivities

determined for these elements in the silicon

or silicon oxide matrix. However, for bulk

type analysis in silicon, the ordinate is

from 1 ppb to 10 ppm going from top to bottom.

This illustrates that secondary ion mass spectroscopy is a very sensitive technique

even when material quantities are small, as is the case in surface analysis.

Figure 11 is a series of spectra taken from a silicon surface with identical primary 02+

beam density, with the analyzed surface depth

controlled by either rastering the primary

beam or moving the sample. The calculated

sputter rate for these spectra was about

50 A/sec. As can be seen in these spectra,

by far the most sensitive method in surface

analysis is to move the sample and thereby

constantly expose a fresh area to the primary

beam. In this spectrum the carbon, hydrocarbon fragments, sodium, potassium and calcium

peaks are quite pronounced. The 400 micron

raster spectrum does show the same peaks but

not nearly so intense. Also, note that the

overall intensity is reduced, which tells us that the whole spectrum was recorded in the

region between the surface and primary ion

implant depth. The 50 micron and stationary

spectra both suggest the surface is quite

clean with only traces of Na and K appearing.

I have not labeled the peaks on these spectra

since the point I wish to make is the number

of peaks and their relative intensity. However, if you care to do some quick mental

exercises, I will remind you Si ha.s 3 isotopes

at mass 28, 29, and 30, and Si related peaks

will always appear as groups of three to five

peaks. Also, you may note that the Si02

group at mass 60 (spectra shifted to the

right) are prominent in the 50 micron and

static beam spectra. This, along with the

higher intensity, tells us that the spectra were recorded after reaching the implant

depth. This molecular spectra is characteristic of structural chemical bonding and has

been investigated by several of my European

colleagues for surface fingerprint spectra.

Other techniques such as decreased beam intensity or faster recording could have been

used to insure the surface wasn't sputtered

away before the entire spectra was recorded,

but it must be remembered that sensitivity is

directly related to ions/sec delivered to the

detector, and that a little integration goes

a long way at the low end.

Figure 12 is the analysis of various fractions

of a monolayer of Ag deposited on a Si surface. The monolayers were deposited by

electroless deposition from an aqueous solution which contained *

1 *Ag tracer ions.

The coverage was then determined by gamma

counting the tracer on the Si slice after

deposition. The analysis was done with a 20 kV, 15 nanoamp, 10 micron, 02+ beam rastered

to a 200 x 160 micron area. By electronic

counter gating synchronized with the primary
beam raster, the count acceptance area was restricted to the central 1/4 of the crater

bottom. The depth assignments were made by

measuring the total crater depth with optical interf erometry and assuming a uniform

sputter rate. This family of curves all

show the same trend, a very intense surface

peak rapidly falling off before the 50 angstroms depth, then increasing slightly as the primary ion implant depth is approached,

and then taking a second sharp decrease to a constant level well above background. These

shapes show that much of the surface Ag is

knocked into the silicon by the primary ion

beam, hence, the second rapid decline. This

is bad if we were interested in a concentration depth profile, but actually beneficial

in a surface analysis since it does contribute to detectability time for quantitative

analysis. However, if we are struggling for

quantitative analysis, which is every mass spectroscopist 's impossible goal, we would

like to minimize this knock-on effect even

though it is only 1 part per hundred. Figure 13 is an attempt to reduce the knock on, or at least the depth, by lowering the primary ion energy. The general trend does

show that the second sharp drop is indeed

decreased as the energy is reduced. Only

the 12 kV curve violates this trend and is

probably in error. The actual depths for

these curves were extremely difficult to measure and could easily be as much as 100%

in error. One final point to note on these

curves is that only the 20 kV curve shows a positive slope after the first surface drop.

This results from the decreased sputter

rate, with the lower energy beams allowing

the ambient oxygen to keep a constant surface oxide throughout the analysis. Figure

14 is a closer look at the surface region as

a function of primary ion energy. The set illustrates that even in the worst case at

20 kV, we will have collected about 90% of

the integrated intensity before 50 angstroms

have been sputtered away. Consequently,

this illustrates that for quantitative surface analysis by secondary ion mass spectrometry, it will probably be for only one mass per area sputtered. If the area is

limited such as on a small signal diode,

this can be a severe limitation. However,

on pilot slices, which encompasses most of

my work, the problem is not limiting as it can be overcome by continual movement of the

sample. Figure 15 is one last look at the curves as a function of primary ion energy.

These are the same curves as Figure 13, but now the abscissa is plotted as a function of

time. This set shows that for qualitative

analysis, as much as 10 minutes was available to detect the Ag which was present at about 5% of a monolayer.

Figure 16 shows how quantitative the results

shown back in Figure 12 were. The line is a least squares fit to the five points indicated. Each point represents the total integrated count under the respective curve, with

the base line taken where the curve leveled

off after the second sharp drop. This empirical calibration curve is very encouraging

even though two of the low points are off by

as much as 100%. No correction was made for

ion pick-up efficiency by means of a matrixion and this is known to vary by as much as 50% from sample to sample. This study does

show that given a set of standards, and I believe that bulk standards can be used,

secondary ion mass spectroscopy surface analysis can be empirically quantitated, provided

the surface impurity concentrations are not high enough to drastically alter the effective

matrix which controls the ion yield efficiency.

Figure 17 turns the topic over to the analysis

of Si02 surfaces, whereas up to now I have

concentrated on silicon surfaces. This is a series of spectra taken under various conditions on a 10,000 angstrom thick oxide film.

The oxide was not coated with a conductive

film, as would be the normal practice for

bulk or in depth analysis, since this would

mask the surface of interest. As can be seen

from this set of spectra the use of a positive

beam on a naked sample gave only a few peaks, none of which are normal patterns expected

from a Si02 matrix. The absence of a spectra

is the result of a high positive charge

buildup at the point of impact by both the

positive ion beam and loss of secondary electrons. Also, without a conductive surface

coating there is insufficient bias on the

sample surface to accelerate and direct the

secondary ions to the pickup electrode. The

negative beam on the naked sample does show

the silicon peaks but they are very weak.

This does show that the secondary electron

loss is compensated by the negative primary

ion which allows the point of impact to come

to an equilibrium state. The absence of the

remainder of the spectrum here is entirely

due to the absence of the accelerating sample

bias. The two spectra on the right were

taken through the open areas on a 300 mesh

electron microscope copper grid which was electrically connected to the 1500 volt sam- ple bias. The primary beam was rastered

rapidly over both the copper bars and the

openings. The theory is that the primary

beam bombarding the copper emitted sufficient

low energy secondary electrons, which neu- tralized the positive charge on the Si02

surface and provided the necessary secondary

ion accelerating potential as well. These

spectra do show that the mechanism works;

both positive and negative beams give

33
satisfactory spectra. However, these electroplate formed copper grids are not noted for

their purity, so a much more pure grid must

be employed to render meaningful results.

Grids formed by e-beam evaporation of high

purity Al or Au should be excellent. This

discussion only applies to oxide films

greater than 1500 angstroms thick. Thinner

oxides behave the same as bare silicon.

Therefore, when searching for the source of

contaminants, such as furnace tubes, the

oxide thickness should be kept below about

1200°C which just happens to be the gate

oxide thickness of most MOS devices.

No presentation concerning oxide analysis

would be complete without illustrating the

mobile ion problem shown by Ron Baxter 4 in

Figure 18. These profiles were taken on the same sample using both a positive and negative beam. The conclusion is that using a positive beam, the Na all moves to the interface and with a negative beam, it is all

drawn to the surface. The grid technique

may help this problem as well, but to date

the grid itself has been the dominant source

of sodium.

To summarize this presentation, a look at secondary ion mass spectroscopy with respect

to five criteria for surface analysis is in

order.

1. Accuracy of Chemical Analysis

The qualitative analysis was shown to be

very good and the Ag experiment showed

quantitative potential is fair.

2. Surface Sensitivity

The sensitivity has been shown to be as good as a few ppm.

3. Non-Destructive Analysis

The method is totally surface destructive.

4. Selected Area Chemical Analysis

Can be very good depending upon instrument .

5. Experimental Complexity

Secondary ion mass spectroscopy is one

of the most complex analytical techniques in use.

Therefore, for surface analysis, secondary

ion mass spectroscopy has not yet reached

the ideal, but it has demonstrated itself to be one of the more powerful techniques.

REFERENCES

1. Anderson, C. A., International Journal

of Mass Spectrometry and Ion Physics

2, 61 (1969).

2. Fralick, R. D., and Roden, H. J., ASMS

Committee VII Workshop, 22nd Annual ASMS

Conference, Philadelphia, (May 1974).

3. Larrabee, G. B. , unpublished work, Texas

Instruments Incorporated, Dallas, Texas.

4. Hughes, H. L. , Baxter, R. D. , and

Phillips, B., IEEE Trans. Nucl. Sci.

NS-19 , 256 (1972).

34
ION MICROPROBE

ION BEAM

9

0 •

oO fP

oo

°G

o O o o oo

°>°o°

°o ° o O

o °OoQ

°ooon ooo0

°

TO MASS

SPECTROMETER

o i

O

^3 »

O

I20A

Figure 1. Artist's conception of the

secondary ion production process.

35

37
ALUMINUM

50 100 TIME (SEC )

38 n +

J L

50 100

Figure 9. Sputter yield comparison for

kinetic and chemical ionization mechanism.

Accel, potential 11 KV

Sample current 4 * 10~ 9 A Probe size 20 ym2

Figure 10 (below). Relative yield of some positive ion intensities from pure metals.

0.0) 40 50 Z ATOMIC NUMBER

90
10 »

SECONDARY ION MASS SPECTRUM

10'

10 *

5 .

10 2

10 0

~~

i

1 r SILICON

400 MICRON

RASTER

0 10 20 30 40

nflSS NUMBER SECONDARY ION MASS SPECTRUM

10 s

10*.

10

i i r SILICON

8 : i

.

10»_

10

10 1_

0 10 20

50 MICRON

RASTER

SECONDARY II IN MASS SPECTRUM

10'.

io 2 _

10 0

~~

i r~

SILICON

Jul

—i r STATIC BEAM

10 20 30 40 SO 60

HR33 NUMBER SECONDARY ION MASS SPECTRUM

Figure 11. Surface spectra of Si as a function of depth consumed.
10 5

'10 *

10'

102

i

1

1 r

3/27/75

J ] U J_

200 400 600 800 1000 1200 1400

, ANGSTROMS RG/S]

Figure 12. Ag on Si surface and depth pro- file ion intensity as a function of monolayer coverage.

450 600 ANGSTROMS

900 1050

AG/SI 4.7E13 ATOMS PER SQ. CM.

Figure 13. Ag on Si surface and depth profile ion intensity as a function of primary

ion energy.

1 1 1 1 1 3/27/75

i 6 KV \ o 8 KV

» 10 KV \ « 12 KV

:^jb--0\ • 20 KV

3

?

1 1 30 40 ANGSTROMS

0 10 20

AG/SI 4.7E13 ATOMS PER SQ. CM. Figure 14. Surface expand view of figure 13.

0 300 600 900 1200 1500 1800 SEGON0S AG/SI 4.7E13 ATOMS PER SQ. CM.2100

Figure 15. Ag on Si surface and sputter

time dependent ion intensity as a function

of primary energy.
0.5 1.0 15

Ag (107) COUNTS X I0" 6

Figure 16. Empirical quantitative curve for monolayer Ag on Si surfaces.

Figure 18. Na profiles in SiC>2 taken with

positive and negative primary beams.

41
10 *_

10 »_

i r Si0o NEGATIVE OXYGEN

MOVING SAMPLE

_j iJbu

10 20 30 40 MASS NUMBERSO 60 SECONDARY ION MASS SPECTRUM

10

10 4

10'

10

SiO,

i

1 1 r~

NEGATIVE OXYGEN

THROUGH GRID

30 40

rtflSS NUMBER SECONDARY ION MASS SPECTRUM

Figure 17. Surface mass spectra of uncoated Si0 2 surfaces and mass spectra taken through

a 300 mesh grid.

42
DISCUSSION OF THE PAPER be accompanied by a corresponding dip in the

silver signal?

Participant : I did not understand what you

said concerning the appearance of the Si02

peaks in the silicon.

Dobrott : The Si02 peak appeared in the

silicon analysis after the initial oxide surface was removed simply because I was ana- lyzing the silicon dioxide formed at the implant depth of the primary oxygen beam.

P articipant : Could you have avoided that

problem by using 180 instead of 1^0 as the

primary ion?

Dobrott : There was no mass interference

problem. I made that observation to point

out the sample depth at which that particular

spectrum was gathered.

Morabito : Should not that dip in the Si02

profile, due to the primary oxygen implant,

Dobrott : The silver profiles determined at lower primary ion energies did not really dip,

they went down and never regained; probably

because of the proximity of the initial implant depth and the surface. The higher

energy profiles dipped and recovered probably

due to the implantation process.

Participant : In the silver case did you

assume the silver was uniform, or could it

have formed islands, or what?

Dobrott : We assumed that there was a uniform

deposition of silver. The analysis was done

to see if we could correlate the known amount

of low temperature electroless deposited sil- ver on the surface with what we get when we sputter it all away, and to see how quickly

this could be done. Yes, the silver could

have formed islands, we do not know that it

did not.

NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 197.5 (Issued March 1976)

SOME EFFECTS LIMITING SIMS DEPTH PROFILE ANALYSIS

AND METHODS FOR IMPROVEMENT

Robert K. Lewis

Cameca Instruments, Inc. Elmsford, New York 10523

INTRODUCTION

Secondary ion mass spectrometry (SIMS) has

proven to be a very attractive method for obtaining surface and bulk concentrations of materials used in semiconductor device technology. The popularity of SIMS has resulted

primarily from the fact that the method has

the sensitivity necessary to detect doping

level concentrations, namely 10+-'-^ to I0+17 atoms/cm-^, and the ability to produce depth

profiles at these concentration levels. 1>2 The disadvantage of being a destructive

method is readily offset by the inherent

depth analysis capability which results from

the sputtering action of the primary beam.

It is this depth profiling capability, in fact, that has provided the principal interest in SIMS for analyzing semiconductor material and the technique is now widely used

for characterizing these materials. However,

there are problems with SIMS depth profiling

which must be taken into consideration.

This paper discusses some of the effects

which have been encountered and tend to limit the usefulness of the technique and

describes methods to minimize them. These

effects have been summarized in Table I.

DEPTH RESOLUTION

To achieve good depth resolution it is obvious that the surface of the crater bottom

must be maintained as flat as possible.

With normal or non-normal incidence of the

primary beam a high degree of flatness can

be easily maintained by rastering the beam.

Schematic diagrams of craters produced by

non-rastered and rastered beams are shown in Figure 1 and Figure 2. Figure 1 shows that

with a non-rastered beam, 250ym in diameter,

incident at 45° (this is approximately the

bombarding angle in the Cameca IMS 300 in- strument used for this study), and a gaussian

distribution, the flatness variation over a typical selected area (50ym) is approximately 200 A. This is adequate for many depth

profiles. However, if the beam is rastered

as shown in Figure 2, the flatness over the

area selected for analysis can be greatly

improved. We have established that the

variation in flatness can be less than 40 A

over an area of lOOym to depths of about

1,000 A by bombarding amorphous Ta205 and

using the technique of observing the variation

in color of fringes produced in white light.

1

This would give a value of less than 5% for

the variation in depth divided by the total

depth sputtered. This value is consistent

with the R value found by Werner 3.

It is clear that ions sputtered from the crater wall must be rejected to have good

depth resolution. This is easily accomplished

in the Cameca ion microscope by mechanically

aperturing in an image plane or in an ion

probe by electronic gating (turning off the secondary detector when the primary beam is on the crater wall) . Mechanical aperturing

is shown schematically in Figures 1, 2, and

16. With electronic gating the primary beam

must be kept smaller than that required with

mechanical aperturing. This tends to limit

the maximum average beam density obtainable

and consequently the maximum sputtering rate available with ion probes.

If the sample surface does not sputter away evenly because of a lateral variation in

sputtering rate, the depth resolution will

suffer accordingly. Lateral variation in

sputtering rate results from different lattice orientations encountered by the primary

beam from grain to grain in polycrystalline

material or from different lattice structures when different phases are encountered. For

this reason the depth resolution attainable

in such material is usually only ten to twenty percent of the total depth sputtered.

Continuously rotating the sample while sputtering could reduce the lattice orientation

effect; however, no commercial instruments

presently provide this capability. When

lateral differences in orientation or in

phase are encountered the use of an oxygen

leak can sometimes minimize the variation in

sputtering rate.

Channeling (lattice effects) in SIMS have

been described by Slodzian and Bernheim^.

This is a variation in ion yield that occurs when the incident beam traverses transparent

and opaque directions in the lattice. A plot

of secondary ion intensity versus the angle

of rotation is shown in Figure 3. Fortunately, this effect is only a second order effect

so it is generally not a problem in depth

profiling. However, it is important to be

45
aware of the fact that it can cause errors especially when comparing samples. A method

of minimizing the effect is to bombard the

surface with a stream of oxygen (oxygen leak)

This produces a less severe variation in ion

intensity with orientation as shown by the

dotted line in Figure 3. Slodzian has

pointed out that the effectiveness of the

oxygen leak is determined by whether the surface oxide formed is crystalline or amorphous.

"Knock on" is the redistribution that occurs

from the physical driving in and sub atomic

mixing of the atoms being analyzed. The effect has been shown by McHugh^ analyzing

Ta2C>5 which had an approximately 50 A thick

phosphorous layer located 230 A below the

surface. His results are shown in Figure 4. The effect is significant only at high primary bombardment voltages (above about five

kilovolts) . The effect would be less, of course, at a higher angle of incidence. The

angle of incidence can be changed for non- normal incident primary beams by changing the

polarity of the primary beam. The combinations positive primary-positive secondary and

negative primary-negative secondary give the

highest angles of incidence.

DETECTION SENSITIVITY

One of the most serious problems encountered

in SIMS is the mass interference from the

complex polyatomic species produced by the

secondary ion process. These species are formed from the clusters of the matrix atoms formed by the sputtering action of the beam

(e.g., M, M2, M3,...) and combinations with

the primary ion when a reactive gas such as oxygen is used (e.g., MO, M02....M20, M2O2,

. . .etc.)

Typical spectra produced when bombarding

silicon with a reactive gas (oxygen) is

shown in Figure 5 due to Evans^. It can be seen that the problem of mass interference

is more severe at the higher masses. This is

just the opposite of the ion production

process in spark source mass spectrometry

(SSMS) where the interferences are more con- centrated in the lower mass region of the mass spectrum. This is because in the spark

source the energy available is much higher

which breaks up the clusters and creates many

multiply charged ion species.

The simplest method for minimizing mass interference from the polyatomic species is to take advantage of the large differences in

energy distribution between the polyatomic

and monatomic species. These distributions

are shown in Figure 6 which is from

Satiewicz^. It can be seen that the complex

polyatomic species can be rejected perferentially over the monatomic species by re- jecting the low energy ions.

The low energy ions may be rejected in the

Cameca instrument by using the low energy

discriminator (L.E.D.) shown in Figure 7 or

by appropriate adjustment of the electrostatic analyzer when the double focusing ar- rangement shown in Figure 8 is used. Re- jection of the low energy ions with the ESA

is possible because the double focusing

geometry used has a first order direction

focus position such that slits can be used

to precisely define the energy limits accepted. The SIMS instruments designed by

Herzog et. al. (GCA) 8 and Tamura et. al. (Hitachi) ^ have this feature while the ion

probe SIMS instruments described by Liebl

(ARL) 10 and Banner et. al. (AEL) 11 do not.

The attenuation of the silicon polyatomics in

the spectra shown in Figure 5 with increasing

attenuation of the parent monatomic peak

using the L.E.D. on the Cameca instrument is

shown in Figure 9. A typical application of

the L.E.D. is shown in Figure 10. Here the

detection of arsenic using the AsO

- species

has been improved by one order of magnitude.

An equally important use of the L.E.D. is

the removal of tails on the low energy side

of a mass peak. The presence of these tails can cause a serious loss in abundance sensitivity (and therefore detection sensitivity)

on the low mass side of intense peaks. This

is shown for the detection of copper and zinc

in gallium arsenide in Figure 11. These

tails result from the post ionization of the

neutral species leaving the sample. This

"kinetic" emission process has been described

by Blaise and Slodzian-^. it is not necessary

to attenuate the intensity of the peak to eliminate the low energy tails with the L.E.D.

A second and more elegant method for eliminating interference from polyatomic ion species

is the use of high mass resolution. The

method of obtaining high mass resolution on

the Cameca instrument is shown in Figure 8. Here the normal imaging mode path, that is

prism - mirror - prism, is interrupted and

the beam is allowed to pass through a hole in

the mirror (the mirror being displaced) into

a sperical electrostatic analyzer (ESA)

This combination of an ESA with the first

magnetic deflection is an inverted NierJohnson double focusing configuration. With

this .analyzer mass resolutions over 5000 have

been achieved. It is possible with high mass resolution to completely separate the interfering polyatomic mass peak ^' AI2 from the

->^Fe peak as shown in Figure 12 because of

the mass differences that occur. Some

46
typical elements, the possibly polyatomic

interferences, and the mass resolution re- quired to resolve the doublet are listed in

Table II. The separation of the

75As/ 29 Si 30 SiO doublet at the peak and lower

limit of the distribution, for arsenic ion

implanted into silicon, is shown in Figure 13.

A second effect limiting detection sensitivity

is the contamination of the surface from the

ambient gas molecules in the vacuum system.

At pressures around 10" 7 to 10"^ torr, which

is the typical vacuum level in commercial

SIMS instruments, it can be readily demonstrated that the low current densities present in the beam periphery produce an artificially high background signal. This is

because the influx of ambient gas molecules

now competes with the primary beam and react with the surface where they are sputtered

away producing various interfering ion

species. Low beam densities will always be

present at the periphery of an ion probe.

In addition there is a significant flux of

neutral atoms (neutralized ions) over a large

area outside the bombarded region that also

sputter the target surface and produce ions

from the ambient gas as shown by McHugh-'- 3 in

Figure 14. A cold plate near the sample

surface will minimize the hydrocarbon con- tamination, but it will not eliminate it and

it will have little or no effect on oxygen,

nitrogen and CC^. The effect is shown

dramatically using the Cameca instrument by

bombarding aluminum with a beam of nonreactive ions (Ar"*") . By keeping the beam

diameter less than the field of view (225um)

it is possible to see the effects of low

primary beam density at the periphery of the

beam on the fluorescent screen as shown in

Figure 15. At the periphery there is a marked increase in the secondary ion intensity due to an ion yield enhancement (chemical effect) from the reactive ambient gas. By mechanically aperturing as shown in

Figure 16 and selecting areas approximately

50ym in diameter at the center (Area A) and

the periphery (Area B) the spectra shown in

Figure 17 were produced. It is obviously

impossible to achieve a low background signal without rejecting this ambient contribution. The only methods that have been demonstrated to be effective in eliminating the

ambient contamination are lowering the vacuum below 10~ 9 Torr as Benninghoven-L^ has

shown which unfortunately is impractical in

commercial SIMS instruments or the ions produced outside the area of high beam density

can be eliminated by the mechanical aperturing technique. However, there is no way to avoid this effect with the electronic gating

technique used with ion probes.

A third effect which can limit detection

sensitivity is the phenomena of resputtering.

This effect is observed when material in high

concentration near the surface at the edge of

the crater is sputtered and the neutral particles are scattered in some manner into the

analyzed area at the crater bottom. This is

essentially a second order effect. It can be

kept below the detection limits by using

crater diameters much larger than the area

analyzed. Evaporating non mass interfering

metal (e.g. gold) on the surface prior to depth profiling can also be helpful.

ACCURACY

The most serious effects limiting the accuracy

of a depth profile are those limiting the

uniformity of the ion yield of the element

being measured or a reference element (e.g. a matrix element) used for normalization.

Any marked change in concentration of a reactive element such as boron, carbon or oxygen can change the ion yield of elements

associated with it. This effect is shown in

Figure 18 where phosphorous enhances the

yield of the tantalum ion. It is important

to determine if such species are present by

in-depth mass spectral analysis whenever

possible, and, if present, to measure the

depth profile of the reactive species simultaneously with the element of interest. This

also applies to any element being measured

for normalization purposes. Sometimes it is

useful to check for the presence of a reac- tive element by using an inert gas as the

primary beam (e.g. argon) which does not en- hance the matrix greatly increasing the

contrast

.

Whenever a reactive primary beam is used

there can be a marked change in the ion yield

over the first few hundred angstroms of sputtered material (15kV bombardment) due to the

implantation effect of the primary beam as described by Lewis et. al.-'--'. The effect is

shown in Figure 19 where pure silicon was bombarded with 0~ primary ions and positive

secondary ions of oxygen and silicon measured

together. This ion yield variation results

from the fact that the implanted primary

oxygen ion is first concentrated at a depth

below the surface as shown schematically in

Figure 20. Until the sputtering front

reaches this depth there will be a variation

in ion yield of all elements which are en- hanced by the primary ion. The effect is

minimized by keeping the bombardment voltage

low (below 5kV) and can be essentially eliminated by the use of an oxygen leak as shown

in Figure 21.

47
Variations in sputtering rate with depth can cause errors in depth measurements if the

variations are not taken into consideration

in the sputtering rate calibrations used to establish the depth scale. Often measurement

of a matrix species and normalization against

this signal will take care of small sputtering rate variations due to beam density

changes. However, variations due to phase

changes in the sample must be determined by

depth measurements outside the instrument.

If sufficient thicknesses of the material in

the individual phases are available (a few

hundred angstroms) it is possible to make

their depth calibrations with a Talystep. If

larger thicknesses (a few thousand angstroms)

are available it is more convenient to make

the measurements with a light microscope

equipped with a Michelson interferometer.

It is not unusual to encounter a situation

where what appears to be a variation in depth

(shown for aluminum in the thin film depth

profiles shown in Figure 22) is actually a variation in lateral distribution as shown in

the ion image of Figure 23. This variation

may be real or a result of variations in the

lateral sputtering rate. The lateral distribution effect is observed most often at metal - metal oxide interfaces or thin film

metallizations. It is difficult to minimize

this effect. However, its presence can usually be taken into account by observing a series of ion images with depth.

One of the major concerns in all depth

profiling work is whether the element of interest is being redistributed by the bombardment process. This is the case of the "knock

on" phenomena described above, which was described as a second order process. The

only first order redistribution effect observed to date has been for the analysis of

sodium in silicon dioxide layers. This re- distribution occurs whenever a potential is

applied to the surface. The sodium migrates

(within milliseconds) to the Si02/Si interface or to the sample surface depending on whether the potential applied at the sample

surface is positive or negative as shown in

Figure 24. It is questionable whether an accurate profile of sodium in Si02 can be obtained at all since even neutral atom bombardment will release secondary electrons and

charge the surface. The surface to substrate

potential must not exceed a few tenths of a volt while the profile is being measured.

Fortunately this ion mobilization phenomena

appears to be a problem only for the case of

sodium in silicon dioxide. No other documented instances are known to this author.

References

1. Morabito, J. M.,and Lewis, R. K. , Anal.

Chem. 45_, 869 (1973).

2. Kofkner, W. K. , Werner, H. W. , Oosthoek,

D. E.,and de Grefte, H. A. M. , Rad.

Effects 17, 88 (1973).

3. Werner, H. W. , Acta Electronica 18, 51

(1975).

4. Slodzian, G. , and Bernheim, G. , Int. J. Mass Spectrom. Ion Physios (1975).

5. McHugh, J. A., Rad. Effects (1973).

6. Blattner, R. J., Baker, J. E., and

Evans, Jr., C. A., Anal. Chem. 46, 2171

(1974).

7. Satkiewicz, R. G. , Air Force Technical

Report, AFAL TR-69-322 (1970).

8. Herzog, R. F. K. , Poshchenrieder , W. P.,

Ruedenauer, F. G. and Satkiewicz, F. G.

Proc. Fifteenth Annual Conf. Mass Spectrometry and Allied Topics, Denver,

Colorado, May 1967, p. 301.

9. Tamura, H. , Kondo, T. and Doi, H.

Advances in Mass Spectrometry, Quayle,

Ed., (Inst, of Petroleum, London,

1971), Vol. V., p. 441.

10. Liebl, H., J. Appl. Phys. 38, 5277

(1967).

11. Banner, A. E. , and Stimpson, B. P.,

Vacuum 24, 511 (1975)

12. Blaise, G., and Slodzian, G. , J. Phys.

31, 93 (1970).

13. McHugh, J. A., to be published.

14. Benninghoven , A., Surf. Sci. _2J5, 541

(1971).

15. Lewis, R. K. , Morabito, J. M. , and

Tsai, J. C. C. , Appl. Phys. Lett. 23,

260 (1973).

48
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C

4J co qj s u 3 TO

QJ

k rH o X o

Piw PS o o rJ

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t-l 00 4->

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1-1 e p. cu CO

a o • o CU

u

u o 4-1 oo rl 3 4J >H <4—1

MH CO 3

4J 01

O c CO CU

cd 00 1 t-l •H cu

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a c C o CO 3 a

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49
TABLE II

MASS RESOLUTION REQUIRED TO SEPARATE DOUBLETS

OF TYPICAL INTERFERING POLYATOMIC ION SPECIES IN SIMS

Nominal

Mass Species M/ AM

14

UN/ 2?Si++ 1,000

27 nB 160/

27Al 1,200

48

24Mg/8Ti 2,200

54

27.. ,54^ Al / Fe

z 2,300

56

24

Mg0

2

/ 56 Fe 1,400

56 28Z ,56 °Si

2 r

D Fe 3,000

59

27 A1 16n ,59„ Al 0^1 Co 1,500

60

28Q .16 n

,60... Si 0^1 Ni 1,700

64

48^.16^ ,64„ Ti 0/ Zn 4,700

75

29 c .30 c

. n/ 75. Si SiO/ As 3,200
To Mass Spectrometer

Aperture -

Secondary Beam

—1 1 ( |(— 50/tm I \^.

/

/ /

/ /

45°

'

<*— Primary Beam

Sample 0.017 fj.m —

350/im

To Moss Spectrometer

Aperture

Secondary Beam

l/xm

Rastennq

Plates

—50/i m

350/xm

50/i.m

45"

Primary Beam

Sample

Figure 1. Crater produced by unrastered

primary ion beam.

Figure 2. Crater produced by rastered

primary ion beam.

0' 90' 180' 270' 360'

Figure 3. Variation in ion yield with

different orientations of primary beam.

Figure 4. "Knock on" redistribution from

ion bombardment

.

~\ r

1

1.75 keV ,60

,78 A

n 1 r- 7.75 keV l60

~i 1 r -i 1 r

18.5 keV ,60

3I P-RICH- LAYER

230 A

I60

ION BEAM

//)y//A

}

To2°5

200 200 400 200 DEPTH (A) 400 600

51
00001

1 1 1———

High Purity Silicon Oxygen Bombardment

I- 1 i *t* I " 1

1 — i 1— " *-

i " — ————

1 0 10 20 30 40 50 60 70 80 90 100 110 120 130 140 ISO

Mass/Charge

Figure 5. Mass spectra of silicon polyatomics

0 200 400 600 800 1000 1200

INITIAL ENERGY eV—

Ion beam

Electro - magnet

GRID FILTER

Figure 7. Low energy discriminator grid

filter on the Cameca IMS.

Figure 6. Energy distribution of polyatomics .

%1
PRIMARY

ION GUN

V

SAMPLE^b^'IAELECTRON MULTIPLIER DETECTOR

ION nin

"OR JJl

MAGNETIC ANALYZER

_tTO IMAGE

J CONVERTER

IT*—SLIT B ADJUSTABLE

'SLITC SPHERICAL CONDENSER

ELECTROSTATIC ANALYZER

Figure 8. Double focusing ion optics acces- sory for the Cameca IMS.

Figure 9. Attenuation of positive secondary silicon

molecular ion species with low energy discrimination.

Si ATTENUATION r

53
111.

mri, L In,

MASS —> 85

Figure 10. Attenuation of the Si20 2 interference with

AsO using low energy discrimination.

F.S. L. E. D.

Counts

1,000

Figure 11. Attenuation of low energy tails

from gallium in gallium arsenide interfering

with the detection of copper and zinc.

^4
Arsenic in Silicon

At/cm 3 As

As

SiO;

-- 10

M

i 9

AM

At/cm 3 As

10 19 "

= 3200

--10 18 10 18 "

SiO;

As t Jl_

Figure 12. The separation of the aluminum

dimer from the iron 54 isotope using high

mass resolution.

Figure 13. Separation of arsenic from SiC>2 at the peak and tail of an arsenic ion implantation profile.

Figure 14. Schematic of neutral bombardment

associated with an ion probe.

Figure 15. Enhancement of the aluminum ion

yield at the periphery of an argon beam bombarding aluminum.

II
Figure 16. Mechanical aperturing in a Figure 18. Phosphorous enhancement of tan- secondary ion image plane to eliminate con- talum signal in Ta2 0 5

. tributions from periphery of the ion probe

and neutral bombardment.

si
Depth

Figure 19. Depth profiles of silicon and

oxygen in pure silicon bombarded with oxygen.

I : 0", V : 15 kV, P: 10_17 Torr

P P

Figure 20. Primary oxygen beam implantation Figure 21. Primary oxygen beam implantation

depth profile. depth profile with oxygen leak.

1?
Figure 22. Depth profiles of platinum - aluminum thin films sputtered onto silicon.

Figure 23. Aluminum ion image of platinum - aluminum thin film sputtered onto silicon at Al/Si interface.

Depth Depth

Figure 24. Schematic diagram of sodium

redistribution in SiC>2 with applied fields.

5!:'
DISCUSSION OF THE PAPER

Hurd : How does the SIMS profiling technique

compare to other profiling techniques, such

as NAA, spreading resistance, or the four

point probe techniques?

Lewis : Actually I should probably refer you

to one of the people who has used SIMS and

those other techniques, Joe Tsai could tell

you more.

Tsai : Our results show very good correlation

as a matter of fact.

Morabito : Another point is that the SIMS

technique can measure the electrically inactive as well as active species.

Lewis : The electrical techniques only measure the concentration of electrically

active impurities whereas SIMS measures the

actual chemical concentration. I should also

say something about the speed of analysis.

A typical SIMS profile can be obtained in a matter of minutes, but your spreading resistance measurement takes how long?

Participant : About a day or so.

Lewis : Right. So I think it is important

to point this out. You can, in a day, do

easily 40 or 50 SIMS depth profiles. While

the high performance SIMS instruments are quite expensive, the speed of analysis also

has to be taken in consideration. This is one thing I like to point out when SIMS

analysis is compared to electron probe

analysis. At the same concentration level

SIMS can typically do analysis five to ten

times faster. This can readily offset the

difference in cost of these machines. The

real interest to semiconductor people, however, is that SIMS has the sensitivity to de- tect the dopant concentration levels and I do not know of any other technique that does

as well.

Harrington : You can also tell what the

dopant species is, whether it is boron, or

aluminum, or phosphorous. With the electrical

techniques you are only guessing.

Lewis : Right, we are identifying the species

and, as a matter of fact, we are correlating

this information with other possible electrically or nonelectrically active species, so the information that you gather this way is

considerably more useful.

DiStefano : What sort of total spatial

resolution do you obtain with your SIMS

technique?

Lewis : The depth resolution can be a few

tens of angstroms if you are dealing with

single crystal silicon. But there are many

limiting effects. Usually we do not like to quote a depth resolution capability better

than 10% of the depth sputtered because of

sputtering artifacts. However, in silicon

you are usually working with a very ideal

situation and much better resolutions can be

obtained.

DiStefano : If you sample a large area you

could obtain a resolution of 10 A, but for a smaller area you obviously would have to sample a greater depth. What combinations

of sampled area and sampled depth are attainable?

Lewis: You must always take into consideration the amount of material sputtered and

the ionization efficiency of the particular

species you are looking at. As the area is

made smaller and smaller you must sputter to greater depths to get enough ions to achieve

an adequate precision. I am talking about a bulk analysis now. In a concentration gradient this does not really apply. The answer

is that the lateral spatial resolution is

about a half a micron in bulk analyses.

However, this is not a practical area to use

for a depth profile. A practical limit to depth resolution for a depth profile is a few tens of angstroms for an area 100 pm or

so in diameter.

5S

NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

Qualitative Assessment of Ion Erosion Damage

by Means of Electron Channeling Patterns

Dale E. Newbury

Analytical Chemistry Division

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

In secondary ion mass spectrometry (SIMS),

the sample is bombarded with energetic ions,

usually in the range 5 to 30 keV. These

primary ions impart their energy to the atoms of the sample via inelastic processes,

resulting in the sputtering of atoms lying

at or near the surface. The ionized fraction of these sputtered atoms, the secondary ions, are then mass analyzed, forming

the basis of this spectrometry.

Concomitant with the sputtering of atoms is the damage to the sample produced by the

primary ion — solid interaction along the

entire range of primary ions. Lattice defects such as vacancies and interstitials

are created by the displacement of atoms

from lattice sites due to energy imparted by

the primary ion collisions. The effect of

ion bombardment is to degrade the crystal

perfection. Such degradation occurs in a shallow layer near the sample surface, the

depth of which is governed by the range of

the primary ions. Typically, this range

will be of the order of 50 nm or less for

the ions, for example, Ar+

, 02+, 0+, N2+

, N*, 0~, and the beam energies used in SIMS

and other surface analysis techniques which

make use of ion bombardment for erosion. 1

Assessment of the damage induced by ion bombardment is possible by use of the electron

channeling pattern technique in the scanning

electron microscope (SEM). 2 ' 3 '

1* Briefly,

electron channeling contrast is obtained in SEM because of the influence of the periodic

structure of a crystal on the initial

electron-specimen interaction. The electron

channeling pattern (ECP) is generated by

varying the beam incidence angle relative to the crystal through the scanning action.

The emitted or absorbed electron current varies depending on the angle between the

beam and the crystal lattice. The information carried by electron channeling contrast

is confined to a region within 50 nm of the

surface since elastic scattering of the beam

electrons rapidly randomizes and decollimates the beam with depth. Moreover, the ECP has been found to be sensitive to imperfections in the crystal. 5 ' 6 ECP's can be

obtained from areas of 10 urn diameter or

less and, thus, the technique offers the

possibility of assessing the crystal perfection of shallow layers with micrometer resolution.

In the present study, the ECP technique has

been used to qualitatively assess the damage

induced in a silicon crystal by the ion bombardment in an SIMS instrument. The ion

bombardment conditions used were: 16 0~ primary ions, beam energy 18.5 keV, beam current

10 nA, beam astigmatically focused and scanne

in an area 50 urn x 50 pm for 1000 seconds.

The resulting eroded crater is shown in fig- ure 1. A region of constant current density

was obtained over most of the scanned area,

as evidenced by the uniform erosion (figure 1 region I) observed by an optical interference

pattern. An ion flux of 2.5 x 10 18 /cm2 passei

through this region. The astigmatically

focused beam produced a gradation in ion flux

at the edge of the crater (figure 1, region

II). ECP's were obtained from the original

crystal surface, the region of limited ion

bombardment (region II) and the center of the crater by translating the specimen under the

beam. The ECP from the original surface,

figure 2, shows the fine detail and strong

line definition characteristic of highly perfect crystal. The ECP obtained from the

region of limited ion bombardment, figure 3, with identical SEM operating conditions as those used to generate figure 2, is greatly

degraded. The ECP from the center of the pit was totally lost, indicating that as far as the electron channeling effect is concerned,

the specimen is amorphous.

Additional experiments are being carried out

to relate ECP quality to the defect concentration. However, from the present qualitative comparison of an ion bombarded region

with the original crystal, it is clear that

ion erosion leaves behind a severely degraded

surface layer. This fact must be considered

when questions of solute mobility during

analysis or of the physical nature of the

region under analysis are important.
References

Carter, G. , and Colligon, J. S., Ion

Bombardment of Solids, pp. 151-156

(Elsevier, New York, 1968).

Booker, G. R. , Modern Diffraction and

Imaging Techniques in Material Science,

Amelinckx, S. , et al., ed., pp. 613-653

(North Holland Publishing Company,

Amsterdam, The Netherlands, 1970).

Newbury, D. E. in SEM/1974, Proceedings

of the 7th Annual SEM Symposium,

Johari, 0., ed. , IITRI, Chicago,

Illinois, 1974, pp. 1047-1054.

4. Goldstein, J. I. , Yakowitz, H.

Newbury, D. E. , Lifshin, E., Colby, J.,

and Coleman, J. , Practical Scanning

Electron Microscopy, pp. 150-179 (Plenum

Press, New York, 1975).

5. Schulson, E. M. , van Essen, C. G. , and

Joy, D. C. in SEM/1969, Proceedings of

the 2nd Annual SEM Symposium, Johari, 0.

ed., IITRI, Chicago, Illinois, 1969,

pp. 47-55.

6. Wolf, E. D. , and Hunsperger, R. G. in

SEM/1970, Proceedings of the 3rd Annual

SEM Symposium, Johari, 0., ed., IITRI,

Chicago, Illinois, 1970, pp. 457-463.

82
1 I

I 50pm

Figure 1. Optical micrograph of ion eroded

crater. Region I is uniformly bombarded;

Region II is lightly bombarded.

<?k>

Figure 2. Electron channeling pattern obtained from original crystal surface prior

to ion bombardment

.

Figure 3. Electron channeling pattern from

edge of crater region (Region II in figure 1)

.

13

NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

THE EFFECT OF SPECIMEN COOLING ON THE MIGRATION

OF SODIUM IN THIN FILM Si02

Bradway F. Phillips

Naval Weapons Support Center Electron Optics Laboratory

Crane, Indiana 47522

Alfred E. Austin Harold L. Hughes

Battelle Columbus Labs Naval Research Laboratory

Columbus, Ohio 43201 Washington, D.C. 20375

INTRODUCTION

Being able to profile element concentration

changes as a function of depth in thin film

oxide structures, particularly Si02 , is of vital concern to the understanding of MOS

device characteristics. Early in the development of techniques to measure these

element profiles directly using the ion

probe, it was shown that both Na and K were mobilized by the primary ion beam while Al

and P were essentially undisturbed by the

charge built up in the thin Si02 film [1],

[2], [3]. This discovery helped to explain

how a number of widely varying results could

be obtained on specimens which were sup- posedly identical in preparation since K and

Na were moving through the Si02 films quite

readily. This effect was causing Na and K surface contamination to show up at different points remote from the surface, particularly at the Si02 - Si interface.

EXPERIMENTAL CONCEPT

As a step toward providing better data on Na distributions in Si02 a number of possible experiments were proposed, including:

A. Use a primary ion beam with a low ionization potential compared to the 0~ , 02+ or Ar+ beam used regularly.

B. Increase or decrease the primary ion energy greatly compared to the nominal

5.5 keV 0 2 + or Ar+ ions or 14.5 keV 0" ions

used in the Cameca IMS 300.

C. Cool the specimen as near to liquid nitrogen temperatures as possible to slow down

the "drift" of Na under the surface charge

built up with ion bombardment

.

It was felt that all of these experiments

were important but that choice C was best

suited to the study we wished to make.

Therefore, a liquid nitrogen dewar and

specimen holder were constructed to fit the

specimen cooler port of the Cameca IMS 300. The specimen holder was designed to hold

specimens 0.010 to 0.020 inch thick, 3/8

inch wide, and up to 1.5 inches long, and

included a translation mechanism to go to any spot in a vertical line across these

specimens. This allowed us to load three

3/8 inch square silicon slices, in addition

to the copper grid pressed into aluminum for

aligning the instrument and making sure the

secondary ion optics were properly focussed.

Specimen cooling rates and ultimate temperatures were measured several times, and the averaged temperature value obtained at the

spot being analyzed was -180°C.

EXPERIMENTAL RESULTS AND ANALYSIS

As a reference to the work which lead to this study we see the results of profiling

the Na distribution in Si02 at room temperature in figure 1, which was mentioned pre- viously today in Bob Dobrott's talk and

which is typical of analyses we were seeing

in 1970, 1971, and early 1972. According to the predictions of the LSS theory we would

expect the sodium to peak at 350 A from the

surface and obviously it is not doing that. Therefore, for our first room temperature

and low temperature experiments we chose a series of specimens, which ranged from as- grown oxides to samples containing several

levels of aluminum implantation and residual

sodium. Results of these analyses are presented in table 1. First, we have an as- grown oxide analyzed at -180°C, in which 27% of the sodium concentrated at the Si02~Si

interface; then at room temperature, where

70% of the sodium concentrated at the interface. The remainder of the specimens are aluminum implantations performed at 20 and

25 keV ranging in three steps from 4 x 10 15 to 6.8 x 10 12 atoms per square centimeter

total dose. As we can see with the lower

implant dose, NRR 371, we have at -180°C the

smallest amount of sodium (32%) moving to the Si0 2 -Si interface of this group of

specimens. However, as we go to the higher

dose specimen, NRR 372, 45% of the sodium moves and at the much higher dose, NRR 437,

46% of the sodium moves, where at room temperature these values are 75, 88, and 90%

respectively. So we can see that there is a definite effect produced by using the low

temperature specimen holder. As an example

of the profiles obtained, figure 2 shows the

results of analyzing an as-grown oxide at
room temperature. The silicon secondary ion

yield changes by several orders of magnitude

as the interface is penetrated. At all

points in the specimen except at the SiC>2-Si

interface the sodium profile shows a count rate which is essentially at the background

noise level of the instrument.

In figure 3 we see the results of cooling

the same specimen to -180°C and this analysis demonstrates that the drift of sodium is

retarded. The instrumental conditions were

kept exactly the same in both cases. Here

we see that when the specimen is cooled the

sodium does not move rapidly to the SiC^-Si

interface during analysis, but is retarded;

in fact, the sodium concentration is two orders of magnitude higher throughout the

oxide in the spot analyzed on the cooled

specimen than it is in the spot analyzed at room temperature. Although this technique

does not yet give the true sodium profile,

it does give some encouragement to continue

on this path of analysis, and table 2 describes the next set of specimens which were

prepared. The first specimen, JR1C, the as- grown oxide control sample, was cut up and

used as a base material for all of the so- dium ion implants which ranged from

3.2 x 10 11+ to 1.4 x 10 15 atoms per square

centimeter total dose; some of which were

annealed at 200 or 900°C. The results obtained on these samples are tabulated in

table 3. In specimen JR1C analyzed at 25°C

with the 0~ beam, 100% of the sodium

appears at the surface of the oxide; at -180°C with the 0" beam, 97.5% of the

sodium appears at the surface. Therefore, we see that in the control specimen essentially all of the sodium is present at the

surface of the oxide as contamination. The

next experiment on JR1C using the 02+ beam

demonstrates that we still have a fair

amount of the sodium mobilized both at 25°C

and at -180°C. At 25°C, 10% of the sodium

is at the surface and 89% is at the Si02~Si

interface. Going to -180°C on the same sample, we see that 54% is on the surface of

the oxide, 26% at the interface, and 20% is

included within the bulk oxide. So we have

retarded the movement of the sodium although

we do not feel that we have a perfect profile yet. Working with specimen JR1A1,

which was implanted to 3.2 x 10 lt+ atoms per

centimeter and annealed at 200 °C, we see

that there is a very small amount of the so- dium at the surface in both the -180°C and

the 25°C specimens. Most of the sodium has

gone to the interface. We interpret this as being caused primarily by the ion induced

channels from the implantation process and

by the fact that the 200°C treatment does

not anneal out these channels. We still

have an inducement for the sodium to move, even though we have cooled the sample.

JR1A2 has been annealed at 900°C for 30 minutes, and in this case we see at -180°C that

29% of the sodium is on the surface. 65% is at the interface. Essentially the same appearance was seen at 25°C, showing that

there must have been some prior movement of

the sodium at room temperature and then at 900 or 200°C before we did the analyses.

In order to show the differences a little

bit more clearly we have made up differential plots between the work done at room temperature and the work done at low temperature. Figure 4 shows the results obtained

with a differential plot of JR1C, our control specimen. This plot demonstrates that

cooling holds some Na at the surface of the

Si02, retards its diffusion through the

Si02 , and greatly reduces the pileup at the

Si02-Si interface. In figure 5 we are looking at the JR1A1, which is an implanted and

low temperature-annealed sample. We see

that there is a sizeable portion of the so- dium present which is retained at the surface by cooling the specimen during analysis

In the sodium profile through the sample

there is a large decrease in the peak right

before going into the silicon substrate and some reproducible small peaks as you sputter into the silicon itself that must be

caused by interface states and the differences in sodium activity between 25°C and

-180°C. In figure 6 we are looking at the

differential plot of JR1A2, which was an- nealed at 900°C for 30 minutes. Here we see somewhat the same results as with JR1A1 except that there is a wide peak of the so- dium at the surface compared to the specimen

analyzed at room temperature and a very deep

valley when we get close to the interface.

The integrated sodium intensities of the

two JR1A2 specimens were 26,000 and 29,800

counts, compared to 54,000 and 44,000 for

the JR1A1 specimens. This difference may be

due to the loss of sodium from the surface

during the high temperature annealing process. Such loss caused by diffusion of so- dium to the surface and partial escape into

the furnace atmosphere would also account

for the increased surface concentrations of

sodium in the JR1A2 specimens, 22 to 29%

compared to 6 to 7% of the JR1A1 specimens.

That the sodium concentration at the surfaces of the JR1A1 specimens is lower than

would be expected for the shallow 200 A range suggests that the sodium ions may be

driven into the Si02~Si interface by surface

charging during the ion implantation.
A run on the JR1A1 specimen with O

- primary

beam at -180°C showed a high but continuously decreasing sodium concentration through

the oxide, and no sodium at the Si02~Si interface. As with the control, the highenergy 0~ beam apparently makes the sodium

more mobile in the oxide. We see in our

third example, JR1A2, a fairly uniform distribution of the sodium throughout the Si02 , with the peak reduced at the interface and a peak showing at the surface. Therefore, we

can conclude that when a specimen which has

little implantation damage is analyzed at -180°C, we can get a much better profile of

the sodium which we would say does exist on

the surface of the SiC>2 as contamination

from the growth technique, and throughout

the oxide from the implantation, than we can

at room temperature.

SUMMARY

In summary the following points have been

established

.

1. A liquid nitrogen-cooled specimen holder

was fabricated for a Cameca IMS 300.

2. Cooling thin-film Si02 specimens does

retard the diffusion of Na during ion probe

analyses

.

Although it does not give a perfect profile,

such cooling is a step in the direction of

attaining sodium profiles that are reproducible and real.

3. The interactions of primary beam energy

and species, ion implantation dose and

energy, annealing temperature, and a variety

of other factors not yet known, combine to affect the Na profiles obtained and to improve or degrade the benefits of cooling the

specimens

.

ACKNOWLEDGEMENTS

We would like to acknowledge the contributions of Mr. Ronald Baxter of Leeds and

Northrop Company, North Wales, Pennsylvania

in the design and initiation of these ex- periments, and the sponsorship of the Defense

Nuclear Agency under subtask Z99QAXTD033,

work unit 51, work unit title, "Radiation

Resistant MOS Technology," technical monitor, Mr. H. L. Hughes, of the U. S. Naval

Research Laboratory.

REFERENCES

1. H. L. Hughes, R. D. Baxter, and B. F. Phillips, IEEE Trans. Nue. Soi. NS-19 , 256 (1972).

2. D. V. McCaughan, and V. T. Murphy, IEEE

Trans. Bug. Sci. NS-19 , 249 (1972).

3. R. D. Baxter, Investigation of Thermally

Grown Si02 by Ion Microanalysis,

Battelle Memorial Institute (1971).

II
TABLE 1

Na DISTRIBUTIONS AT -180°C VS 25°C

(a)

Specimen

Ion Implantation

% Na at Si02 /Si Interface

-180 C 25 C

NRL-434 None 27 70

NRL-436 25 keV Al to 4 X 10 15 cm" 2 46 90

NRL-371 20 keV Al to 6 8 x 10 12 cm

-2 32 75

NRL-372 20 keV Al to 5 X 10 13 cm" 2 45 88

(a) Sputtering by 0* primary beam

TABLE 2

SPECIMENS FOR Na DISTRIBUTION MEASUREMENT

AT BOTH 25°C AND -180°C

1) JRIC - as grown

2) JR1A1 - Implant 10 keV Na 23 to 3.2 x 10 14 cm" 2 annealed

200 C 30 minutes

3) JR1A2 - Implant 10 keV Na 23 to 3.2 x 10 14 cm' 2 annealed

900 C 30 minutes

4) JR1B1 - Implant 30 keV Na 23 to 1.0 x 10 15 cm" 2 annealed

200 C 30 minutes

5) JR1B2 - Implant 30 keV Na 23 to 1.0 x 10 15 cm" 2 annealed

900 C 30 minutes
TABLE 3

Na DISTRIBUTIONS AT -180 C VS 25 C

Percent of Na in

Temp. Primary Surface Bulk Si0 2 /Si

Specimen C Beam of Oxide Oxide (interface) Si

JR1C 25 0" 100

JR1C -180 0" 97.5

JR1C 25

4- 0+

2 10.7 0.23 89

JR1C -180 0+

2 54 20 26

JR1A1 25 0+

2 7.2 2.3 90.5

JR1A1 -180

J. 0+

2 6 24 70

JR1A2 25 0+

2 22 1 77

JR1A2 -180 0+

2 29 6 65

JR1B1 25 o+

2

5 6 74 15

JR1B1 -180 o+

2

3 64 26 7

JR1B2 25 0+

2

21 4 74 1

JR1B2 -180 0+

2 25 40 34 1
SILICON DIOXIDE (3100 I) ON SILICON Na* IMPLANTED <20k»V, IxlO15 ionl/cm*)

• '*0~ PRIMARY BEAM

"O* PRIMARY BEAM {

\

SIOj Si

.1 I L

O 500 1000 I5C0 2000 2500 JOOO S500 DEPTH, I

LU 10

CO —I

I- <

z o

d tn 8 o

Q ,„2

Figure 1. Sodium profiles of a sodium - implanted

silicon dioxide film for both 16 0+ and 16 0~.

104

zo

Z><0 Oo

» 30 Si • 23 Na

Wax*'*!

•

•

• •

• • • '•• «•* *•• ' * •< •

•

• •

•

•

•

RELATIVE DEPTH

Figure 2. Sodium profile at 25°C of an as-grown

oxide.

X 30Si

• 23 Na

xxxxxxxxXXXXXXXXXX,xxxxxxxxx^ xxxxxxxxxxxxxxxxxxxx xxxxxxxxxx "*x

X

X .'. X . y

*•••

K

• »x »»xx«,xXXXXXXXxXX (XXXxXX»»J

• •

•

0 20 40 60 80 100 120 140 160 180 200 RELATIVE DEPTH

Figure 3. Sodium profile at -180°C of the same specimen as in figure 2.
eoOp

700 I

600

500

400

0LO 300 - CN

200 -

Na 100

u -100 -

0O

00 -200 - H

1 -300

nj S3 --•GO -

-500 -

-600 -

-700 -

-800 1—

~i i r

J 1 L

Figure 4. Plot of Na (-18.0°C) - Na (25°C)

distribution in JR1C.

30 40 50 60 70 60

Chonnol

Figure 5. Plot of Na (-180°C)

distribution in JR1A1.

Na (25°C)

IB00

1600

1400

1200

1000 u coo oLO

CM 600

400

200

1 0 O

0 -200 O

00

t-H

1

-400

-600

-800

-IOC0

-1200

-K00

-1600

lo 26001 "1 1 1

1 r

go so ico no

Figure 6. Plot of Na (-180°C) - Na (25°C)

distribution in JR1A2

.

71
DISCUSSION OF THE PAPER

Participant : What was your primary beam

current density?

Phillips : The primary ion beam intensity

was approximately 250 nanoamps rastered out

to cover an 800 micron square.

Cock : For one of your samples you postulated that the sodium motion was enhanced by

dislocations. Would you could elaborate on

that model.

Phillips : I would say that at the 200°C

anneal, you have not removed the channels

that are caused by the ion implantation

process itself. At the 900°C anneal you

have removed some of the channels but you

have also mobilized the sodium. Data, which

we did not present today, using the picture

mode in the Cameca IMS 300 to study the

location of the sodium after a 200°C anneal

and after a 900°C anneal shows the difference.

At 900°C you have mobilized the sodium and

you actually have it diffusing back through

the channels to the surface and escaping

into the furnace atmosphere, making your

total sodium content in the entire oxide

half as large as it is with the 200°C anneal.

Cock : I see, so this is a dislocation in

the oxide structure.

Phillips : Yes. If you implant sodium to a high enough intensity in the oxide structure,

somewhere around 1 x 10-L5 atoms per square

centimeter, you can provide a state in the

Si02 which is of such a high energy that

the silicon and sodium, - in fact, all ions

- coming off the specimen from a section

extending from the surface to close to the

depth at which the maximum of the implantation peak should occur, are of such a high

energy compared to the secondary ion band

pass of the instrument that you do not see any secondary ions. So you know that you

are producing an extremely strained state in

the Si02 film that does shift even your

secondary-ion energy distribution.

Participant : The temperature of 180°C that

you quote is, I gather, a bulk sample

temperature

Phillips : That is correct. The microtemperature at the sputtering spot has not been

determined. The temperature at which the

sputtering actually takes place, or the

plasma in which the ions are produced, would

be significant for the theory devised by Dr.

C. A. Andersen. The bulk temperature has

been determined to be 180°C.

Participant : Although the temperature at the sputtering location has not been determined do you have any idea of approximately

how large of a difference you might expect?

Phillips : Well, in the plasma itself, or at the place at which the sputtering is occurring, the temperature is several thousands

of degree Kelvin.

Evans : You are talking of the plasma or electronic temperature, not the thermal tem- perature. The thermal rise at the surface

due to ion bombardment is only of the order

of 5 to 50°C above the bulk sample temperature. The plasma "temperatures" are thermal

spikes: they last the order of 10 seconds; they are not a temperature in the true

sense, they are an electronic temperature.

Phillips : The temperature rise is extremely

small because we are dealing with a fairly

heat-conductive medium, silicon, and a fairly

thin film of Si02 with a massive heat sink

compared to the size of the sample. So I would say that even going through the Si02

there would be an extremely small temperature

rise in the local area surrounding the

sputtering; of the order of a few degrees at the most.

12
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg , Maryland, April 23-24, 1975 (Issued March 1976)

SILICON-ON-SAPPHIRE IMPURITY ANALYSIS 5

'1

D. Howard Phillips

Electronics Research Division

Rockwell International Corporation

Anaheim, California 92803

INTRODUCTION

The presence of undesirable impurities had

detrimental effects on both the electrical

and radiation performance of silicon-on-sapphire (SOS) integrated circuits. Conversely,

device electrical and radiation characteristics can be improved by minimizing the number and types of impurities in SOS films.

Ion microprobe mass analysis (IMMA) techniques are being utilized for impurity analyses in a program associated with the development of a radiation-hardened CMOS/SOS (complementary MOS on sapphire) technology. From

an applications point of view, four goals

have been established for this program. The

first is to improve lot-to-lot repeatability;

the second is to improve wafer-to-wafer re- peatability; the third is to improve the

electrical stability of the devices (i.e.,

the bias-temperature stability of threshold

voltage) ; and the fourth is to improve the radiation stability of the devices.

IMPURITY SOURCES

One source of contaminants is the sapphire

substrate. Impurities are sometimes added

during the growth process to minimize crystal

defects such as twinning.

A second source of impurities is the epitaxial film itself. During the epitaxial film

growth process, impurities come from silane,

hydrogen carrier gas, and dopant gases.

A potential major source of impurities is the processing steps that are used during

the wafer fabrication process.

MASS-ANALYSIS SILICON I MPURITY DATA

Figure 1 shows recent data comparing siliconon-sapphire films with an epitaxial silicon

layer on bulk silicon. The shaded areas indicate the impurity concentration measured

in silicon films on a sapphire substrate

grown using the Czochralski process. The

* Partially sponsored by USAF/AFCRL (Contract

F19628-75-C-0108) and by Rockwell Inter- national Corporation (IR&D)

.

unshaded lines on the bar graph indicate the impurity concentrations that were measured

on epitaxial films grown on bulk silicon.

The data are plotted as a function of elemental impurities versus the relative con- centration, where each increment on the horizontal axis indicates an order-of-magnitude

change in impurity concentration.

These data show that aluminum concentrations

in silicon-on-sapphire are orders of magnitude greater than aluminum concentrations in

bulk silicon. In the case of sapphire

(AI2O3) , there is more aluminum present in the system to act as a potential contaminant,

either from top-side diffusion or from backside diffusion out of the sapphire substrate.

Another very important observation from these

data is a consideration of sodium content. Sodium and potassium are known sources of electrical instability in MOS transistors.

Compared to bulk silicon, typical silicon on sapphire contains much higher sodium concentrations. The data in Figure 1 show very little difference between silicon-on-sapphire

epitaxial films and bulk silicon epitaxial

films, in terms of potassium impurity con- centration .

DIFFUSION PROFILES IN SILICON ON SAPPHIRE

IMMA analysis techniques are being used in the analysis of intentional dopants. Figure

2 shows phosphorous concentration as a function of diffusion drive-in time. The drivein time at 1,000 degrees centlgrate was varied from 0 to 15, 30, 60 and 80 minutes.

Initially (at T = 0) , phosphorous was con- centrated close to the surface. Following

high-temperature processing, the profiles

changed as shown in Figure 2. We typically

use silicon-on-sapphire material with a silicon epitaxial thickness in the range from

0.8 to 1.0 micron. For typical wafer-processing thermal cycles, phosphorous diffuses

completely through the silicon-on-sapphire

film. The diffusion coefficient in sapphire

is very low and phosphorous does not diffuse

through the silicon-sapphire interface. The

sapphire interface acts as a diffusion stop

causing phosphorous to remain in the silicon

film.
MASS ANALYSIS DATA ACQUISITION

Figure 3 illustrates the instrument used to obtain impurity data presented in Figures 1 and 2. The ion microprobe mass analyzer

generates an ionized beam that is swept electrostatically both in the Y direction and

the X direction in much the same manner that

a TV raster is generated. The ion source is

allowed to impinge on the sample of interest.

Beam diameters can be adjusted from 2 microns

to 500 microns. The rate at which the surface is eroded depends on several experimental factors, i.e., the ion source, the intensity of the ion beam, the energy of the

ion beam and the nature of the sample that

is being analyzed. Erosion rates vary over

a range from less than one angstrom per second up to 1000 angstroms per second. The

secondary ions are analyzed electrostatically

(momentum) and magnetically (mass-to-charge

ratio) . The ions impinge on a target where

an electron beam is produced. The electron

beam generates light in the scintillator,

and this light is coupled to a photomultiplier tube. The signals from the photomultiplier tube can be recorded and displayed in

several ways. One way is to record them on

a CRT which utilizes the signal from the

photomultiplier tube for Z axis modulation.

CRT Y and X signals are derived from the

waveforms used to deflect the primary ion

beam. Data can also be recorded on a stripchart recorder by using the signal from the

photomultiplier tube to modulate the Y axis

of the recorder. Data can be recorded by

using a scaler to count pulses from the

photomultiplier tube. Data can be recorded

in the MHz range without significant data

loss due to scaler dead time.

SPECTROCHEMICAL-ANALYSIS SAPPHIRE IMPURITY

DATA

Table I lists data showing impurity levels

in the sapphire substrate, rather than the

silicon film. These data were obtained using

spectrochemical analysis techniques rather

than the ion microprobe mass analysis technique just described. These data were obtained by using Crystal Systems sapphire

which was grown using the gradient furnace

technique. The gradient furnace technique

is one of three techniques that is used com- mercially to grow sapphire crystals that can

be used to fabricate silicon-on-sapphire

integrated circuits. Table I compares pre1971 sapphire impurity data with data for

the improved process that is being used at this time by Crystal Systems in the growth

of gradient furnace sapphire. The number of

impurities present to a significant extent

are greater for the improved process, but

the impurity concentrations, in general, are much less than for the sapphire-growth process used by Crystal Systems prior to 1971.

MATERIAL COMPARISON AND DATA CORRELATION

Considering the different analysis techniques

that have been used to analyze both sapphire

and silicon films, data correlation is needed.

Figure 4 indicates some of the work that is

being done to correlate the data that have

been accumulated. Data are being obtained

for three types of sapphire substrates

Czochralski-grown sapphire substrates fabricated by Union Carbide, ribbon sapphire substrates fabricated by Tyco, and gradientfurnace sapphire substrates fabricated by

Crystal Systems. These sapphire substrate

growth techniques are quite different, and

preliminary indications are that the impurity

concentrations are quite different for these

three techniques.

Figure 4 describes a correlation study that

is being performed as a cooperative effort

with NBS and Rockwell participating in the

data acquisition activities. This study includes the analysis of silicon-on-sapphire

wafers from three sources. Each wafer will

be divided into halves. One-half of the

sapphire wafer will be analyzed using the

flame emission spectrometry techniques in use at the National Bureau of Standards.

The other half of the wafer will be analyzed

by Rockwell using the ARL-IMMA analysis

technique. An attempt will be made to cor- relate the results of these two measurement

methods with the objective of learning more

about the sapphire substrates and the silicon epitaxial films, and using that information in our attempts to improve the material

characteristics and the device electrical

and radiation characteristics. Wafers from

the same lot will be used in integrated circuit fabrication. These integrated circuits

will be characterized to determine the electrical characteristics and the radiation

performance of SOS integrated-circuit devices. The data from the integrated circuit

fabrication and test activities will also be

correlated with the analysis of the SOS

starting material using the flame emission

spectrometry and IMMA analysis techniques.

At the same time, we have other silicon-onsapphire integrated-circuit fabrication and

characterization programs that are continually providing additional data that can be

used for correlation.

SUMMARY AND CONCLUSIONS

Unintentional impurities in silicon-on-sapphire films and substrates are relatively

74
high, based on mass analysis data. Concentrations of the unintentional impurities are equal to, or greater than, the dopant concentration of the most lightly doped region of

the semiconductor.

There is a significant difference between the

impurity concentrations in typical SOS devices or SOS wafers as compared to bulk silicon wafers.

Preliminary data indicate that there is a good deal of impurity-concentration variability depending upon the sapphire source

and the sapphire growth technique. In addition, considerable lot-to-lot and wafer-towafer variability has been observed in silicon films and has been attributed to growth

conditions and device fabrication methods,

as well as impurity concentration variations.

75
TABLE 1

SPECTROCHEMICAL ANALYSIS

CRYSTAL SYSTEMS SAPPHIRE DATA

GRADIENT-FURNACE SAPPHIRE

IMPROVED PROCESS (1971 1975) EARLY PROCESS (PRE-1971)

Element Concentration Element Concentration ELEMENT CONCENTRATION (ppm)

(ppm) (ppm)

Cr 0.4 N 3

Cr 1 Cu 2 F 0.3 Cu 10

Fe 8 Na 3 Fe 1 Mg 0.3 P 0.05 Mg 1 Mo 1 K 3 Mo 10-100

Ti 0.3 Ca 0.7 Pb 1 Y 0.5 V 0.04 Si 10-100

Sc 4 Mn 0.7 Ti 10100

Ce 6 Ni 0.2 Elements sought and not detected (< 0.05 ppma):

B 0.02 Zn 0.6 Ag, As, Au, B, Ba, Be, Bi, Ca, Cd, Co. Ga, Ge, Hf, Hg, In, Ir.K, Li, Mn. Na Nb Ni Os P Pd Pt Rb Re. Rh Ru. Sb Sn. Sr Ta. Te. Ti. Nominal detection limit = U.Ub ppma.

V, W, Zn, Zr.
RELATIVE ELEMENTAL SCANS OF EPITAXIAL S I LICON ON BULK

SILICON AND SAPPHIRE SUBSTRATES

B

C

Na

Mg

Al

K

Ca

Cr

Fe

Mn

Cu

Tzzzzzm

mZTZZL

7ZZL Z ;y,' 7 ////ay/,;;;; aaaI

'

1

IMMA DATA 0 Si ON SAPPHIRE (CZ)

Si ON BULK

ZZZZZZZZZZZZZZZ 7Z2 22 ' 2 .'-V/V; >>w?w//??

x

io* 10x+2 10X+3 RELATIVE CONCENTRATION

10*+4 10x+5

Figure 1. Relative elemental scans of epitaxial silicon on bulk

silicon and sapphire substrates.

OTHER SOS APPLICATIONS OF IMMA

PHOSPHOROUS CONCENTRATION VS

PROCESSING TEMPERATURE: 1000°C

4 6 DEPTH X 103 A

Figure 2. Phosphorous diffusion profiles in silicon on sapphire.

77
ION MI CROPRORE MASS ANALYZER (IMMA)

Figure 3. Ion microprobe mass analyzer (IMMA)

CORRELATION

CORRELATION APPROACH

• Czochralski • Ribbon

•Grad-Furnace

FES

(NBS)

WAFERS FROM

SAME LOT

IMMA (ROCKWELL

CORRELATION

TEST

• ELEC •RAD

OTHER SOS CHARACTERIZATION PROGRAMS

Figure 4. Material comparison and data correlation.

18
DISCUSSION OF THE PAPER

Participant : I have a comment and a question. Without having seen your original

spectrum of the epitaxial silicon/bulk sili- con composite, what I suggest is that your

detection of species such as chromium,

manganese, iron and copper may in reality be

the following: at mass 52, since you have a large carbon peak, it could be SiC2+ 5 at mass 55 instead of manganese you may have

AISi; at mass 56 you may have Si2» and since

you have a large sodium peak this may not be

out of the question, and at mass 63, 65 the

polyatomic peak SiCl. Now if your energy

bandwidth is broad, you are going to be in

trouble with polyatomic species, and therefore I suggest you look once again to the

possibility of interferences.

The question is the following: In the profiling of phosphorus and silicon how do you

discriminate between SiH at mass 31 and 31 P especially at low concentrations?

Phillips : Your comments are well taken.

Although I cannot give you a comprehensive

answer to your question, the need for correlation experiments is recognized. As I indicated, mass analysis impurity data are

being correlated with impurity data from

flame emission spectroscopy measurements.

Morabito : If you did not have this chemical

technique you would not know from the SIMS

technique what these impurities are because

of these mass resolution problems. Although

we have looked at phosphorus and silicon and_

it is not impossible to resolve SiH from P . The technique there is to look at the P0~

species

.

Dobrott : I have a few comments on applications with the IMMA. At that phosphorus

level the SiH is nearly completely absent

from the spectra provided the coldplate was

run to gather most of the water. The SiH

will not get you into trouble until you get

much below the 10 17 /cm 3 region for phosphorus. The mass 52 which you identified

as Cr could be SiC2> however if that were

the case you would see a much stronger mass

40 peak and also the mass 50 peak of Cr would be absent. You are probably perfectly

right that mass 56 corresponds to iron

because if you had much Al you would mask

the iron 54. However, again, the intensity

of molecular species is down considerably

from what the monotonic species is, so you

could get some feel of iron. Now when it comes to the copper of mass 63 and 65, it is true that silicon 28, together with chlorine

35 and 37 would dominate at 63 and 65, and

unfortunately as nature would have it, the

isotope ratio is about the same as what you

would expect for copper. But it is unbelievable to me that he could take an epitaxial

grown silicon and have enough chlorine incorporated into it that it would combine

with the silicon and still come out to anything detectable by the IMMA.

The only other comment that I have is more

or less a question. It was stated that the

concentration level of anything which was

not detected was below 0.05 ppma. I do not believe that this value is quite correct

since a considerable number of noble ions were present. I would say the correct

value that I would put on that with confidence is less than 10 parts per million

unless you did specific long term counting

on those particular ions.

Phillips : One comment on chromium and

titanium. Often times those two elements

are intentionally added to the melt during

sapphire growth for the purpose of minimizing the number of defects in the crystal as it is grown. It has been experimentally observed that the number of twins per unit

volume can be reduced if either or both of

those two elements are added in reasonably

large trace quantities.

n
II
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

SURFACE COMPOSITION BY ANALYSIS OF IMPACT RADIATION

Clark W. White

Bell Laboratories

Murray Hill, New Jersey 07974

Several speakers at this conference will discuss analysis techniques that are based on phenomena that occur when the low energy ions

impact on solids. These include the techniques of low energy ion backscattering and

secondary ion mass analysis, both of which

utilize beams of low energy ions to bombard

the surface of interest. In addition to the

backscattering of the ion projectile and the

secondary emission of positive and negative

ions, one also observes the production of intense optical radiation in the low energy particle-solid collision process. In collaboration with Norman Tolk and Douglas

Simms, we have been studying the radiation

that is produced in collisions of low energy

ions and neutral atoms with solids. These

studies give fundamental information on the

low energy particle-solid collision phenomena and at the same time provides us with

the basis for a sensitive technique to identify the constituents of the surface and

near surface regions. The radiation that

will be discussed is produced as a result of the sputtering process and, therefore, if sputtering is being used for any purpose

these optical signals and the optical information can then be used to serve as a monitor

for the sputtering process.

Figure 1 schematically illustrates the type

of experimental apparatus which has been used

in this work. This apparatus produces beams

of ions or neutral atoms of well-defined en- ergy in the energy range from about 50 eV to about 8 keV. The neutral beam is used to impact on insulating surfaces to avoid any problems that are associated with charge ac- cumulation on the nonconducting target. In the target chamber, the target to be bombarded is oriented with the surface normal

making an angle of from 45 to 60 degrees

with respect to the beam direction, and photons that are produced in the collision process pass through a quartz window and are focused by a quartz lens into a small, fast

monochrometer . The monochrometer and the photomultiplier are used to record the spectral distribution of radiation that is pro- duced in the collision process and single

photon counting techniques are used for sen- sitive radiation detection.

Table 1 lists at least three possible

sources or mechanisms for producing optical

radiation in the low-energy collision region.

First of all, there is radiation from ex- cited states of atoms and molecules that

have been sputtered off the surface. Basically, interaction of the beam with a solid

results in sputtering of atoms and molecules

off the surface. A significant portion of the sputtered fragments leave the solid in excited electronic states which then decay

and give rise to lines and bands that are characteristic of surface constituents including surface contaminants. Secondly,

there is radiation from excited states of backscattered beam particles. This obviously

gives information on the quantum mechanical

states of the backscattered projectiles.

Then, finally, there is radiation which re- sults from the excitation of the electrons

in the solid. These, in general, are observed to be very broad continuum of radiation which are produced most efficiently by

high velocity projectiles impinging on insulating targets. The radiative continuum

that are observed here are localized to the

solid and result from the transfer of pro- jectile energy to the electrons in the solid.

This gives rise to the creation of electron^

hole pairs which then produce the radiative

continuum upon recombination. The first

mechanism provides the basis for a technique

to identify surface constituents. This is the technique which we call the SCANIIR technique which stands for Surface Composition

by Analysis of Neutral and Ion Impact Radiation. As an analysis technique this method

is very sensitive to the first few monolayers

of the solid. The spectrum of radiation that one obtains is relatively simple and uncom- plicated, but there are effects due to non- radiative de-excitation processes that one

has to be aware of.

Figure 2 shows examples of the spectrum of radiation that are produced when argon ions

at an energy of 4 keV impinge on surfaces of copper and nickel. The prominent lines that are shown in these two spectra arise from

low lying excited states of neutral copper

and neutral nickel. Therefore, we are detecting the sputtering of neutral atoms from

the surface by detecting the radiation that

they emit when they leave the surface. In addition to radiation from copper and nickel, we also see radiation from common surface

contaminants such as sodium, from the sodium

D lines, as well as organic contaminants

from a molecular band that arises from an

8!
excited state of the CH molecule. There is

also radiation from hydrogen, which presumably results from either the fragmentation

of the organic contaminants or possibly water

vapor. The molecular band from the CH molecule is the most intense radiative feature

that is seen from any of the common organic

contaminants, and it serves as a unique fingerprint to identify specifically the fact

that one has organic contamination on a surface. We have used this on numerous occasions to look for such things as photoresist

residues on elemental surfaces.

Figure 3 shows results that were obtained

using the neutral beam on three insulating

targets: sapphire, lithium fluoride, and

fused quartz. The optical lines that are shown here arise from low lying excited

states of aluminum, lithium and silicon, re- spectively. The widths of all of these

lines, at least the prominent ones, have

been measured to be less than or

o equal to the instrumental resolution (^1 A) and this

reinforces the conclusion that these lines

originate from atoms which are radiating after they leave the surface. Experimentally we find that the photon production efficiency

is substantially greater for these insulating

targets than in the case of metal targets.

In the case of metals, a typical prominent

line will be produced with an efficiency of

approximately 10

-

^ to 10~5 photons produced

per incident projectile. In the case of insulators, these same optical lines are produced with an efficiency that is measured to be some two to three orders of magnitude

higher.

A plausible explanation for the large difference in the excitation efficiency on metal

and insulating targets can be found by con- sidering the effects of nonradiative deexcitation. When an atom is in the vicinity

of a solid, there are nonradiative electronic

processes which result from the interaction

of the excited atom with the solid which can very efficiently compete with radiative decay of the excited atomic states. In general, there are two basic types of nonradiative processes, and they are illustrated

schematically in Fig. 4. This shows a potential well diagram that is appropriate to a metal with an excited atom located a distance

S away from the surface. If the excited

atomic state lies above the Fermi level, then

the excited atomic electron can tunnel

through the potential barrier into one of the

unoccupied levels. This is a one electron

resonance tunneling process and is known as resonance ionization. However, if the ex- cited atomic state lies below the Fermi level

(Case B) , then tunneling is prohibited since

the levels are filled. Even in that case

the excitation energy can still be transferred nonradiatively to one of the available

electrons in the solid and the electron from

the solid may or may not appear as a secondary electron; it depends on the excitation

energy of the atomic state and the depth in

the well from which the upper electron was

drawn. This is a two electron nonradiative

process and is known as Auger de-excitation.

For both of these processes the nonradiative

de-excitation rate is a very strong function

of the distance of the excited atom from the

surface, and it can be shown that these pro- cesses will preferentially de-excite those

excited atoms which leave the surface with

low velocities. This leaves only the small

fraction ejected at higher velocities to efficiently contribute to the radiation. However, the band structure of a typical insulator is such that in many cases these non- radiative processes are not energetically

allowed. Under these conditions, then, we expect to see radiation from all excited

sputtered atoms even those that leave the

surface with low velocities. We believe then

that this accounts, at least in part, for the

very large difference in the optical excitation that one commonly observes on insulating

targets as compared to metal targets.

A good example of the effects of these non- radiative processes is illustrated in Fig. 5 where we have profiled a composite structure

consisting of a 1200 angstrom thick film of

Si02 on an infinite thickness of silicon.

We have done this by measuring the intensity

of a prominent silicon optical line as a function of the bombardment time. These

measurements have been normalized to unity

and then plotted on a logarithmic scale.

The location of the Si02 silicon interface

region is determined by that time at which

the very rapid decrease in the intensity of

the silicon optical line occurs. We see a factor of 50 reduction in the intensity of

this optical line as we sputter from the

S102 film into the silicon substrate. The

explanation for the results of Fig. 5 follows

from the consideration of Fig. 6. Figure 6 shows the same type of potential well diagram, this time appropriate to an excited

silicon atom in the vicinity of a silicon

target and an SiC>2 target. As illustrated

here, only in the case of the silicon substrate are the nonradiative de-excitation

processes energetically allowed. Therefore

in the silicon case, we expect to see radiation contributed efficiently only by those

atoms which are leaving the silicon target

with high velocities; only those atoms that

lie in the high velocity tail. In the case

B2
of Si02» because of the very large 8 or 9 eV bandgap it is energetically impossible for

these nonradiative processes to take place.

Under those conditions we expect to see radiation from all excited sputtered atoms, even

those that leave the surface with low velocities. And we believe that this accounts

for the large difference in the optical intensity from SiC>2 as compared to silicon.

This model basically says there should be a substantial difference in the velocity distribution of the radiating atoms, and you

might expect to see this reflected in the

width of the optical emission line due simply

to the Doppler effect, i.e., the Doppler

broadening of the emission lines. Experimental results are shown in Fig. 7 where we

have measured the emission line profiles for

the same optical line in the case of the

SiC>2 target and in the case of the silicon

target. These measurements were done at higher bombarding energies, a bombarding

energy of 80 keV, and they were done in second order to enhance any possible differences

in the emission line profiles. In the case

of the Si02 target we measure an emission

line that has a full width at half maximum o of about 1 A, which is the instrumental res- olution that we are using for these measurements. Under the exact same experimental

conditions from the silicon substrate, we measure a full width at half maximum of o about 5 A. Since the instrumental resolution

is the same 1 A then most of the width of

the optical line from the silicon target is

real and this definitely indicates that

there is a substantial difference in the

velocity distribution of the radiating

atoms. Experimentally, we find that the effects of these nonradiative processes can be

minimized by either using oxygen as a bombarding projectile or by purposely backfilling the target chamber with oxygen gas. These, of course, are techniques that historically have been used in secondary ion

mass analysis for essentially the same purposes and they drastically reduce the effects

of the nonradiative processes in the optical

emission.

To use this optical technique to identify

the constitutents in the near-surface region

of unknown solids, the procedure that we use

is to simply impact the surface with the

projectile beam, record the spectral distribution of the radiation, and identify the

prominent lines and bands. A typical result

is shown in Fig. 8. This shows the spectrum

of radiation that was produced when nitrogen

molecules impinged on two silica targets

containing various oxide impurities. These

two spectra were produced using an equivalent

neutral current of about 1x10"^ amps/cm^ and

it took about 20 minutes to accumulate this

data. We are using a scanning monochromator

and therefore the full wavelength range must

be scanned. By identifying the prominent

lines that are shown in Fig. 8, one can show

that there are common oxide impurities such

as sodium, aluminum, and calcium in both of

these. However, in one sample we have substantial optical lines from elements such as magnesium and iron, and those are not seen with any intensity in the other sample.

An indication of the sensitivity of the

optical technique was obtained by using Si02

samples containing oxide impurities which

were distributed homogeneously at known con- centrations in the Si02 matrix. These sam- ples were impacted with a beam of neutral

atoms and the signal to noise ratio of the

most prominent optical line from each impurity was measured. The detection limit

for each impurity was estimated from the

known impurity concentration and the measured signal to noise ratio. Table 2 sum- marizes results for some 10 or 12 oxide impurities in Si02- The left hand column of

Table 2 lists the oxide impurities which

were present in several different Si02 samples, and the right hand column gives the

detection limit for the corresponding impurity. The numbers in the right hand column

of Table 2 refer to the weight fraction of

the oxide impurity which is necessary to be

seen with a signal/noise ratio of one to one. An integration time of 10 sees was used and

the particle flux was approximately lxl0~6

amps/cm2. These results were obtained using

a monochromator to isolate the prominent

optical line characteristic of each impurity.

However, if one is willing to use a narrow

band interference filter to isolate the

prominent optical lines, then 2 orders of

magnitude improvement can be expected simply

because this increases the photon collection

efficiency by 2 orders of magnitude.

One other way that we have used to estimate

detection limits is to use substrates which

contain impurities distributed on the surface at known concentrations. Figure 9 shows results that were obtained using silicon samples which had chromium purposely deposited on the surface. For these measurements chromium was deposited on the surface

of three silicon wafers at different con- centrations and each of these samples was measured by 2 MeV Rutherford backscattering

to give the chromium concentration. These

range from approximately 2x10-'-^ to 2x10-'-^ chromium atoms per square centimeter. The

samples were then put into the apparatus and

the intensity of the prominent chromium
optical line was measured as a function of

bombardment time. The sample with 2xlO XD chromium atoms /cm^ was almost infinitely

thick during the time scale of these measurements. For the sample that had 2x10-^

atoms /cm^, the chromium optical line was observed with a peak signal to noise ratio of

approximately 200 to 1. The noise level was approximately 20 counts per second. Since

the chromium coverage is about one-tenth of

a monolayer on the silicon surface then one might expect to see 5 parts in 10^ of a monolayer of chromium on silicon, and that

could be improved also by using a narrow

band interference filter.

One can also use this technique to provide

depth profile information. Figure 10 shows

the profile of a composite structure con- sisting of a layer of Si02, on AI2O3, on

Si02, on an infinite thickness of silicon.

These results were obtained by measuring the

intensity of a silicon and an aluminum optical line as a function of bombardment time. In Fig. 10 you can easily distinguish the

three different interface regions, but this

is not a particularly good example because

we obviously have problems related to the

cratering phenomena. Presumably that is the reason that the aluminum optical intensity

does not decrease to the noise level. Nevertheless, I think this example illustrates

the potential for the optical technique and

it certainly shows, for example, that if

ion etching is being used for any purpose

and you are interested in milling away a given layer to an interface region, then detection of the optical radiation produced in

the collision process provides an in situ

method to know when to turn off the beam.

One other possible application for this

optical technique is to measure the range of

implanted ions in silicon dioxide. Figure

11 shows some very preliminary results for

an aluminum implant into silicon dioxide.

The implantation dose was 10 J aluminum ions

per square centimeter at an energy of 25 keV. This sample was then profiled by

measuring the intensity of the aluminum optical line and the silicon optical line as

a function of bombardment time. The time

at which the rapid decrease in the silicon

optical line occurs is a measure of the

Si02/silicon interface region. If the initial film thickness is known, then the film

thickness and time can be used to measure

the sputtering rate for the conditions under which the experiment was performed.

Knowing the sputtering rate, you can then

determine the depth of the peak in the

distribution of the implanted aluminum,

which in this case is about 350 A. From

the full width at half maximum you can estimate the projected standard deviation

which in this case is about 160 A. The gaussian profile that one would expect is not indicated in Fig. 11. It has been fit and

there is very little deviation until you get

into the wings of the distribution and then

it does start to deviate somewhat from the

gaussian profile presumably due to effects

associated with recoil implantation.

1 would like to conclude by indicating the

results of an experiment that was done to determine how impurity profiles of a mobile

species such as sodium, which we have already heard about this morning, can be modified as a result of several different types

of heavy particle bombardment. These experiments were done in collaboration with

Richard Kushner and D. V. McCaughan. Previous work by these two had shown that if one took Si02 films on silicon, purposely coated

on the outside surface with sodium, and bombarded the contaminated film with positive

ions, then a large fraction of the sodium

which was initially on the surface of the

Si02 film was transferred as a result of the

bombardment to the Si02~silicon interface

region. They observed that up to 10 to 20%

of all the sodium could be driven through

the oxide film to the interface, simply as a result of the low energy ion bombardment.

They attributed this transport to effects

associated with the neutralization of the

incoming ion at the Si02 surface. This, of course, gives rise to a positive charge on

the outside surface which then provides a source of driving electric field to move the

sodium to the interface. The experiment we

did was to compare sodium transport in Si02

films subjected to ion and neutral particle

bombardment. The results are shown in Fig.

12. The films that were used in this work

were 5,000 angstrom Si02 films on silicon,

purposely contaminated on the outside surface with approximately 10^ sodium atoms/

cm , the sodium contained 22 Na as a radio

tracer. The slices were then bombarded with

nitrogen ions at 2 keV, nitrogen neutrals at

2 keV, and a third case was to use nitrogen

ions at 2 keV but to flood the surface of the

S102 film with thermal electrons from a heated filament during bombardment . This

was done in an effort to keep the surface

neutral during bombardment. Each film was impacted with a particle dose of vLO^

heavy particles /cm^ . Following bombardment,

the sodium profile in the film was determined using planar etch and radio tracer

counting techniques, i.e., counting the

etching solutions for the ^^Na removed in a given etching step. The profile results

are given in Fig. 12. In the case of ion

84
bombardment we observed that sodium is left

distributed throughout the oxide and is

piled up very near to the interface region,

certainly within the last 500 A of the

interface region. In the ion bombardment

case, approximately 10% of all the sodium

that was initially on the surface was moved

through the oxide film to the interface.

Under the conditions of neutral particle

bombardment, we cannot detect any sodium in

the Si02~silicon interface region. After

removing a few hundred angstroms of the Si02

film, there is very little or no sodium observed above the detection limit for the

radio tracer technique which is approximately 7x1014 Na atoms /cm^. Therefore in neutral bombardment there is no detectable

sodium at the interface, and this is substantially less, by at least 3 orders of

magnitude, than in the case of the ion bombardment. The case of using nitrogen ions

with the surface flooded with thermal electrons is an intermediate case. The sodium

at the interface is down by almost two order

of magnitude compared to the case of positive ion bombardment, but this is substantially above that which was observed in the case of the neutral particle bombardment.

This work was done using 2 keV ions or neutral molecules of nitrogen and the tar- gets were at room temperature. Other work,

using argon as a bombarding projectile, gave

essentially the same results. This work

therefore suggests that using neutral particles for bombardment may substantially re- duce problems associated with ion induced

mobile impurity transport in Si02 films.

However, a word of caution is in order. In our work a particle dose of ^lO^-Vcm^ was

used. However, much greater doses,

10 17-10 1 ^/cm2

, are required to profile a film of a few thousand angstroms, and it is

dangerous to extrapolate our results to that dose range.

In conclusion, I want to emphasize that the

study of optical radiation produced in particle-solid collisions provides fundamental

information on particle-solid interactions

and serves as the basis for a technique to identify constituents and contaminants on surfaces. This radiation is produced in

the sputtering process and the optical information is available for a variety of

diagnostic purposes in any type of

sputtering experiment.
TABLE 1

SOURCES OF OPTICAL RADIATION IN PARTICLE - SOLID COLLISIONS

1. Radiation from excited states of sputtered surface constituents.

2. Radiation from excited states of backscattered beam particles.

3. Radiation resulting from the excitation of electrons in the solid.

TABLE 2

SCANIIR DETECTION SENSITIVITIES IN Si0 2

Element Sensitivity

Al 7 X 10" 5

Fe 8 X 10" 4

Ca 2 X 10" 5

Mg 1 X TO" 4

K 8 X 10" 4

Na 3 X 10" 5

Rb 7 X 10" 5

Li 1 X 10" 5

Sr 7 X 10 -5

Ba 8 X 10" 5

Mn 1 X lO -4

Ti 1 X 10" 3

Be 4 X 10" 5

II
INCIC ION BEAI

Figure 1. SCANIIR Surface Analysis Apparatus.

fcjgsa— r~ i- t r-~~i - i ; i , , , i i i i i i i i t*" v 'i^-1 7800 7400 7000 6600 6200 5800 5400 5000 4600 42 00 3800 34 00 3000 WAVELENGTH (A)

Figure 2. Spectra of radiation produced by the impact of Ar+ (4 keV) on copper (top) and

nickel (bottom). Lines arising from neutral Cu, Ni and various contaminants are identified

by element and wavelength.

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ARGON ON Si0

2 (I200A ) / Si (~)

SII - 2882 JS

80 120 BOMBARDMENT TIME (MINS)

Figure 5. Profile of Si02 (1200 A) on Si (<=°) . Intensity of SiI-2882 A

line is plotted on a logarithmic

scale as a function of bombardment

time. A current density of ^ 10

pA/ cm2 was used for these measurements .

Figure 6. Energy level diagram for

excited Si atom (4s 1 P° state) near

silicon and Si0 2 surface.

A - RESONANCE IONIZATION

B,B' — AUGER DE-EXCITATION

DETECTION OF Cr ON SILICON

10 5 I

BOMBARDMENT TIME (SECS)

Figure 9. Detection of Cr on silicon.

Measurements were made as a function of

bombardment time using Ar+

(8 keV) , 4 x 10~6 amps/ cm2

)

.

PROFILE OF SiO ? (2000A)/At 0 (500A)/SiO (IOOOA)/Si(-)

0 A * rf^o

o - Sil - 2882 A a - All - 396I

A

80 ^

I 320 _J

400

I BOMBARDMENT TIME IMINS)

Figure 10. Profile of Si02 (2000 A)/A12 0 3 (500 A)/Si02 /(1000 A) /Si (°°) . Measurements

were made using Ar+ (5 keV)

.

11
10' 10 20 30 40 50 60 BOMBARDMENT TIME(minutes)

70

10'

Figure 11. Profile of Al+ (25 keV a 10 15 ions/cm2

) implanted in SiC>2 (1500 A thick).

Aluminum and silicon optical lines were measured as a function of bombardment time.

Profiling was accomplished using Ar+ (8 keV, ^ 5 x 10-6 amps/cm2

) bombardment.

Figure 12. Sodium profiles produced by

bombardment of Si02 /Si. Initial sodium

surface concentration on the Si02 surface

was 10 11+ /cm2

. Each sample was impacted with

a dose of 10 15 heavy particles/cm2

. The

sodium profile following bombardment was determined by planar etch and radio tracer

counting techniques.

4000 3000 2000 \000 DISTANCE FROM Si INTERFACE (A)

11
DISCUSSION OF THE PAPER

Participant : I wonder if you could comment

on the detectability of hydrogen?

White: Hydrogen is definitely detectable;

no question about it. In fact we have difficulty in getting rid of it. Our vacuums

are typically of the order of 10~° Torr, and

under those conditions, one has difficulty

getting rid of the hydrogen.

Participant : How does the photon production

efficiency depend on the primary beam kinetic

energy.

White: In the case of metal targets it is a very rapidly decreasing function of energy

as you go toward lower energies. From a metal target, in fact, we have been able to see radiation down to about 300 or 400 eV primary energy using an argon projectile.

In the case of the insulators, we do not know where the threshold is. It is below

50 eV for actually producing this radiation.

It appears to rise linearly with energy and

then apparently flattens out somewhere between one kilovolt and 1.5 kilovolts. In

the case of the metals it continues to rise

and does not peak until you get up into the

range of about 10 to 12 keV.

Participant : How do you measure the particle

beam current for a neutral beam?

White : You can do that in one of two ways.

One of the detectors that we use is a bolometer, where basically you measure the

temperature rise. This is calibrated against

an ion beam of known current. But you can also measure particle current by using the

photons that come from a metal surface be- cause when a low energy ion impinges on a metal, the ion is actually neutralized several angstroms in front of the surface before the close-in sputtering encounter takes

place. Therefore it does not matter whether

you start with an ion or a neutral at infinity. By the time the sputtering encounter

takes place, the incident projectile is, for

the most part, neutralized. On the other

hand photons are produced as a result of the

sputtering encounter. And therefore the

photon production efficiency from a metal

target is essentially independent of the

charge state of the incident projectile.

And, thus, you can use a metal target and

calibrate it in terms of its photon production efficiency using an ion beam of known

energy and known current. Then by charge

exchanging the ion beam and deflecting the

ions out of the way to let the fast neutrals

come into the target chamber you can use the

photon count rate that the neutral beam produces on that same metal surface as a means

of determining the neutral particle flux.

Both of these techniques work very well.

Lin : As we understand, the sputter ion yield

varies very much for elements in different

matrices. How about the photon yields of

elements in different matrices? This is very

vital to the capability of the depth profile

technique you described before.

White : The answer is that there are definite

substantial variations in the photon yields

as a function of different matrices. Impurities in a metal, for example, sodium impurities in metals, appear to have a different

photon production efficiency than they do in

a good insulator. How severe these problems

really are at this point I honestly do not

know. I do not know whether you see the

gigantic variations that you see in secondary

ion yields or not, but you do see variations.

Lin : Do you, then, expect it to be proportional to these sputter ion yields?

White : Not in all cases, no, definitely not. Other work that we are doing now shows that

there are definite differences between the

photon yields and the secondary ion yields,

at least in some cases, and we are trying to understand that right now.

Feigel : Is the instrumentation as simple and

as straightforward as you have shown or do

you have your magic or money buried in there

somewhere?

White : I can give a yes or no type answer to that. The apparatus that I am using has a lot of money invested in it. The reason is

that it was designed to be a general purpose

accelerator. It was built to study gas phase

atomic collision processes and particle-solid

collision processes. However, this radiation

can be produced for as little money as you

need to generate an ion beam, and then how

you use the information is obviously up to you. You can go to some very complex photon

detection schemes if you want to try to cover

all possibilities. If you are interested in

looking specifically for a given element, you

can isolate optical lines from all of the

elements of interest by using narrow band

interference filters. These can be purchased

for something like $200, a photomultiplier

cost you $500, and then you have either got

to count photons or you have to dc record. For specific applications, if you are interested in getting to an interface region, you

should be able to do that unbelievably cheap.

You may even be able to do it by eye, or using

a solar cell type detector.
Harrington : Both ion and neutral beams will

give secondary electrons and charged surfaces. Evidently from your experience, this

is a second order effect in charging surfaces. Would you care to comment on that?

White : You can, under the conditions of

neutral bombardment drive a lot of the

sodium through an SiC>2 film to an interface

if you put a collector in front of the target and collect the negative charge that

comes off the bombarded surface. By biasing

the collector to a certain voltage, you can control, or at least roughly control, what

the surface potential will be on the bombarded surface. Now, our experience has

been that if the collector is biased to approximately 100 V, you can then begin to see significant effects associated with

driving sodium through the oxide. However,

if you are not purposely collecting the

secondary electrons that are emitted from

the surface, then I do not expect that the

bombarded surface is going to increase to a potential which is any greater than the

average energy of the secondary electrons.

This, I think is of the order of just a few

volts, approximately 4 or 5 volts. Therefore, I think that this would be the only

driving field that you would have, at least

under the conditions that we have seen. This does not appear to be enough to give us sodium transport.

Participant : Do you have any techniques

for reducing the cratering effect of neutral

beams?

White : No. The only techniq ue that we have

is to narrow down the field of view of the

optical detection system to the point where

we are looking at a very small region. And

then we hope that we are in the middle of

the bombarded region. However, you can play

games. We produce neutral beams by accelerating positive ions up to the desired

energy and then partially neutralizing the

ions by resonant charge transfer in their

own gas. Resonant charge transfer is an atomic process in which there is very little

momentum transfer so the resulting neutral

has essentially the same energy and direction

as the initial fast ion. If the apertures

in the charge exchange region are opened up, then a certain amount of rastering of the

ion beam before it goes through the charge

exchange region is possible. By this technique a certain amount of rastering of the

neutral beam is possible. But we have not

done that yet and we have no plans to do it.

Participant : Have you looked for, or have

you seen, any effects of damage while profiling implanted impurities?

White : No, as a matter of fact we have not. We have just gotten into the implantation

business and we definitely have not looked

for topography changes in the bombarded region

yet. That is obviously something we should be

doing, and in the future we will be. But in terms of the ion implantation profiles, we have done so few of them at this point that

it is really meaningless to talk about them.
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Caithersburg , Maryland, April 23-24, 1975 (Issued March 1976)

NUCLEAR RESONANCE AND BACKSCATTERING SURFACE ANALYSIS OF SILICON AND RELATED INSULATORS*

Kenneth L. Dunning

Naval Research Laboratory

Washington, D.C. 20375

INTRODUCTION

Methods for surface analysis which are based

on nuclear reactions and scattering generally require more complex and expensive apparatus than do other methods considered at this workshop and, in addition, are more

limited with respect to the number of hostimpurity combinations that can be profitably

examined. On the other hand, these methods

have some distinctive features which make

them very useful in investigations pertaining to the technology of silicon and related

insulators. For example, distributions of sodium in thin layers of silicon dioxide on silicon substrates can be accurately and re- peatedly profiled; such measurements cannot

be made by means of methods which sputter

away the specimen in the depth-profiling

process

.

Two of the most useful nuclear methods for quantitative surface analysis are: nuclear

resonance profiling and Rutherford backscattering. This paper is concerned with the principles and applications of these methods,

with emphasis on the first since it is less

familiar, and with the discussion of the

second confined mainly to a recently-introduced improvement in the depth resolution

very near the surface. These methods can be

used to measure impurity depth-concentration

profiles with depth resolution of the order

of tens of angstroms (very near the surface)

.

In interior regions, depuh resolution is poorer, changing roughly as the square root

of the depth. The depth scale is based upon

the average rate of energy loss of the bombarding particles as they penetrate the

specimen. This basis for the establishment

of depth scales is better understood and

provides greater reproducibility than do methods based upon sputtering rates. Although charged particles with energies in the neighborhood of one million electron

volts do alter the crystalline structure of solids which they penetrate, their effect

* This research was partially supported by

the Defense Nuclear Agency under subtask

Z99QAXTA007, work unit 52, work unit title

"Electronic Transport and Trapping in Thin-Film Oxides."

upon the distribution of elements is generally

negligible and can be investigated by repeatedly measuring a particular depth-concentration profile.

References [1] through [3] are concerned with

materials analysis by means of nuclear reactions in general; references [4] through [8] are specialized to nuclear resonance reactions; reference [9] is a recent review of backscattering and nuclear-reaction analysis

methods

.

INTERACTION OF ONE-MeV PROTONS WITH SILICON

The sketch of Figure 1 is intended as an aid

to a brief review of the way charged particles

with energies in the neighborhood of 1 MeV

interact with silicon. For def initeness

,

protons are taken to be the incident particles; remarks on the interaction of alpha

particles (which are commonly used for Rutherford backscattering analyses) will be inserted

as the discussion proceeds. The sketch rep- resents a layer of silicon 1,000 A thick. As

the protons penetrate the layer, they lose energy in a large number of small-angle inelastic scatterings depicted at the top of the figure. The average loss of 1-MeV protons traversing 1,000 A of silicon will be 4 keV; the energy loss is linear with the depth

of penetration and, hence, the depth can be

conveniently measured in terms of this energy

loss. Energy loss serves as a basis for a depth scale in both nuclear resonance pro- filing and Rutherford backscattering. As the protons pass through the silicon layer, they

excite and ionize the atoms of silicon,

losing a few electron volts in each interaction, and emerge in essentially the same direction in which they entered. At these

velocities , a proton may occasionally pick up an electron but more than 99% of the time it

is simply a stripped hydrogen nucleus. If the protons were to enter as hydrogen atoms,

they would be stripped almost immediately and

proceed as stripped nuclei. Alpha particles

penetrating matter behave in a very similar

way; however, the rate of energy loss with

depth for 1-MeV a particles is about 7 times

as great as for 1-MeV protons. This is an advantage in measuring depth by means of energy loss.

35
Although the average loss for 1-MeV protons

traversing the layer is 4 keV, the losses of

individual particles will be distributed

about this average in a predictable way. For penetration depths of the order of tens

of keV, these loss distributions are Gaussian

shaped; for small penetrations, the loss

distributions are asymmetric about the mean

but can still be predicted with great accuracy by means, for example, of the theory of

P. V. Vavilov [11]. The full width at half

maximum of the energy loss distribution of

the protons emerging from the layer of silicon depicted is about 2.8 keV. It is this

phenomenon of energy-loss fluctuations which

is the chief source of depth resolution

degradation in nuclear resonance profiling

and Rutherford backscattering. The process

represented at the middle of the sketch of

Figure 1 is that of large-angle scattering.

If this large-angle scattering occurs because of the Coulomb field interactions and

not because of nuclear interactions, we call

it Rutherford scattering. The probability

for the occurrence of this process in a layer of the type illustrated is much less

than the first process described but sufficient so that detectors can be located to measure the number and the energies of the

scattered particles and this information

used to establish quantitative depth-concentration profiles.

A third process of interest is illustrated

at the bottom of the sketch; a proton is

captured into the nucleus of one of the target atoms to form a compound nucleus of a new atomic species. Many different modes of

decay of the compound nucleus are possible;

two are illustrated. The decay mode in

which a gamma ray is emitted is very useful

for nuclear resonance profiling but that in

which a nucleon is emitted can also be used.

The probability for nuclear capture is often

an extremely sharp function of the energy of

the incident protons and the process is sometimes described as resonance capture.

PRINCIPLES AND APPLICATIONS OF NUCLEAR

RESONANCE PROFILING

It is convenient to present the principles

of nuclear resonance profiling with an ex- ample which makes use of a sharp resonance

in the 27Al(p, y) 28 Si reaction at a proton

energy of 991.9 keV. The sketch of figure 2 shows a proton beam incident upon a silicon

target which has a layer of aluminum atoms

at the surface and another at a depth of

500 A. A depth scale in angstroms is shown

at the top of the target and at the bottom

of the target is a depth scale in keV of the

energy loss for the proton beam as it

penetrates the target. The resonance in this

reaction is very sharp having a full width at half maximum of 0.10 keV. If the proton en- ergy is advanced in small steps from 1/2 keV

below resonance energy to about 6 keV above resonance energy and a fraction of the gamma

rays from the target is counted by means of a sodium iodide detector (with an energy window

9 to 12 MeV) for a fixed number of protons on target at each energy step, a gamma-ray yield

curve of the kind illustrated in the lower

half of the sketch will be obtained.

To understand the role of the resonance, con- sider the interior target layer containing

aluminum atoms. When the bombarding energy

is 2 keV above resonance, the most probable

energy loss for the protons that have arrived

in this layer is 2 keV; thus, the protons in

this layer are very near resonance energy.

Most of the protons to the left of this layer

are above resonance energy and those to the

right are below resonance energy. The first

peak in the gamma-ray yield curve is associated with the first layer in the target con- taining aluminum atoms. The first peak is

higher and narrower than the second peak because there are more aluminum atoms in the

first layer than in the interior layer and

there is less energy-loss straggling. The area under the gamma-ray yield curve is proportional to the number of aluminum atoms in

the layer. The abscissa scale on the gammaray yield curve is a measure of the incident

proton energy in keV above resonance energy.

The figure has been drawn so that the abscissa

scales of keV energy loss and incident proton

energy above resonance are numerically equal.

The yield curve, when plotted as a function

of bombarding energy above resonance, provides an approximate depth-concentration profile. It is not a perfect depth-concentration profile because the incident proton beam

has a small but finite width to its energy

distribution and the resonance also has a small but finite energy width. The most

serious impediment to precision is the sta- tistical fluctuation in the energy loss of

the protons as they penetrate the solid, that

is, energy-loss straggling. This straggling

becomes more pronounced with depth, increasing approximately as the square root of the

depth of penetration.

Although the approximate profile provided by

the yield curve is often accurate enough for some purposes, a much more exact profile can

be obtained with the help of the following

equation which describes the yield Y(Eb) as

a function of the bombarding proton energy,

Eb [10].
r(Eb

) = Const^ j I

o(E) gCE^E^ w(E

1 ~E,A) C(A) dE dE

i dA

where o(E) is the cross section for the

27Al(p,y) 28 Si reaction in the vicinity of

the resonance (Breit-Wigner formula)

,

gCE-L.Ejj) is the beam-energy distribution,

w(E-l-E,~A~) is the energy-loss distribution

which gives the probability that for a depth

in the target for which the average energy

loss is ~K, the exact energy loss is E^-E,

and CCA) is the concentration function.

That is, the protons which are in an energy

bin of width dE^ at E-^ before entering the

target have the probability, w, of being degraded into energy bin of width dE at E in

the neighborhood of resonance energy on _ their way to the target layer of width dA at depth ~E. The beam-energy distribution depends upon the characteristics of the beamenergy analyzer and, for the case at hand,

the energy-loss distributions can be calculated by the theory of P. V. Vavilov [11,12].

One point of view that may be helpful is to think of performing the integration with

respect to E^ first. This would give a new

function of E and ~E. This proton energy

distribution in the target broadens with

depth, making the uncertainty for the calculated yield greater in interior regions than

at the surface. Note that an additional integration with respect to E^ will give the area under the yield curve, this area is

proportional to the total number of aluminum

atoms represented by the profile. The foregoing equation may be used to calculate, to within a constant, ordinates on the gammaray yield curve if the concentration function, C(K) is known. The constant can be

determined from a measurement of the yield

from a homogeneous target of known thickness.

If a yield curve is measured for a target

with an unknown concentration profile, we

can equate the experimental yield to the

right-hand side of the equation but are then

faced with the task of extracting the con- centration function. Since the measured

gamma-ray yield curve is an approximation of

C(A) , we may construct a trial function,

calculate the yield function, and compare it

with the experimental yield curve. If the

comparison is not satisfactory, the comparison may be used as a guide to construct a new trial concentration function and the

process iterated until satisfactory agreement between the calculated yield and experimental yield is obtained. With the aid of

a digital computer to make the calculations

and plot the experimental yield curve, the

calculated yield curve, and the concentration

function, a very satisfactory representation

of the concentration function can be obtained

with a few iterations.

Figure 3 illustrates results obtained with

the procedure described for a target which

was made by depositing 600 A of aluminum

oxide on a silicon substrate. A gamma-ray

yield curve was then measured by counting the

gamma rays as the bombarding energy was advanced in small steps from below resonance to

10 keV above resonance; these data are represented by the circles. The concentration

function of the aluminum in the oxide is represented by the solid line. The calculated

gamma-ray yield curve is represented by the

triangles. The fact that the yield rises before the profile (near resonance) is due to the finite widths of the resonance and the

incident beam-energy distribution. The

"spike" at resonance is due to the asymmetry

of the energy-loss distributions. The divergence of the calculated and experimental

yield curves from the profile near the back

face is due to the energy-loss straggling.

The area under the experimental curve is proportional to the number of aluminum atoms represented by the profile. The constant of

proportionality can be easily obtained and

experience has shown that the number of atoms

represented by the profile can be determined

with an accuracy of about 5%. This determination does not depend upon the shape of the

beam, the shape of the resonance, or the effects of energy-loss straggling.

Figure 4 shows a profile for an unknown sample

determined by this method. The sample was made by implanting 60-keV aluminum into silicon carbide and then annealing at a temperature of 1400°C. The sample was profiled before annealing and it was found the profile

had a Gaussian shape with a peak at about

700 A. The annealing drastically altered the

profile, bringing aluminum to the surface

region. A comparison of areas showed that

there is much less aluminum in the region of

the profile than there was before the annealing. The width of the spike in the profile

at the face of the target is about 20 A at half height. This gives some idea of what

the depth resolution can be at the very face

of the target. In interior regions, it is

poorer than this due to energy-loss straggling. If such a spike concentration were

located in the interior of the target then the

yield curve representing it would be smeared

out; the accuracy to which the profile could

be constructed depends upon the depth at

91
which the spike occurs and upon the statistical uncertainty of the experimental data.

It is feasible to profile sodium in thin

silicon dioxide layers by means of this

method. Reproducible profiles have been obtained for sodium in thin layers of silicon

dioxide on silicon, both for test structures

which have thin gold layers on the silicon

dioxide surface and for test structures

which have no metallic layer. There is,

however, a limit to the fluence of protons

which can be passed through the silicon dioxide layer without disturbing the sodium

distribution. Figure 5 shows the results of

measurements on a test structure consisting

of 3,000 A of silicon dioxide on a silicon ' o substrate with a 175-A layer of gold evaporated onto the surface of the silicon dioxide. Prior to the deposition of the gold

surface layer, the structure was implanted

with 60-keV sodium to a fluence of 10 15 atoms/cm2

. The profile represented by the

solid line was obtained for the as-prepared

structure. It was re-measured to insure

that the measurement process did not alter

the profile. The specimen was then irradiated with 13-keV electrons (with the gold at +30 V relative to the silicon) to a fluence

of 10 13 /cm2

, the profile measured, and then

re-measured to insure reproducibility. The

electron irradiation has caused a migration

of the sodium from the region profiled. It appears that some sodium has moved deeper

into the specimen and that sodium has migrated through the gold film on the surface

of the silicon dioxide. However, there is

the possibility, of sodium contamination to the gold surface during the electron irradiation. These data are presented simply to demonstrate the profiling technique.

PRINCIPLES AND APPLICATIONS OF RUTHERFORD

BACKSCATTERING SURFACE ANALYSIS

Much has been written concerning the application of Rutherford backscattering to the

analysis of surf aces [2 , 9 ] .Nuclear resonances

are not used in this method to improve the

depth resolution as was done in the previously discussed method. Instead, the depth

scale is established on the basis of the

energy loss of the incident charged particle

(usually an a particle) in its trip to the

scattering atom in the solid being analyzed

and, after scattering, to the detector. In some situations, scattering can be enhanced

by nuclear resonances but this is generally

avoided so that the cross section for scat- tering can be computed from the Rutherford

formula. The method is best suited to the

profiling of heavy impurity atoms in light

hosts

The principles upon which the method is based

can be quickly reviewed by referring to fig- ure 6. Mono-energetic a particles (typically

in the range 1 to 3 MeV) are shown incident

from the right upon a target consisting of an aluminum oxide substrate covered by a thin

nickel film which is, in turn, covered with a thin gold film. Scattered particles are usu- ally detected by a semiconductor detector

shown at the upper right. The voltage pulses

from such a detector are proportional to the

energy of the scattered particle. If these

pulses are sorted by a pulse-height analyzer

and displayed on an oscilloscope, a pulseheight (energy) spectrum such as that displayed in the lower part of the diagram is

obtained. The energy, EQ , of the incident

particles is an upper limit to the energy of

the scattered particles. The high-energy

side of the peak in the scattered spectrum

marked Au is due to the scattering from the

surface of the gold since the gold atoms are

the heaviest atoms in the target and the incident particles have lost very little or no energy in the target before being scattered

(billiard-ball kinematics) . Alpha particles

that are scattered from the interior of the

gold film lose energy on the way to the scattering center and again on the way to the

detector by small-angle inelastic scattering

such as that discussed in connection with

figure 1. Thus, if there are no atoms in the

target which have nearly the same mass as the

gold atoms, a peak will be formed in the

scattered spectrum whose height is proportional to the number of gold atoms/cm2 and

whose width is proportional to the thickness

of the gold layer. A similar peak appears in

the scattering spectrum due to the presence

of the nickel. The a particles scattered

from nickel all have lower energies than any

particle scattered from gold because the

nickel atoms are not as heavy as the gold

atoms and, in addition, the incident particle

must lose energy in the gold layer before any

scattering from the nickel can take place.

Since the aluminum oxide substrate is con- sidered to be very thick relative to the

nickel and gold layers, the a particles scat- tered from aluminum atoms form a continuum

which extends essentially to zero energy and,

similarly, the contribution to the backscattered spectrum from the oxygen appears on top

of this continuum. The signal-to-background

ratio is much better for the heavy elements

than for the oxygen. Since the cross section

for scattering can be calculated from the

Rutherford formula, and the rate of energy

loss of the a particles in various materials

has been measured and cataloged [13] , the dat

in the scattered spectrum may be used to con- struct a quantitative depth-concentration

profile
The depth resolution for the method just

described is determined by the finite energy

width of the incident a-particle energy distribution, the energy-loss straggling in the

target, and the energy-loss straggling introduced by the detector. The energy loss

in such detectors is usually in the range 10

to 15 keV with the width of the energy-loss

distribution being in the neighborhood of 10

keV; this can cause an uncertainty in the

depth scale at the face of the target, typically, in the neighborhood of 100 A.

There are a few laboratories where the semiconductor detector in a scattering apparatus

has been replaced by a magnetic spectrometer

which determines the energy of the scattered

particle without introducing energy loss.

Figure 7 shows such an arrangement which has

been developed and tested by Dr. J. K. Hirvonen and his colleagues at the Naval Re- search Laboratory. The spectrometer can be

rotated about a verticle axis and positioned

to accept a certain fraction of the particles

scattered from a target. Of the particles

scattered, those within a small range of

momentum (determined by the strength of the

magnetic field which is perpendicular to the

paper) will pass through the spectrometer

and be detected in a position-sensitive de- tector in the image plane of the spectrometer. The voltage pulses from this detector

are routed to a pulse-height analysis system. The energy coordinate of the contribution of a particle to the scattered-particle

spectrum is determined by the position at which the particle strikes the detector

which is located in the image plane of the

spectrometer. With this arrangement, one obtains only a small part of the semiconductor

energy spectrum but the degradation of depth

resolution introduced is an order-of-magnitude smaller than that introduced by a semiconductor detector.

In order to measure the effect of the scat- tered-particle detector on the depth resolution, it is useful to have a very thin target of heavy atoms. The data of figure 8 were obtained with a target made by depositing gold onto silicon dioxide to an average

thickness of one monolayer. However, gold

deposited in such thin layers tends to form

"islands" which means that some of the incident 2-MeV a particles were scattered by

several monolayers of gold. At any rate, we

have a target which is quite thin.

The upper part of figure 8 shows a spectrum

obtained with a semiconductor detector. The

left half of the spectrum is due to the silicon and oxygen and is not of interest. At

the high-energy end of the spectrum is the

contribution from the gold, a peak with a full

width at half maximum of 16 keV. This width

is due almost entirely to the energy-loss

straggling in the semiconductor detector.

The spectrum at the lower right of the figure was obtained with the magnetic spectrometer.

The full width at half maximum of the peak is now 5 keV and less than half of this is due

to the uncertainty in the scattered energy

due to the spectrometer arrangement which is

about 2 keV and corresponds to an uncertainty

in the depth scale of a depth-concentration

profile of about 25 A. Thus, the depth reso- lution that may now be achieved with Rutherford backscattering at the surface of a specimen is comparable to that illustrated in

figure 4 with nuclear resonance profiling.

The magnetic spectrometer arrangement has

another advantage; it passes scattered particles whose energies fall only within a narrow

band, whereas the semiconductor detector sees

a much larger number of particles per unit of

time. This results in a higher background in

the semiconductor detector scattered spectrum

due to pulse "pileup" and degrades the detection level. The "spectrum" at the lower left

of figure 8 was obtained by bombarding a silicon dioxide target with no gold layer. Although five times as many incident a particles were used in this case as were used to obtain the spectrum at the lower right, not a single event was recorded in the 40-keV re- gion spanned by the abscissa. If one count had been recorded, the detection limit for

the gold would be 10 10 atoms /cm2. The detection limit with the semiconductor detector

is about 10 12 gold atoms/cm2

. Thus we see

that we can expect a depth resolution very

near the surface of about 25 A with a detection limit of less than 10 10 atoms/cm2 with

this spectrometer arrangement for a target

that is free from contaminating atoms with

atomic weights near that of gold.

CONCLUSION

Some of the more important features of the

two discussed methods of obtaining depth-concentration profiles can be summarized and

compared as follows.

(1) Both methods discussed are nondestructive

in the sense that the atoms struck are not moved very far (at least almost never) . However, the probing ions do move the struck

atoms from their lattice sites far enough to degrade the crystalline structure.

(2) Repeated measurements can be made on the same test specimen and these measurements can be interspersed with various other processes

such as annealing and irradiation.
(3) Depth scales do not depend upon surfaceremoval techniques and can be very accurate.

(4) The total number of atoms represented by

a profile can be determined, often to accuracies better than 5%.

(5) The apparatus for these methods is com- plex and, hence, these methods are suitable

only for some problems related to the silicon-device development and production. They

can be very useful in support of the development of measurement technology which, in

turn, can be more directly and easily applied

to device problems.

(6) The number of host-impurity combinations

that can be analyzed is limited and detection limits vary from 10 10 at oms/cm2 upwards.

The nuclear resonance profiling method is

isotope specific and useful for light impurities in any kind of substrate that does not

have interfering resonances. The Rutherford

backscattering method is best for impurities

that are heavier than the host impurity.

ACKNOWLEDGEMENTS

Support of this work by the Defense Nuclear

Agency is gratefully acknowledged. The part

of this report dealing with Rutherford backscattering is based on the work of Dr. James

K. Hirvonen, Naval Research Laboratory.

Contributions of Mr. Harold Hughes, Naval

Research Laboratory, to the work are gratefully acknowledged. The experimental work

with accelerated ions was carried out with

the 5-MV Van de Graaff accelerator of the

Ion Beam Applications Branch, Radiation

Technology Division, Naval Research Laboratory, with assistance from the staff. Dr.

E. A. Wolicki provided helpful comments and

suggestions

REFERENCES

1. Amsel, G. , Nadai, J., D'Artemare, E.

David, D. , Girard, E.,and Noulin, J.,

Nucl. Instr. Meth. 9^2, 481 (1971).

2. Wolicki, E., Nuclear and Ion Beam Techniques for Surface and Near-Surface

Analysis, NRL Report #7477, (December

13, 1972).

3. Wolicki, E., Surface Analysis Using

Nuclear Reactions in New Uses of Low

Energy Accelerators, (Plenum Publication

Corporation, New York, to be published).

4. Moller, E.,and Starfelt, N. Nucl. Instr. Meth. 50, 225 (1967).

5. Witton, J., and Mitchell, L. , Can. Jour.

Phys. 49, 1125 (1971).

6. Bernett, M. K. , Butler, J. W. , Wolicki,

E. A., and Zisman, W. A. , J. Appl. Phys.

42, 5826 (1971).

7. Dunning, K.,and Hughes, H. , IEEE Trans.

Nucl. Sci. NS-19 , 243 (1972).

8. Dunning, K. , Hubler, G. , Comas, J.,

Lucke, W.,and Hughes, H. , Thin Solid

Films 19, 145 (1973).

9. Chu, W. , Mayer, J., Nicollet, M.

Buck, T. , Amsel, G., and Eisen, F. , Thin Solid Films 17_, 1 (1973).

10. Gove, H. , Nuclear Reactions, p. 293,

Endt and Demeur, Eds. (North Holland

Publishing Co., Amsterdam, The

Netherlands 1959)

11. Vavilov, P., Zh. Eksp. Teor. Fiz. 32,

320 (1957); Transl. JETP 5, 749 (1957).

12. Seltzer, S., and Berger, J., Nuclear

Sciences Report 39, p. 187, National

Academy of Sciences - National Research

Council, Washington, D. C. (1964).

13. Williamson, C. , Boujot, J., and

Picard, J., Tables of Range and

Stopping Power of Chemical Elements

for Charged Particles of Energy 0.5 to

500 MeV, Report CEA-R-3042 , Centre

d' Etudes Nucleaires de Saclay, France

(1966).

100
INTERACTION PROCESS PROBABILITY

ONE MeV

PROTONS

INCIDENT

SILICON

"°^NUCL

SMALL-ANGLE

INELASTIC SCATTERING

NUCLEAR CAPTURE

-0.9999

LARGE-ANGLE ~I0" SCATTERING

HOT*

(SINGLE RESONANCE)

Figure 1. Sketch illustrating tlje interaction of 1-MeV protons in a 1,000 A thick slab

of silicon.

(a)

Figure 2. Diagram illustrating the relationships between depth-concentration profiles

and yield curves.

(b)

DEPTH (ANGSTROMS)

500 1000 1500

I 11 PROTON BEAM

AL ATOMS

SILICON

DEPTH (KEV ENERGY LOSS)

0 2 4 PROTON ENERGY (KEV ABOVE RESONANCE ENERGY)

Figure 3. Experimental yield, calculated

yield, and "profile" for an A12 0 3 target

4.6-keV (600 A) thick.

2.0 4.0 6,0 8.0 RELATIVE PROTON ENERGY (KEV)

101
DEPTH (ANGSTROMS) 400 600 900

SAMPLE NO 31. (60KEV) 4i lO^AL^/CM 2 IMPLANTEO INTO SIC ANNEALEO

-RELATIVE CONCENTRATION

EXPERIMENTAL YIELO CALCULATED YIELD

RELATIVE PROTON ENERGY 1KEVI

Figure 4. Profile, experimental yield, and

calculated yield for SiC implanted to a flu- ence of 4 * 10 14A1 atoms/cm2 at an energy of

60 keV and then annealed at 1400° C.

A IN SiO 2000 3000

I75ANolp.x) Mg

1417 keV RESONANCE

TARGET 5b UNSTRESSED

AFTER IRRADIATION WITH I0

l3/CM 2 13- keV ELECTRONS WITH AU AT

+ 30V RELATIVE TO SI

0.0 4.0 8.0 12.0 16.0 20.0 24.0 BOMBARDING ENERGY RELATIVE TO RESONANCE ENERGY (keV)

SEMICONDUCTOR

DETECTORFigure 5. Depth-concentration profiles for

sodium implanted at an energy of 60 keV to a fluence of 10 15 /cm2 into an Au/Si02 /Si test structure. The solid line shows the measured

yield curve for the as-prepared specimen; the

dashed curve shows the yield curve after the

specimen has been irradiated with 13 keV

electrons to a fluence of 10 13 /cm2 with the

gold layer at +30 V relative to the Si during

the irradiation.

ENERGY OF BACKSCATTERED PARTICLE

Figure 6. Diagram illustrating the Rutherford

backscattering method for obtaining depthconcentration profiles.

102
Figure 7. Diagram illustrating the use of a charged-particle magnetic spectrometer in

the detection leg of a Rutherford backscattering apparatus

.

AU MONOLAYER ON SILICON DIOXIDE

Figure 8. Rutherford backscattering spectra

illustrating the superiority of a magnetic

spectrometer over a semiconductor detector

in the detection leg of a Rutherford backscattering apparatus. The spectrum at the

top was obtained from a very thin layer of

gold on silicon dioxide with a semiconductor

detector; the gold peak at the right hand

side of the spectrum has a full width at half maximum of 16-keV. The lower-right

spectrum was taken with a magnetic spectrometer; only the gold peak is shown and it has

a full width at half maximum of 5 keV, less

than half of which is introduced by the detection system. The lower left "spectrum"

is for a silicon dioxide target with no gold

layer and the magnetic spectrometer and

shows that with five times as many incident

a particles as for the lower-right spectrum,

no background counts were registered.

-I 1 l 1

Solid State Detector -2000 V Energy Spectrum

10 fj.C charge

-1000 I Detector

^""'•*'""*<*i Resolution —

\ 16 KeV

_j iA 1 1

i AE

MAGNETIC

SCATTERING YIELD (COUNTS)

1

1 BACKGROUND A J Si Wafer

-2 50^C ,

icvc I

i n

-150 J system / resolution _J

j— 5 KeV

103

it
DISCUSSION OF THE PAPER at MeV energies as being almost completely

stripped.

Harrington : The magnetic spectrometer that

you use depends, of course, on ions coming

from the surface. From what depths can these

reflected particles originate and still be

ions? Are not most of the particles from,

say, the silicon substrate coming back as neutrals?

Dunning : At this energy many particles come

off as charged particles and so I would say

one can analyze at least to micron depths.

Harrington : I was under the impression that

most of the particles were coming out as neutrals from those depths. Are they still

ionized?

Dunning : My impression is that , at MeV

energies, as soon as an ion enters a solid it

is almost entirely stripped. It will pick up

a little charge and loose a little charge as

it goes through but we always think of ions

Dobrott : You showed us a profile of a sodium

ion implanted oxide. You told us that the

sodium flux was lO^-Vcm^ but do you happen to remember at what energy the implant was made?

Dunning : The energy for the sodium implant I think was 60 keV.

Dobrott : Then the actual peak that you did

see was probably a little bit deep for the

theoretical projected range of the sodium in

the Si02-

Dunning : It is not exactly according to theory but I think it is not far off.

Dob rott : The peak was shown to fall between

two thousand and twenty-five hundred Angstroms

into the sample, and that is pretty deep.

Dunning : No, I would say that that was

fairly near the theoretical peak.

104
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

APPLICATIONS OF SCANNING AUGER SPECTROSCOPY (SAM) TO THE SILICON INTEGRATED CIRCUIT (SIC) TECHNOLOGY

Joseph M. Morabito

Bell Laboratories

Allentown, Pennsylvania 18103

I. INTRODUCTION

Auger electron spectroscopy is a technique

based on the emission and energy analysis of

secondary electrons produced by a high energy

(3-10 keV) electron beam. The energy of a small fraction of the secondary electrons can

be related to the core levels of a target

atom; these electrons, the Auger electrons,

have an escape depth in the range of approximately 5 to 20 A, making Auger electron spectroscopy a truly surface sensitive technique

which can detect all elements with exception

of H and He down to approximately 10 12 atoms

per centimeter squared or .1 atom percent.

The technique, in principle, is nondestructive; but electron beam damage and electron

beam desorption effects can occur. In com- bination with insitu ion sputtering, indepth profile analysis can be obtained and

by the use of multiplexing, several elements

can be monitored while sputtering through the

film. By the use of finely focussed electron

beams, or optical methods, selected area surface analysis can also be obtained, which is very important when analyzing actual integrated circuits. Quantitative analysis is also possible along with Auger imaging, which

provides a surface distribution map of the

impurities on the surface.

II. INSTRUMENTATION

Figure 1 shows some of the basic components

and associated electronics of the commercially available PHI scanning Auger microprobe. C

1 ) This consists of an electron gun used to excite the Auger electrons and a spectrometer for energy analysis. Electrons

of a well-defined energy are focussed on the exit aperture of the spectrometer and enter

an electron multiplier where they are detected by phase sensitive detection using a lock- in amplifier. The electron gun consists

of 4 basic components; a tungsten triode

source, condenser and objective lenses, and

deflection plates. Its operation is somewhat

similar to a scanning electron microscope.

The condenser lens and the objective lens are used to demagnify the crossover formed by the

triode source. The deflection plates are used to raster the beam over the sample.

Electron beam spot sizes down to approximately 5 microns can be obtained and the beam

currents possible are in the range of 10 nanoamps to 10 microamps. To locate areas of interest on the sample, absorbed current micrographs or secondary electron micrographs are

used. Useful magnification is about 1,000 X.

Line scans and Auger imaging provide information on the variations of surface impurities.

The Auger imaging is obtained by rastering

the primary electron beam and intensity modulating the z axis of a CRT. Table I shows the

minimum time constant required for the analysis

of 10 and 1 atom percent, corresponding to signal-to-noise ratios of 10 to 1 and 100 to 1, as

a function of varying beam size (in urn 2

) for current densities of 10 7 amps per micron

squared and 10 & amps per micron squared. A practical limit on the lock-in amplifier is approximately 10 seconds . This is due to am- plifier drift. This limit is reached at 1 atom percent for a beam area of approximately

10_microns squared for a current density of

10 7 A/um2

. If_the current density is increased from 10 7 to 10 6 A/um2

, this limit

of 1 atom percent is reached at a beam area

of approximately 1 micron squared. Beam areas

in the range of 20 to approximately 1 um

2 are very useful in the analysis of actual integrated circuits. As the beam size is decreased, the time that it takes to do the analysis increases, and this time is inversely

proportional to the_beam current. Current densities in the 10 7 A/um2 range can be obtained, as shown by the work of Brandis^ 2 ), with the tungsten triode source down to ap- proximately 1 micron squared. Current densities in the range to 10 6 can only be obtained

with a lanthium hexaboride source, which is a brighter source. However, at these beam current densities there is a possibility of beam

desorption and beam damage effects,

particularly in the case of insulators.

III. MODES OF OPERATION

The Auger technique, like SIMS, has basically

three modes of operation: a) Auger spectra,

b) Auger images and c) in-depth profile analysis. These three modes of operation (illustrated in Figure 2) provide a truly threedimensional microchemical analysis. Auger

spectra are used to detect and identify surface

impurities and impurities in the bulk of the

sample. Spectral interferences are less of a 05
problem in Auger Spectroscopy than in SIMS,

but they can occur since the inherent energy

spread of Auger electrons is such that it is

difficult to resolve peaks that are within 1 or 2 eV of each other. When this becomes a problem, the best solution is to add an X-ray

source and to obtain photoelectron or ESCA

spectra. Some of the unique features of the

ESCA technique have been described in the

literature.

3

) One of its unique features

is that it can provide information on chemical environment, but it is also possible to occasionally obtain information on chemical

environment using the Auger spectra. > 5

» 6 ) This information is contained in peak energy

shifts, changes in peak shape, and in fine

structure on the low energy side of the peak.

Figure 3 shows Auger spectra which were obtained on a silicon nitride layer of a silicon diode array. The spectrum at top is the

spectrum obtained prior to sputtering. Prior

to sputtering the silicon is partially present in oxide form( 5

) and the low energy silicon transistion is occurring at approximately

78 eV. After sputtering we have removed the

oxide, the silicon is now no longer present

as oxide and its energy has shifted from 78 eV to approximately 90 eV which is a shift of

12 eV. In the case of pure silicon the low

energy peak occurs at 92 eV. Although such

changes in the Auger spectra are often rather

difficult to interpret, they do provide a fingerprint like the SIMS spectra of the

chemical environment

A lateral resolution of approximately 0.5

microns has been demonstrated

2 ) in Auger

imaging. Auger imaging does not require sample consumption, but is limited to approximately 1 atom percent concentration.

Auger profiling is particularly useful when

monitoring concentration from the surface

into the bulk of a thin film, or when the extent and amount of interdif fusion in multilayer film is of interest. Figure 4 shows an Auger profile of a titanium-copper-nickelgold (TCNA) thin film system'' 7 '' after 4 hours

of 350°C. The copper and nickel have com- pletely diffused through the gold layer and

are present on the surface as oxides. The

metal oxides can interfere, for example, with

TC bonding to this system. Copper is present throughout the nickel diffusion barrier

layer and has also diffused into the titanium

glue layer. The depth resolution possible

with Auger profiling is primarily limited by

artifacts inherent in the sputtering process.

In polycrystalline films the depth resolution

is in the range of 10 to 30 percent of the

thickness removed by sputtering. (8) Some or the problems inherent in sputtering are: First, the possibility of differential

11

sputtering; this is particularly true in the case of alloys.

(

9 ) It is not always possible

to eliminate, but it can be minimized by the

use, when possible, of higher energy (larger

escape depth Auger electrons) because these

transitions are less sensitive to the change

in surface composition

(

9 ) which results from

differential sputtering. Next, the recoil or "knock on" phenomenon^ 0 ); this is an effect

whereby lattice atoms are pushed deeper into

the solid during the sputtering process. It is a cumulative affect, i.e., a function of

film thickness. This can sometimes be minimized by the use of low energy primary ions

in the range of around 500 eV to about 2 KeV. There is also a possibility of the redeposition

of sputtered atoms. This can be avoided by

careful instrument alignment. Topographical

changes can also occur during the sputtering

operation. The best way to avoid these is to use smooth substrates when possible. And,

finally, probably one of the worst problems,

a varying sputtering rate throughout the film

due to changes in composition or changes in the

alloying that occurred during the interdiffusion process. This greatly complicates the

assignment of an accurate depth scale. Accurate depth scales are important when extracting

diffusion coefficients from Auger profiles . ^ 8

> *jf

:? However, the work of Dahlgren and McClanahan' 12 ' has shown, in the case of alloys where the so- lute is about 10 atom percent or greater, that

the lower sputtering yield element actually

controls the sputtering rate. By using this

fact, it is often possible to assign a rather

accurate depth scale from measurements of

sputtering rate (z) and time.( 8 )

IV. STATUS OF QUANTITATIVE ANALYSIS

There are essentially three methods of

performing quantitative Auger analysis. Method A is a method based on the use of standards

prepared by techniques such as ion implantation, ( 13 ) reactive sputtering, and con- trolled doping.

15 ) Method B is the use of

pure elemental samples and sensitivity factors

with an equation of the following form:

I./AC. =

l

rel m

I I . /A

(1)

j=l

rel

where C^ is the concentration of element i,

Ij is the measured Auger peak height, ^Te ± is

the relative Auger yield of element i referenced or normalized to some matrix element,

and j is a running index summed over all the

elements present in the sample. Method C is

a technique which is independent of any

standards - called the first order approximation
method. Generalized models, e.g., models

such as those used for quantitative analysis

using the electron microprobe (17) or SIMS( 18 ) technique, do not at the present time exist

for Auger electron spectroscopy. The major

reason for this is that we do not have an ac- curate data base on the ionization cross sections, the escape depths, or on the backscattering correction. Other reasons include: the sputtering artifacts mentioned

and the fact that it is difficult to perform

background correction procedures in Auger

spectroscopy. Figure 5 shows a calibration

curve

(

15 ) obtained on samples which were

prepared by bulk doping. With these calibration curves it is possible to measure

the concentration of phosphorus and boron

in silicon. This data also establishes

the detectability limits for these dopants in

silicon. In this case, the detectability

limits for boron and for phosphorus in silicon is approximately 8 x 10 18 atoms per cubic

centimeter. Method B has been successfully

applied to numerous multicomponent systems. ( 8 » 19 ) The first order approximation

method does not depend on standards, but is

based primarily on the fact that the Auger

yield is determined primarily by the ionization cross section ($) , the transmission of

the spectrometer (n) , and the Auger escape

depth (d) . Work to date has shown that the

ionization cross section ($) is proportional

to one over the square of the Auger energy,

at least for K and L transitions. For the

CMA analyzer, the transmission (n) is linear

with energy. Escape depths (d) in a range

of 100-1,000 eV appear to be proportional

to the square root of the energy. The product, $nd, is therefore proportional to one over the square root of the energy. We can use this fact to estimate the Auger yield or sensitivity factors. Table II shows some of

the results using this method for magnesium

oxide, copper oxide, cadmium sulfide, gallium

phosphide, and potassium chloride. The re- sults are quite good, particularly since no corrections

(

9 ) for sputtering artifacts have

been made.

Table III shows a comparison between the

bulk detectability limits of Auger and SIMS

for the common dopants in silicon and for

the light element in tantalum. In general,

the detectability limits for Auger are in

the range of 10 19 and 10 20 - which is much

less sensitive with regard to bulk detectability limits than SIMS, which can detect

boron down to lO 14 atoms/cm 3

. It is very

important to be able to detect concentrations

in the range of 10 17 -10 14 in the case of

transistor fabrication. This is just not possible at the present time with Auger

electron spectroscopy. Possible ways to

1

increase the sensitivity of Auger electron

spectroscopy are: 1) increase the transmission (n) of the spectrometer; 2) improve the

detection scheme, e.g., the use of multichannel analyzer, data smoothing techniques,

deconvolution, background correction, integration techniques, etc.; 3) signal processing - analog averaging or digital averaging;* 4) use

a different source to excite the Auger electron. Some possibilities include high energy

proton sources and ions. With these sources

the background could be lower.

V. APPLICATIONS OF SCANNING AUGER MICROPROBE

IN SILICON INTEGRATED CIRCUIT (SIC)

FABRICATION

a) Analyses Of Contaminated Gold Beam Leads

Gold beam leads are used to attach SIC's to gold metallized ceramic. Any contamination

of these leads or of the gold on the ceramic

can prevent intimate gold-to-gold contact,

which results in unreliable bonding. Figure

6 shows an Auger spectrum and absorbed current

micrograph of a contaminated beam lead. The

beam lead is ^ 4 mils in width. The bright

lower line on the micrograph indicates that

most of the beam current is absorbed by the

conducting beam lead which appears dark. The

major surface contaminants are carbon and

oxygen with trace amounts of copper, sodium,

and silicon. Although the high energy gold

transition is detectable, the low energy gold

transition at 69 eV is not. This low energy

gold transition has an escape depth of ^ 5 A

and is completely attenuated by the carbon

layer. The carbon layer does not, however,

completely attenuate the high energy gold

transition at 2,017 eV, which has an escape

O depth of ^ 20 A. This suggests thaf the

carbon layer is no greater than 20 A in

thickness. The carbon Auger image obtained

showed the carbon to be homogeneous across

the surface, but the copper and silicon Auger

images suggest a heterogeneous distribution

for these elements. Figure 7 shows the indepth profile of carbon. As the carbon film

is- removed by sputtering, the intensity of

the low energy gold transition increases.

The thickness of the carbon layer can be

calculated by the following expression^ 20 )

Au Au L \ Au/ \ c Au I

J

* This is not a very practical approach since

a factor 10 increase in sensitivity would

require a factor 100 increase in analysis

time

.
where I c /IAu ^s t'ie measured carbon peak

height normalized to the high energy (2,017

eV) gold peak, and I£/Ia.u is tne ratio or tne

272 eV carbon Auger peak height to the 2,017

eV peak height for carbon and gold in the

pure state. xc is the thickness of the car- bon layer, d c is the escape depth of the

carbon 272 eV Auger electrons and dAu is the

escape depth of the 2,017 eV gold Auger

electrons. Equation 2 can be simplified to

x = d. In

c Au

(3)

which has been plotted for dAu equal to 4, 6,

and 10 A in Figure 8. These curves can then

be used to estimate the thickness of the car- bon layer from the data shown in Figure 6. The thickness of the carbon layer for

I c /I^u

= 35.2 is in the range of 8 A (for

dAu = 4 A) to 20 A (for dAu = 10 A) . It is

doubtful that this amount of carbon could

seriously interfere with bonding. Therefore,

the carbon layer is most probably polymerized

by the high energy electron beam, leaving

behind a carbon layer of reduced thickness.

There is, in addition, the possibility of

beam desorption of some of the carbon present

on the surface during analysis. Therefore,

the actual thickness of the contamination

layer prior to electron beam excitation is

perhaps a factor 5 higher, i.e., in the range

of 40 to 100 A, which would definitely inhibit bonding.

b) Analysis of Contaminated Gold Plated

Ceramic

Figure 9 shows the Auger results obtained on

a gold plated ceramic after pre-bond cleaning. Pre-bond cleaning which includes a heat

treatment of 120°C for 1 hour has increased

the Cu (920 eV),Au (239 eV) and Ag (354 eV)/Au - (239 eV) ratios from an average of .36 to 1.24 for the "as-received" sample to .49 and

1.70 respectively. This increase is due to diffusion from the bulk via grain boundaries.

There is also a large increase in the oxygen

peak. The large increase in oxygen suggests

that these impurities are present as oxides.

There is, however, no characteristic change

in the Auger spectra of Figure 9 (e.g., peak

energy shifts, additional plasma loss peaks)

which can be used to verify that oxidation

has actually occurred, but Schon^ 21 ) has

recently used X-ray induced Auger emission

and photoelectron spectroscopy to identify

the presence of Cu, Cu20 and CuO. Unfortunately, the energy resolution required^ 21 )

is beyond the capability of most spectrometers

used in Auger electron spectroscopy.

The surface concentration of Ag and Cu varied

in the range of 1 to 4 at. %. This magnitude

is consistent with the fact that 50 ppm (at. %) in the bulk, if completely concentrated within

the first 10 A of the surface, would correspond to 4.95 at. % and to 9.9 at. %. The

above calculations were made for a gold bonding

pad with an area of 2.25 x 10-2 cm2 and a gold

thickness of 10-1+ cm or 1 ym. Figure 10 shows

the calculated surface concentration, Cg ca ^,

which would result if all of the bulk impurities should concentrate within the first 10 A

of the surface. The bulk concentration was varied from 5 to 300 ppm (at. %) . Although it

is highly improbable that all or even most of

the bulk impurities will concentrate on the

surface after processing prior to bonding, a fraction (a) of the impurities will, and this

fraction can be measured by quantitative Auger

analysis

The presence of Sn (Figure 9) is apparently

due to inadvertent contamination introduced

during pre-bond cleaning, but the Cu and Ag

are codeposited during plating. A regression analysis (shown in Figure 11) of the

bonding and Auger data showed the correlation

between the presence of Ag, Cu, 0, and Sn on the surface and poor bonding to be 0.74. The

impurities present in the bulk are therefore

segregating to the surface during thermal processing, oxidize, and act as barriers to intimate gold-to-gold contact.

c) SAM Analysis of Laser Debris On Ti-Pd-Au

Resonator Electrodes

The frequencies of Ti-Pd-Au resonator electrodes

can be raised by vaporizing portions of mass

from the resonator electrode.

22 ) Using the focussed beam from a Q-switched Nd:YaG laser (A = 1.06 ym) , a series of 25 ym diameter holes on approximately 125 ym centers are vaporized as shown in Figure 12. When nominal frequency has

been reached, laser trimming is terminated. The use of an on-line computer and optical beam

steering techniques permit raster scanning by

which each resonator is adjusted in three suc- cessive passes. In this manner the effects of

thermal shock can be minimized and the adjustment control is as precise as the vacuum operation (± 12 Hz). All initial filter electrical

requirements can be met; i.e., adequate control over resonator frequencies can be exercised.

However, since the laser machining involves the

melting and evaporation of material from the

108
acoustically active region of the resonator,

the enhancement of long-term aging might

result from this operation.

A pair of thin film electrodes which form the

resonator region are composed of Ti-Pd-Au

tri-metal thin films. As the focussed laser

pulse strikes the top electrode, a small hole

is produced, and material from this region is

melted and evaporated. Most material is

evaporated; however, some material melts,

pulls away and forms a lip surrounding the

hole. Other material is ejected and redeposited in the immediate vicinity of the

hole. An additional problem is encountered

since the quartz substrate is only 200 ym

thick. When the laser beam strikes the top

electrode, it vaporizes in a matter of a few

nanoseconds.

(

23 ) The pulse width of the Qswitched TaG laser is approximately 300 nanoseconds. After the top electrode is

vaporized there is still sufficient beam

energy to vaporize a hole in the bottom electrode (Figure 12 - 'laser exit') even though

the beam is defocussed and scattered by two

'rough' quartz surfaces. The vaporization

of the bottom electrode is markedly different from that of the top electrode, as the Ti layer next to the quartz is vaporized

first while the outer Pd-Au layers are still

momentarily intact. The diameter of the

laser exit hole is (15 ym) considerably

smaller than that produced on the top electrode (25 ym) . SEM micrographs (Figure 12) were taken of both laser entrance and exit

holes. The entrance hole has a minimal lip

formation and shows no obvious damage to the

substrate. The laser exit hole has a some- what larger lip formation, but again no obvious damage to the substrate can be seen.

Figure 13 (insert) shows an adsorbed current

image of entrance laser machined craters.

The bright line across the micrograph indicates the dynamic position of the primary

electron beam which is within a beam diameter

(^ 5 ym) of the laser hole. The point analyses obtained in the vicinity of the crater

defined by the bright line is also shown in

Figure 13. The selected area analysis showed

that Ti, Pd, and Si (in the form of Si0 2 ) are present in the vicinity of the crater, but

are not present about ^ 50 ym away from the

craters as shown in Figure 14. Carbon, an ubiquitous contaminant present in cleaning

solvents and laboratory air, is present over

the entire surface. S and CI, which are com- mon impurities in Pd' 24 '' and Ti( 25 ), are also

detected (see Figure 13). The impurities

CI and S are also common impurities in the

solvents used in cleaning (such as trichloroethylene, acetone and ethyl alcohol) . The

intensity of the Ti, Si and Pd peaks varied

10

in the vicinity of each crater. This is

attributed to pulse-to-pulse variation in

laser intensity and in point-to-point variations in beam attenuation in focussing lens.

The laser is operated at power levels which

are near threshold values to initiate vaporization; thus, small differences in power levels could easily account for the observed

hole-to-hole variations. Auger images of Ti

and Pd could be taken in the vicinity of the

craters, but surface charging due to exposed

quartz (Si02) at the bottom of the crater prevented the taking of well-defined images at that region. The images obtained in the

vicinity of the crater prior to charging indicated that the Ti-Pd debris layer is uniform

and extends approximately 25 ym beyond the crater. This fact and the fact that the high

energy gold transitions (approximately 20 A

escape depth) were also detected indicates

that the debris surrounding the crater is on

the order of 20 A in thickness. It is also

possible to determine from Auger data that

the titanium debris is completely oxidized,

since the oxidation of titanium has a definite

effect on the titanium LMM Auger triplet

(387, 418, 460 eV) . The 418 eV transition

is the most sensitive to oxidation and will

decrease in amplitude relative to the 387 eV peak, which is smaller in amplitude for

unoxidized titanium.

(

25 » 25 ) Selected area in-depth profile analysis showed that the

entire titanium debris layer was oxidized;

i.e., the amplitude of the 387 eV Ti transition

was larger than the 418 eV Ti transition and

the oxygen peak (510 eV) was also high. Re- moval by insitu on sputtering of the debris

layer (^ 20 A) showed complete absence of the

387 and 415 eV Ti peaks, which also was ac- companied by almost complete attenuation of

the oxygen peak.

The analysis in the vicinity of the exit laser

machined holes was very similar to that previously discussed for the entrance holes. The

Ti-Pd debris again extended approximately 25 ym around the exit crater. Thus, from the

debris distribution analysis, the laser exit

hole did not exhibit different characteristics.

VI. SUMMARY

Auger electron spectroscopy is capable of

three-dimensional selected area analysis. Of

the techniques described in this monograph,

with the possible exception of SIMS, it is

the only analytical technique which can perform a selected area analysis on dimensions

which are of interest in integrated circuit

processing. Quantitative analysis is possible

by means of calibration standards, pure standards, and a first order approximation method.

Detectability limits are in the range of 0.1

atom percent. . 9
Future possibilities include the development

of techniques to increase the sensitivity of

Auger electron spectroscopy, the commercial

availability of instrumentation which combine

Auger with SIMS and ESCA, and the development

and routine use of ^0.5 pm electron beam

sizes with sufficient current for practical

analysis times.

REFERENCES

1. Physical Electronics, Inc., Eden

Prairie, Minnesota.

2. Brandis, E. K. , Proc. of Ninth Annual

Conf. Microbeam Analysis Soc. , Ottawa,

Canada, 45A, July 22-26, 1974.

3. Electron Spectroscopy, D. A. Shirley,

Ed. , (North Holland Publishing Company,

Amsterdam, The Netherlands, 1972).

4. Bassett, P. J., and Gallon, T. E., J. Electron Spectroscopy and Related

Phenomena^, 101 (1973).

5. Amelio, G. F. , Surf. Sci. 22, 301

(1970)

6. Grant, J. T., and Haas, T. W. , Surf. Sci.

24, 332 (1971).

7. Morabito, J. M. , Thomas, J. H. , and

Lesh, N. G. , to be published in special

issue of IEEE Parts, Hybrids, and

Packaging

8. Hall, P. M. , Morabito, J. M. , and

Poate, J. M. , to be published in Thin

Solid Films.

9. Shimizu, H. , Ono , M., and Nakayama, K.

Surf. Sci. 36, 817 (1973).

10. Schulz, F., Wittmaack, K., and Maul, J.,

Rad. Effects 18, 211 (1973)

11. Hall, P. M., and Morabito, J. M. , to be

published in Surf. Sci.

12. Dahlgren, S. D.,and McClanahan, E. D,

J. Appl. Phys. 43, No. 4, 1514 (1972).

13. Morabito, J. M., and Tsai, J. C. C.

Surf. Sci. 33., 422 (1972).

14. Morabito, J. M. , Anal. Chem. 46, 189

(1974) .

15. Thomas, J. M., and Morabito, J. M. , Surf.

Sci. 41, 629 (1974).

16. Morabito, J. M. , Surf. Sci. 49, 318

(1975) .

17. Castaing, R. , These de Doctorat,

Universite de Paris (1951) , Publication

Onera (1955).

18. Anderson, C. A., and Hinthorne , J. R. , Anal. Chem. 45, 1421 (1973).

19. Palmberg, P. W. , paper presented at Symposium on Surface and Near Surface

Chemical Characterization, Am. Inst. Metal. Engrs., Toronto, May 18-22,

(1975).

20. Holloway, P. H. , to be published.

21. Schon, G. , Surf. Sci. 35, 96 (1973).

22. Hokanson, J. L.,and Unger, B. A., J. Appl. Phys. 40, 3157 (1969).

23. Masumura, R. A., and Achter, M. R.

Appl. Phys. Lett. 16, 395 (1970).

24. Tracy, J. C, and Palmberg, P. W. , J. Chem. Phys. 5, 852 (1969).

25. Bishop, H. E. , Riviere, J. C, and

Coad, J. P., Surf. Sci. 24, 1 (1971).

26. Morabito, J. M. , Thin Solid Films 19,

21 (1971).

110
TABLE I

THE MINIMUN TIME CONSTANT, Tm REQUIRED FOR AUGER

SURFACE ANALYSIS AS A FUNCTION OF BEAM SIZE (^m 2

) FOR A CONSTANT CURRENT DENSITY (j)

j 10-7 A/^m 2

A A A A

10 W lOVm 2 10 ^m2 1 /xm

2

S/N C(at%) rMIN TMIN TMIN TMIN 10/1 10 1X10*SEC. .1 SEC. ISEC. 10 SEC.

100/1 1 .1 SEC. 1 SEC. 10 SEC. 100 SEC.

j -lO'

6 A/^m2

A A A A 10V 10V 10^m2 Urn 2

S/N C(at%) TMIN. TMIN. TMIN. TMIN.

10/1 10 1X10-

3SEC. .01 SEC. .1 SEC. 1 SEC.

100/

i

1 01 SEC. .1 SEC. 1 SEC. 10 SEC. TMIN.

0C OR

-Jjj111
TABLE ii

CALCULATED COMPOSITIONS OF COMPOUND STANDARDS

SAMPLE A

Rej CALCULATED COMPOSITION* MgO 1.50 56 at. % 0 44 at. % Mg

Cu 2 0 1.34 35 at. % 0 65 at. 7. Cu CdS 1.57 48.5 ol. 7. S 51.5 at. 7. Cd GaP 2.99 47 at. % P 53 at. 7. Ga

KCI 1.18 47.4 at. 7. CI 52.6 at. 7. K

*AIL DATA (EXCEPT MgO) WERE TAKEN WHILE SPUTTERING.

TABLE III

COMPARISON OF BULK SENSITIVITY LIMITS FOR

P, B, AND As IN SILICON AND FOR N, 0, C IN TANTALUM USING AES AND SIMS

ELEMENT TECHNIQUE DETECTABILITY LIMIT

N, O, C AES 1019 atoms/cm 3

N SIMS 10

19 atoms/cm 3

C, O SIMS 10

17 atoms/cm3

P, B AES

As AES 10 atoms/cni

B SIMS ,14

P SIMS lO

17 atoms/cm3

1 Q Q As SIMS 10 atoms/cm

112
Figure 1. Schematic of SAM instrumentation.

Figure 2. Modes of operation for AES.

113
__l I I I I I I I I I I I I I I I I I I—

0 200 400 600 800 1000 1200 1400 1600 1800 2000 ELECTRON ENERGY, eV

Figure 3„ SAM analysis of a silicon nitride passivated layer of a "good" diode prior to and after sputtering. N ^ 55 at.% or Si 3 Nit

.

Figure 4. Auger in-depth profile

of TCNA after 4 hours at 350° C. 123456789 10 11

SPUTTERING TIME (HOURS)

114
Au

Au Au ABSORBED CURRENT IMAGE 350

X

ELECTRON BEAM

POSITION

BEAM LEAD

J I I I I I L. I l I l 0 200 400 600 800 1000 1200 1400 1600 1800 2000 ELECTRON ENERGY eV

Figure 6. Auger spectrum and absorbed current image of contaminated

beam lead priT to in-situ ion sputtering.

Figure 7. In-depth profile of a carbon con- taminated beam lead.

VmulT = 1560

Vmult = 1500

2 4 6 SPUTTERING TIME (MINUTES)

15 20 25 30 CARBON/GOLD AUGER PEAK RATIO

Figure 8. Carbon layer thickness

vs. the normalized Auger peak

heights (C/Au) for carbon escape

depths (dj of 4 A, 6 A and 10 A.

115
Figure 9. Auger analysis after pre-bond

cleaning of gold plated ceramic.

0 500 1000 ELECTRON ENERGY (eV)

24

BULK (ppm at.) LEAD FAILURES PERCENT

Figure 11. Regression analysis of Auger

data vs. lead failures (percent).

Figure 10. Calculated surface concentration

(C ) assuming complete segration to the

surface as a function of bulk concentration

(ppm at. )

116
LASER FINE TUNE

Figure 12. Schematic of laser fine tune process. SEM of laser entrance hole shows

minimal lip formation and absence of substrate

damage. SEM of laser exit hole shows smaller

hole with somewhat larger lip formation, but

no damage to substrate.

LASER EXIT 2000X

-BEAM POSITION

ADSORBED CURRENT MICROGR

0 200 400 600 800 1000 1200 1400 1600 1800 2000

ELECTRON ENERGY. oV

BEAM POSITION

ADSORBED CURRENT MICROGRAPH

0 200 400 600 800 1000 1200 1400 1600 1800 2000

ELECTRON ENERGY, ev

Figure 13. Absorbed current image (insert)

and SAM spectrum in vicinity of laser machined holes

.

Figure 14. Absorbed current image (insert)

and SAM spectrum away from laser machined

region.

117
DISCUSSION OF THE PAPER

Participant : I wonder if you could clarify

exactly which of the impurities was causing

the beam lead bond failures? Was it the carbon, the silver, the tin, or the copper - and how did they manage to get all those impurities in that gold?

Morabito : At different times, different impurities were causing the bond failures.

For the example I showed, it was primarily

the carbonaceous material. Now various

cleaning techniques can be used to remove

the carbonaceous material. When this is

done there is a possibility that other impurities can start to have an effect. As far as the metallic impurities are concerned,

other people have been able to identify that

during the plating operation impurities such

as copper, iron and silver are codeposited.

Silver is actually a part of the gold cyanide

solution. These impurities will rapidly

diffuse to the surface. In fact, for a cer- tain dimension of beam lead (for example

5,000 A in thickness and so many angstroms

in width) you could actually calculate that

100 ppm of copper in the bulk relocated on

the surface would represent a 10 atom percent concentration. This could have a drastic effect on TC bonding and it does.

Dobrott : Would you care to comment on the

fact that all the Auger spectra showed gold

but were conspicuously absent of sodium or potassium. Could you comment on the sensitivity of Auger spectroscopy these elements?

Morabito : I can show you many spectra where

we do see sodium. The one unique thing

about sodium is that it is very rapidly beam

desorbed. Anyone doing Auger spectroscopy

has seen sodium; not routinely, but very

often. Carbon is perhaps the more routine.

I do not really expect sodium to have any

affect on bonding. The sensitivities of

Auger, as indicated, are not as high as those of SIMS. I am really talking seriously

when I say 0.1 atom percent is about the

very best Auger can do.

Participant : Where does your contamination

come from?

Thomas : Very often you find carbon on top

of the gold surface. The explanation which

is often given is that carbon diffused to the gold surface very readily, even though

there is only a small amount of carbon present in the bulk. Even at moderately low

temperatures it can come to the surface.

Even if there is no hydrocarbon contamination you will see some carbon on top of the

gold.

I would also like to agree with you about the

bonding problem. The most frequently recurring problem is really due to the oxides of

copper, nickel or iron, rather than the carbon. All of our experiments show that carbon

is not really a bonding problem as such, if

it is less than 20-25 A thick, as you said.

The real problems are nickel oxide, copper

oxide, or iron oxide. Some of the leads

have gold plate on copper; the copper can then diffuse through the gold. The copper

can also be codeposited in the gold plate,

as in your case. There are many possibilities, however carbon is not a problem as far

as I can see.

Morabito : In my particular case, the carbon

did not come from the plating film.

Whitf ord : Many semiconductor manufacturers

use, as part of their processing, the technique called low temperature ashing, they

also do etching work with it. It has been

used by the connector industry when they

want a good bond and a hermetic seal to glass, rubber or polymer. Just a quick

spritz for about 5 minutes in a nitrogen

atmosphere at room temperature in a RF plasma

insures the bond and takes off any carbon

films or any materials that have diffused to the surface. I know the technique has been

used in some electron microscopy work instead

of ion sputtering to clean surfaces, and it

is a lot cheaper too.

Morabito: We are very familiar with that

technique.

118
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg , Maryland, April 23-24, 1975 (Issued March 1976)

USE OF AUGER ELECTRON SPECTROSCOPY TO DETERMINE

THE STRUCTURE OF SILICON OXIDE FILMS

J. S. Johannessent and W. E. Spicer

Stanford Electronics Laboratories

Stanford University

Stanford, California 94305

and

Y. E. Strausser

Varian Associates

Palo Alto, California 94305

In the present paper we present experimental

results on phase separation in unsaturated

silicon oxides 1 and the chemical structure

of the Si02 -Si interface. 2

The chemical structure of Si0x

is a controversial subject which has received considerable attention in the past. From interpretation of experimental observations, two chemical structure models have emerged.

These models are commonly referred to as the microscopic mixture model 3 and the random

bonding model. ^

The microscopic mixture model (MMM) states

that unsaturated silicon oxides SiO^, where

0 < x < 2 is a stoichiometry factor, are

phase separated mixtures of amorphous Si and

amorphous Si02. Interpretation of radial

distribution functions of films of SiO

x give

support to this model. 3 ' 5

Optical reflectance data from unsaturated

silicon oxides are interpreted in terms of

the random bonding model (RBM) 14 , which

states that the local atomic arrangement in

SiOxis statistically distributed among five

coordinations of Si atoms.

Figure 1 shows how these two models predict

the relative amount in each coordination C

z as a function of sample stoichiometry, given

by x.

3 The MMM is represented by the broken

line. The RBM, represented by full drawn

curves, is based on a statistical distribution of tetrahedra of the type Si - (Siz-! 0 5 _ z ), where z = 1, 2, 3, 4, 5. The

relative amount in each coordination C

z is determined by the overall stoichiometry, as shown in figure 1. Both models are identical

for x = 0 and x = 2.

On leave from Electronics Research Laboratory (ELAB) , University of Trondheim,

Norwegian Institute of Technology, Norway.

In the present paper we present experimental

evidence of phase separation of Si02 and Si in unsaturated silicon oxides of stoichiometry SiO

x

, where x < 2. We make use of high

energy resolution Auger electron spectroscopy

(AES) in a study of evaporated films of SiOx and thermally grown Si02

. Significant chemical structure effects are observed in both

the LVV and KLL series of Auger transitions.

However, we limit the discussion to the Si^-^

spectra because these spectra reflect directly

the effect of the changes in the local atomic

arrangement around each silicon atom as a result of changes in the overall stoichiometry.

Chemical structure effects or chemical shifts

are well known from XPS and ESCA analysis of oxidized silicon6 and evaporated films of silicon oxide 7

, and appear as a 3-4 eV shift

in the binding energy of the Si(2p) energy

level. We see similar chemical shifts in the silicon KL2 3^2,3 Auger transition as indicated in figure 2. In figure 2 we show two experimental transition energies 1618 eV corresponding to elemental silicon x = 0, and

1611 eV corresponding to silicon in Si02,

x = 2. The straight line between these two values gives an indication of the expected

chemical shift in the silicon ^2^312^ tran- sition energy as the local coordination around

each silicon atom changes. We use the extra- polated Auger transition energies in figure 2 and the coordination factors in figure 1 to calculate the energy dependence of the silicon KL2 3L2 3 Auger transition of the two structural models. The result for two stoichiometrics is shown in figure 3 where we, for convenience, have chosen a Lorentzian

shape of each of the Auger transitions. The

half width in this case is slightly less than

the experimental half width. The broken

curves are the contribution from each of the

separate silicon coordinations. There is a pronounced difference between the two models.

The MMM shows two distinct Auger transitions,

where the peak-to-peak heights depend on overall stoichiometry. The RBM, shows one

119
broadened peak, whose peak-to-peak height

and energy position depends on the stoichiometry factor x.

In figure 4 we show a series of experimental

Sj-KLL spectra obtained from a thermally

grown oxide x = 2, and from evaporated oxide

films 1.2 < x < 1.9. The Si (100) substrate

is used as reference x = 0. Notice that the

spectra for x = 0 and x = 2 are similar, but

displaced in energy by AE ^ 7 eV. This is

the chemical shift due to change in local

coordination of the silicon atom. Similar

shifts are reported for oxidized silicon. 8 The Auger spectra for intermediate values of x are significantly different. This is

particularly pronounced for the lowest value

of x, which show two distinct transitions of

1611 eV and 1618 eV. We were not able to obtain SiO, probably due to a finite 0 2 background during deposition at 10-7 Torr. The experimental conditions are described

elsewhere.^ The stoichiometry factor x is

estimated using the peak-to-peak height

normalized to the thermally grown oxide.

The experimental results presently in figure

4 strongly support the MMM, as can be seen when compared with figure 3. Based on the

above discussion and our experimental data we are led to the conclusion that unsaturated silicon oxides SiOx , where 0 < x < 2, are phase separated and best described by

the microscopic mixture model.

We have also performed chemical depth profiling of thermally grown Si02 on Si (100)

In figure 5 is shown a typical depth profile

of an oxide grown to a thickness of 1000 1

at 1200 A in dry 02. After oxidation the

sample was annealed in dry N 2 and slowly

pulled from the furnace. Figure 5 is obtained by automatic multiplexing of the

peak-to-peak heights of the Ojy^ (502 eV)

SiKLL (!611-1618 eV) , and SiLW (78-92 eV)

Auger transitions. 10 The data points to the

left of the origin are steady state values

before the onset of the 1 keV Argon ion

beam. The small step in the Ojq^ and Si^yy

transitions is most probably due to removal

of surface contaminants such as water vapor

and carbon. The depth profile of figure 5 shows a uniform Si02 film, except in the

interface region when there is a dip in the

silicon KL2j 3L 2 3 peak-to-peak height. The

reason for this dip is clearly shown in the

actual real time Auger spectra plotted above

the depth profile. The numerals indicate

the position of the spectra in the interface

legion of the Si02 -Si depth profile.

There are two characteristic features of the

depth profile near the interface. Firstly,

the Sijy^ transition "sees" the interface

before the and Si^yy transitions, and

secondly it "sees" a broader interface than

the two other transitions. These features

are related to the escape depth of the Auger

electrons and to the nature of the KLL and

LW Auger transitions. 2

When the electron escape depth 5 ' 11 is ac- counted for, we find that the width of the

Si0 2 -Si interface is on the order of 35-50 A

wide. Interface broadening by the ion beam

is estimated to be on the order of 10 to 15 A.. 2

f Hence, we are left with a "real" interface

width on the order of 20-30 A.. In the interface region, the average stoichiometry is

SiOx

, x < 2, and the material appears to be

phase separated in a similar manner as un- saturated silicon oxides.

In conclusion, we mention that we have made

the first observation of chemical shifts in

the silicon KL 2 3 L2 3 Auger transition in

unsaturated silicon'oxide films. We interpret our results in terms of a microscopic

mixture of Si02 and Si. Chemical depth profiles through thermally grown Si02 on Si (100)

show a phase separated Si02 -Si interface of

finite width on the order of 20—30 A. We

believe that the positive oxide charge observed in MOS-devices 12

, are confined to the

connective regions between the two phases at the interface, Si0 2 and Si, respectively.

ACKNOWLEDGEMENTS

Work in part supported by Army Research Office,

Contract No. DAHC 0474 G 0215. One of us (J.S.J.) has been partly supported by ELAB

and the Royal Norwegian Council for Scientific

and Industrial Research.

REFERENCES

1. Johannessen, J. S., Spicer, W. E. , and

Strausser, Y. E. , to be published.

2. Johannessen, J. S., Spicer, W. E. , and

Strausser, Y. E. , to be published.

3. Temkin, R. J., J. Non-Cryst. Solids 17,

215 (1975).

4. Phillips, H. R. , J. Non-Cryst. Solids,

8-10, 627 (1972).

5. Coleman, M. V., and Thomas, A. J. D.,

Phys. Stat. Sol. 22, 593 (1967).

6. Flitsch, R.,and Raider, S. I., J. Vac. Sci. Technol. 12, 305 (1975).

120
Hollinger, G. , Tousset, J., and Tran

Ming Due, Tetrahedrally Bonded Amorphous Semiconductors, M. H. Brodsky,

S. Kirkpatrick, S. and D. Weatre, Eds.,

p. 102 (Am. Inst. Phys., New York,

1974)

.

Wagner, C. D. , Anal. Chem. 4_7, 1201

(1975)

.

Strausser, Y. E., and Johannessen, J. S.

,

An Auger Electron Spectroscopy Study of

Silicon Spectra From Silicon Monoxide,

Silicon Dioxide and Silicon Nitride,

ARPA/NBS Workshop IV, Surface Analysis

for Silicon Devices, A. G. Lieberman,

Ed., NBS Spec. Publ. 400-23 (March 1976).

10. The notation Si^yv ^ s use d to show that

the LW spectrum of silicon oxides is

composed of molecular orbitals originating both from Si-Si bonds and Si-0

bonds

.

11. Lindau, I., and Spicer, W. E. , J. Elec.

Spectr. 3, 409 (1974).

12. Deal, B. E. , J. Electrochem. Soc. 121 , 198C (1974).

121
0.5 10 1.5 STOICHIOMETRY FACTOR X

Figure 2. The energy of the Si.,, L KL2>3 2.3

Auger transition is plotted as function of x the stoichiometry factor of silicon oxides.

The two circles are experimental values.

Also shown is the calculated excess charge

on the silicon atom, according to the Paulinj

electronegativity theory. The straight line

behavior is an approximation.

Figure 1. coordination Si-(Si z _^ 0^_ z

) is shown as functions of x the stoichiometry factor of

silicon oxides. The broken line represents

the microscopic mixture model (MMM) and the

full drawn curves represent the random bonding model (RBM)

EXCESS CHARGE ON Si (units of e)

0.5 1.0 1.5 2.0

1 1

1 1

1618 < >r SH(Si4

)

1616 S:Si-tSi 30 1

)

\ Si-(Si 202l

1614 1 \.

1 \s

Si-(Si

1 03

)

1612 1

1

1

1

v Si-(04)

1610

1 1 , 1

1

1

1

1 1 1 1

1

1 "1

0.4 0.8 1.2 1.6 2.0 2.4 STOICHIOMETRY FACTOR X

1611 1618 1611 1618

Si KLL TRANSITION ENERGY (eV)

Figure 3. Model calculation of the

SivT t Auger transition in two unsatu-

^2,3^2,3

rated silicon oxides, x = 1.0 and x = 1.5,

according to the microscopic mixture model

(MMM) and the random bonding model (RBM)

The density of the material is not corrected

for in these calculations.

122
O S

Figure 4. Experimental SiKLL spectra of

silicon oxides. Thermally grown oxide is

denoted by x = 2. Vacuum deposited silicon

oxides are x = 1.9, 1.6, 1.4, and 1.2. The

silicon (100) substrate is used as reference

for x = 0.

' ' i i i i i i i i i i i i i i_ 1500 1550 1600 1650 (eV! AUGER ELECTRON ENERGY

It 11

Figure 5. Chemical depth profile through

1000 A thick thermally grown Si0 2 on Si (100)

substrate. Above the depth profile is shown

the change in the Si L transistion in

KL2»3 2>3

the interface region between the dioxide and

the substrate.

sirLVV

too0 ' 0003° OOODoQgO^Ot ,000° aoooot^ooo!*5 oO°d3oo rf^rf1^)0!

I I I I I V-r/

0 SILICON (LINDAU 8. SPICERI A SILICON DIOXIDE I IFLITSCH & RAIDER]

3 6 9 12 (min) SPUTTERING TIME

Figure 6. The energy dependence of the electron escape depth in Si and Si02- The broken

part of the curve is an extrapolation by the

authors. The data are collected from the

literature (6,11).

100 ELECTRON ENERGY (eVI

123

NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

AN AUGER ELECTRON SPECTROSCOPY STUDY OF SILICON SPECTRA FROM

SILICON MONOXIDE, SILICON DIOXIDE, AND SILICON NITRIDE

Y. E. Strausser

Varian Associates

Palo Alto, California 94305

and

J. S. Johannessen

Stanford Electronics Laboratories

Stanford University

Stanford, California 94305

INTRODUCTION

AES, in conjunction with ion sputtering, is of great value in semiconductor technology,

for determining impurity doping profiles and

chemical stoichiometry of overgrown layers

on semiconductors. The technologically most important overgrown layers today are silicon

oxides and silicon nitride on silicon substrates.- The electronic properties of these

insulating films are strongly dependent on their stoichiometry. Only small deviations

from Si02 stoichiometry of vacuum deposited

silicon oxide is detrimental to its properties as an insulating layer [1].

The general purpose of the work reported in

this paper is to explore the use of AES and

ion sputtering in determining the chemical

states of these overlayers, the relative

amounts in each chemical state, and how this

is changed by the AES measurement itself.

Particular emphasis is put on chemical structure effects, electron beam distortion of the

sample composition and ion beam effects during simultaneous ion-sputter etching. By

chemical structure effects we mean all of the

changes that occur in the silicon Auger spectra as a result of the chemical environment

of the Si atoms.

It is well known from photo-electron spec- troscopy (PES) of silicon and silicon oxides

that both the valence band spectra and the core levels are effected by the presence of an oxide layer on silicon [2, 3, 4]. Similar

"chemical shifts" are also observed in Auger

spectra of silicon in a large variety of chemical environments [5].

* Permanent Address: Electronics Research

Laboratory, Norwegian Institute of Technology, University of Trondheim, Trondheim,

Norway

The chemical structure effects may be grossly

understood on a basis of charge transfer ac- cording to the Pauling electronegativity

theory [6], although the details of the structure of the valence bands of silicon

oxide and silicon nitride are determined by

the chemical bonds.

EXPERIMENTAL

In the present analysis we have used an Auger

spectrometer with a CMA of high energy resolution, in conjunction with argon ion sputtering for depth profiling of the oxide and

nitride films. The samples investigated are pure silicon, silicon oxides and silicon

nitride on silicon (100) substrates of n- type doping (1 < p < 5 Q cm)

.

Experimental Apparatus

The measurements reported in this paper were made in a Varian Auger spectrometer. This

system uses a cylindrical mirror analyzer

(CMA) with an energy resolution of 0.25% and

a transmission of > 8% of the electrons emitted from the sample. The primary electron

beam was produced by an electron gun mounted

inside the inner cylinder of the CMA. It is capable of producing electron beams at energies up to 10 keV and with diameter less

than 5 p. The important system parameters

used in these measurements were as follows:

a) Si LVV, N KLL and 0 KLL spectra: All of these spectra were obtained using a pri- mary electron beam of 5 yA total current

at 2 keV. The electron beam diameter

was < 20 um FWHM. The sample holder

used held the samples so that there was

an angle of 30° between the sample surface normal and the axis of the CMA. The modulation used to obtain the differential spectra was equivalent to 1 eV peak-to-peak. The sweep rate and the

lock-in-amplif ier time constant were varied in such a way that we always

125
swept - 1 eV per time constant.

b) Si KLL spectra: These spectra were taken

with a primary electron beam energy of

4.5 keV and a current of 5 yA. The beam

diameter was < 15 ym FWHM. In these

spectra, a modulation equivalent to 2 eV peak-to-peak was used. The other parameters were the same as in a)

There are two other features of the electron

gun which are important to these measurements. First, the use of the scanning sample

positioner, normally used to scan the electron beam to make absorbed current images,

enabled us to record spectra while the electron beam was being scanned over small, well

defined areas of the sample. In this way, we could control the average beam current

density incident on the sample and obtain

data on the effect of total electron beam exposure. Additionally, this removed the effect of the current density distribution

through the electron beam.

The second important feature of the electron

gun is the existence of a deflection plate

between the two lenses of the gun which permits one to electronically chop the primary

electron beam. Pulsing the electron beam

then makes it possible to directly record

the energy distribution curves under the same conditions as the derivative spectra were

taken.

The ion gun used in the profiling measurements produces a beam with ion current densities from 0 to > 200 uA/cm2 at energies up

to 3 keV. As in the case of the electron

gun, the ion beam can be scanned over well

defined areas of the sample to provide control of the ion beam current density and to remove crater shape effects. The geometry

used resulted in an angle of 53.5° between

the ion beam and the sample surface normal.

Sample Preparation

Saturated silicon oxides (Si02) were grown

by thermal oxidation of Si in a dry O2 atmosphere at temperatures between 1000°C and

1200°C, in the thickness range from 250 & to 1000 A. Slow and fast cooling in, and without, a protective N2 atmosphere was used.

Non-saturated silicon oxides SiO^, where

0 < X < 2, were vapor deposited in a residual

oxygen atmosphere ranging from 10-7 Torr to 10~ L> Torr, on silicon (100) substrates held

at 210°C or 240°C. The thickness of the SiOx

films has been kept approximately constant

at 1000 A. The silicon nitride films were

grown by CVD at approximately 750°C for a large variation of silane to ammonia ratio.

All samples have been exposed to air for at least several hours before each AES analysis.

All the overgrown layers discussed in this

paper are X-ray amorphous. That is, they

possess no long range crystalline order.

CHEMICAL STRUCTURE EFFECTS

We have observed pronounced chemical structure effects in the Auger spectra of silicon

in compounds such as silicon oxides and silicon nitride. By chemical structure effects, we mean all of the changes that occur in the

Auger spectra of Si due to changes in the

chemical environment of silicon atoms in the

compounds

We have studied the KLL and LVV series of

Auger transitions in pure silicon, and of

silicon oxides, and silicon nitride grown on

Si (100) surfaces. The presentation and

discussion of the results will be divided

into three parts. First, we discuss Si (100)

spectra, then silicon oxides of stoichiometry

SiOx , where 0 < X - 2, and finally we deal

with silicon nitride, Si3N4-

Pure Silicon

AES spectra of silicon are well known, hence

we shall only briefly discuss the characteristic features. of the KLL and LVV series of

transitions and relate them to the electronic

structure of silicon. The gross features of

the band-structure of silicon are determined

by the local tetrahedral coordination of each

Si atom, and the co-valent Si-Si bonds, giving no net charge on any individual Si atom

(according to the Pauling charge transfer

theory)

The series of peaks in the KLL spectrum of

Si is shown in Figure la and lb, where the

Auger current, 1^, is differentiated with

respect to the energy E in Figure la, and

Figure lb is a plot of 1^. versus electron

energy (the background current has been subtracted) . There are three distinct peaks in

the 1^ spectrum one at 1562 eV corresponding

to KLjL2 transitions, one at 1602 eV corre- sponding to KL2L2 transitions, and one at 1618 eV corresponding to KL3L3 transitions,

in good agreement with theoretical predictions [7]

The LVV spectrum of silicon is most commonly

presented in differentiated form (dl^/dE)

showing a strong peak at 92 eV [8]. In Figure 2 are shown both dl^/dE (a) and 1^ (b) as a function of energy for a silicon sample

where the natural oxide has been removed by

argon ion-sputtering. The spectrum shown in

Figure 2 is composed of L2,3W transitions

126
at higher energies and LjI^^V transitions at the low energy side of the spectrum. Detailed analysis and identification of the different transitions is difficult due to the com- plexity of the silicon valence band. However, some of the characteristic features of

the valence band are reflected in the LVV

series of Auger transitions. Using Kane's

density of state calculations [9], by assigning the binding energies V]_ = 3 eV, V2 = 4 eV,

V3 = 7 eV, and V4 = 10 eV to the four major

peaks in the density of states, and using the

binding energies EL ^ = 148.7 eV and El2.3 = 99.2 eV, we are able to give a crude identification of the features of LVV Auger spectrum of Si. This is the same approach as that of Maguire [10]. The major peak in

Figure 2b, at 88 eV, is due to L2 3V

transitions, while the minor peak is due to L2 3V4V4 transitions at 72 eV. The shoulder

between these peaks is largely L2 3V3V3

transitions at 79 eV. The two small pieces

of structure on the low energy side are LjL2 3V2 2 transitions at y 41 eV and

L 1L 2,3V3,4 transitions at y 35 eV. The details of the LVV spectrum are amplified by

differentiation w.r.t. energy as shown in

Figure 2a.

Silicon Oxides

Silicon oxides and oxidized silicon have

been subject to extensive studies in the

past, both by AES and photoelectron spectroscopy (PES) [2, 3, 11]. However, most of

the work has been focused on either silicon

dioxide or oxygen chemisorbed on clean silicon. In this paper, we report on results of studies of thermally grown silicon oxide of

stoichiometry Si02, and evaporated films of

stoichiometry SiO^, 1 % X a, 2.

In Figure 3, we show a series of Si LVV

spectra of silicon oxides (a) , and oxygen

KLL spectra (b) for the corresponding compounds. The stoichiometry of the samples

was determined using the amount of oxygen

present in the films, as indicated by the

peak-to-peak heights of Figure 3 (b) . The

oxygen spectra of Figure 3 (b) indicate that

oxygen is present in only one chemical state,

namely as a bridge between two silicon atoms

[12]. Furthermore, the coordination of oxygen determines the major features of the LVV

spectrum through the density of states in

the valence band of silicon oxides. The two major peaks in the LVV spectra of silicon at

64 eV and 78 eV are associated with transitions from bonding and non-bonding molecular

orbitals of the Si-O-Si (Si20) quasi-molecule. The transition probabilities are apparently determined by the oxygen 2p predominance of these orbitals. The peak appearing

at 92 eV is due to the presence of "free"

silicon, i.e., silicon coordinated by four

nearest silicon neighbors. A small amount

of "free" silicon in silicon oxides is in

this case produced by ion sputter cleaning.

In fact, 500 A of material had been removed

before the AES spectra were taken.

It is important to understand that the ratio

of the peak-to-peak heights of the 78 eV and

92 eV peaks in the ^ spectra does not give,

directly, the ratio of the concentrations of

Si bound to 0 to the concentration of Si

bound to Si ("free" Si) . Observation of the N (E) spectrum shows that the peaks at 64 eV and 78 eV are very broad whereas the 92 eV peak is narrow, with a very sharp high energy

edge. This, of course, results in a large

negative excursion on the 92 eV gS peak without the corresponding large number of electrons in the peak. In fact, N (E) spectra

corresponding to a condition like that shown

in the spectra of Figure 3 will show the

92 eV peak to be only a small, although

sharp, step on the high energy side of the

large, broad 78 eV peak. This does not,

however, prohibit the calibration of the LW

spectrum to determine the relative concentrations of these two states of Si. On the

other hand, the KLL spectra give much more

obvious and direct information about the

relative concentrations of the different

chemical states of Si.

The local atomic arrangement around the silicon atoms is shown by the Si KLL spectra in

Figure 4, for silicon oxides of varying

stoichiometry. For X = 2 (Si02) there appears a well defined KL2 3L2 3 peak at 1611

eV, associated with Si-(d4) coordination.

For X ^ 1.8 there appears a small shoulder

above 1611 eV, and for X ^ 1.3 a broad peak

occurs at 1616 eV. This peak may be associated with silicon tetrahedra of the type

Si - (Siy 04_y) , where y 4- 0, 4. However,

upon electron exposure a single peak emerges

at 1618 eV corresponding to Si - (Si^) tetra- hedra (or "free" silicon)

.

Silicon Nitride

In Figure 5a and b, we show silicon nitride

Auger spectra in the energy range between

25 eV and 525 eV. Figure 5a is the spectrum

from a sample exposed to atmosphere. Note

the presence of carbon and oxygen KLL tran- sitions, the large silicon nitride LVV peak

at 85 eV, and the small silicon oxide peak

at 64 eV. Removal of carbon and oxygen by

ion-sputter etching results in the clean

silicon nitride spectrum shown in Figure 5b. It is interesting to note that oxygen is

associated with silicon oxide in the

127
contaminated surface, and that removal of

silicon oxide results in a "free" silicon

peak at 92 eV, while there is no observable

change in the 382 eV N KLL peak. This indicates that nitrogen exists only in one chemical state associated with a silicon nitride

quasi-molecule (Si-jN) , where the Si-N bonds

have less ionic character than the Si-0 bond.

This observation is confirmed by the Si KLL spectrum of the nitride which has a well defined peak at 1618 eV, apparently identical

to the "free" silicon Si-(Si^) coordination,

indicating a negligible charge transfer from

Si to N in silicon nitride. We are not aware of X-ray photoemission spectra of

Si3N^ or valence band structure calculations

so an analysis beyond this mere descriptive

one is not pursued here.

We have shown above that the Si3N4 LW spectrum is strongly influenced by oxygen adsorbed on the surface. Oxygen adsorption

completely removed the 92 eV "free" silicon

peak. This phenomena suggests that amorphous

silicon nitride is a microscopic mixture of

Si and Si 3N4.

ELECTRON STIMULATED DESORPTION EFFECTS

We have already mentioned in the previous

section that the stoichiometry of silicon

oxides may be changed by electron exposure.

The effect is to reduce the oxygen content in

the sample. This phenomenon is commonly

known as electron stimulated desorption (ESD)

and is thoroughly reviewed in the literature

[13]. We have investigated electron beam

exposure effects on Auger spectra of SiO„

(1.2 < X S 2) and Si 3N4 films on Si (100)

single crystals. Si3N4 appears to be stable

to electron exposure at energies less than

2 keV. Apart from desorption of surface con- taminants (carbon and oxygen) the LVV and

KLL spectra appear unchanged upon prolonged

exposure. Silicon oxides on the other hand

are more volatile, and we will restrict our discussion to these compounds.

Si02 is known to decompose under electron

irradiation. It has been shown that oxygen

desorption from Si02 is accompanied by the

formation of "free" silicon [14]. Our findings are similar for SiO. However, where

the stoichiometry of silicon oxide films is

SiOx with 0 < X < 2, this effect is much

more pronounced. In Figure 6a and b, we show the growth of the 92 eV Si LVV peak as a function of exposure time for SiOi.4 (a)

and Si02 (b) , the parameter is the electron

flux density, and the primary electron energy is 2 keV. There appears to be a linear

increase at small exposure times followed by

a saturation region. The linear slope and

the saturation level both increase with increasing flux density. The former being re- lated to decomposition of SiOx followed by

desorption of oxygen. The saturation level

is most probably determined by oxygen depletion in the surface layer and out diffusion

of oxygen from deeper lying layers. Note

that the saturation level in SiO^^ is more

than three times that of Si02.

Decomposition and desorption of oxygen is

shown in Figure 7, where the 0 KLL peak-topeak height is plotted against exposure time

for the same samples discussed in Figure 6. We find that in order to minimize electron

irradiation distortion of silicon oxides,

the dose should be kept at 10 19 electrons/cm

or below, at 2 keV primary energy. This

corresponds to 2.5 min. exposure at 10-2

A/ cm2

.

CHEMICAL DEPTH PROFILES AND ION BEAM EFFECTS

Chemical depth profiles are obtained by re- cording AES spectra as a function of time

during simultaneous ion sputtering. The

technique is well known, and we shall re- strict the discussion to some of the factors

that limit the depth resolution.

In Figure 8, we show a typical depth profile

through a 1080 A thick layer of SiOx

(X - 1.4) vacuum deposited on Si(100) . Only the

initial and the final part of this particular profile are shown. The initial part of

the curve shows how adsorbed carbon is re- moved, and how a balance between bound oxygen and "free" silicon is reached by oxygen

removal. The saturation level gives a measure of the stoichiometry of the sample. The

final part of the profile shows the interface between film and substrate, and the

finite transition region between the two. In this case, the width is 29 A (defined as distance between 80% and 20% of the oxygen

peak-to-peak height) . The question is now

what factors control and determine the observed width or depth resolution. A recent

paper by Ishitani et.al. [15] deals with

these problems in model calculations of

knock-on atomic mixing. They find a depth

resolution decreasing with increasing ion

energy, and increasing with incidence angle

of the ion beam between 0° and 60°. An ex- perimental investigation of Ta205 on Ta [16]

has shown a minimum transition width of ^ 60-70 A for ion energies between 6 and

7 deV. We have employed A+ ions at 3 keV of

varying current density. In Figure 9, we show the transition with AY as a function of

ion power density for silicon oxides and

silicon nitride on Si (100). The general

trend is that AY decreases with increasing

128
power density for both compounds, towards a broad minimum or saturation level of the oro der of 30 to 40 A. This value is roughly

50% of that reported for Ta205 on Ta [16].

Figure 9 also shows that the nitride interface is sharper than the oxide interface at low power densities. We find that the sputtering yield for silicon nitride is 0.5 molecule/ion, and for silicon oxide ^ 1.5 molecule/ion, and an ion energy of 3 keV. We

find our experimental results in qualitative

agreement with the knock-on mixing model

[15], when we take into account the lower

ion energy and the angle of incidence we

have used. The lower limit of 30 - 40 A interface width may be reduced even further by

reducing the ion energy. However, other

factors such as surface roughness, preferential sputtering, and inhomogeneities may

set an upper limit to the depth resolution.

The ultimate limit is set by the escape depth

of the Auger electrons, which will be discussed in the next paper.

CONCLUSION

The results presented in the preceeding sections may be summarized in the following

conclusions. The electronic structure of

silicon compounds is largely determined by

the molecular bonds and the local atomic ar- rangement. Thus, in silicon oxides it is

the Si-0 bond and in silicon nitride it is

the Si-N bond that gives the major contribution to the electronic structure. The Si-0

bond energy is larger than Si-Si bond*. The

charge transfer associated with the Si-0

bond gives rise to a chemical shift in the

binding energy of the Si L2 } 3 level of 3.9 eV [18]. This gives rise to a shift in the

Si KLL transition energy of 7 eV from

Si-(Si4) coordination to Si-(04> coordination

[18]. We have observed varying amounts of

these two states in electron irradiated SiO x

films, X < 2. Similar chemical shifts have

not been observed in silicon nitride. We

believe that this is due to a much smaller

charge transfer associated with the Si-N

bond.

Electron distortion of silicon oxides by

dissociation and desorption of oxygen can be

reduced to a negligible level by keeping the

total electron dose below a limit of 10 19 el/cm2 at 2 keV electron energy. Silicon

nitride is stable to electron irradiation

within the experimental conditions of this

investigation. Why this is so is not fully

understood. It may be associated with the

higher ionization potential of nitrogen com- pared with oxygen.

profiling of silicon oxides and silicon nitride. The mixing effect results in reduced

depth resolution. By varying the ion flux we were able to reduce the width of the interface between silicon oxide or silicon nio tride and silicon, to 30 to 40 A at an argon

ion energy of 3 keV. This result is en- couraging in that it indicates the possibility

of studying the interface region of M0S and

MNOS devices where charge is accumulated.

ACKNOWLEDGEMENT

We would like to thank Forrest Futtere for

preparing the SiOx

samples, and Krishna

Saraswat for supplying us with thermally

grown Si02 and reactor grown Si3N4 samples.

REFERENCES

1. Hirose, H.,and Wada, Y. , Japan. J. Appl.

Phys. 3, 179 (1964).

2. DiStefano, T. H., and Eastman, D. E.

,

Phys. Rev. Lett. 27_, 1560 (1971).

3. Hollinger, G. , Tousset, J., and Tran

Minh Due, Tetrahedrally Bonded Amporphous Semiconductors, M. H. Brodsky,

S. Kirkpatrick and D. Weaire, Eds.

,

p. 102 (Am. Inst. Phys., New York,

1974)

.

4. Danyluk, S., and McGuire, G. E., J. Appl.

Phys. 45, 5141 (1974).

5. Stupian, G. W. , J. Appl. Phys. 45_, 5278

(1974)

.

6. Pauling, L. , The Nature of the Chemical

Bond, 3rd ed. , (Cornell Univ. Press,

Ithaca, New York, 1960).

7. Siegbahn, K. , Nordling, C. , Fahlman, A.,

et. al. , Electron Spectroscopy for

Chemical Analysis, Technical Report

AFML-TR-68-189 , Inst. Phys., Uppsala,

Sweden (1968).

8. Grant, J. T.,and Haas, T. W. , Surf. Sci.

23, 347 (1970).

9. Kane, E. 0., Phys. Rev. 146, 558 (1966).

10. Maguire, H. G., and Augustus, P. A., J. Phys. C. 4_, L174 (1971).

11. Ibach, H., and Rowe, J. E . , Phys. Rev. B10 , 710 (1974).

We have observed knock-on mixing of elements

caused by energetic ions during depth

129
12. Johannessen, J. S., Spicer, W. E. , and

Strausser, Y. E. , Use of Auger Electron

Spectroscopy To Determine The Structure

of Silicon Oxide Films, ARPA/NBS Workshop IV, Surface Analysis for Silicon

Devices, A. G. Lieberman, Ed., NBS

Spec. Publ. 400-23 (March 1976).

13. Madey, T. E.,and Yates, Jr., J. T. , J. Vac. Sci. Technol. 8_, 525 (1971).

14. Thomas, S. , J. Appl. Phys. 45^, 161 (1974).

15. Ishitani, T.,and Shimizu, R. , Appl.

Phys. 6, 241 (1975).

16. Ishitani, T. , Shimizu, R.,and Tamura,

H., Appl. Phys. 6, 271 (1975).

17. Roberts, J. D. , Stewart, R. , and

Caseuio, M. G. , Organic Chemistry

(W. A. Benjamin, Inc. , Menlo Park,

California 1971).

18. Wagner, C. D. , to be published.

130
1500 1520 1540 1560 1580 1600 1620 1640

E (eV)

Figure 1. The Si KLL spectrum (B) and the

differentiated Si KLL spectrum (A). Also,

shown as vertical lines are theoretical Auger

transitions [10].

Figure 2. The Si LVV spectrum (B) and the

differentiated Si LW spectrum (A) . The

vertical lines indicate the most probable

LVV transition based on the density of states

in the valence band of silicon.

(Bl

0 20 40 60 80 100 120

E (eVI

131
Figure 3. LVV spectra of SiO films (A) and KLL spectra of oxygen in

SiO films (B). The oxygen peak-to-peak height is used to calculate the

stoichioraetry of the samples. 500 A has been removed by argon ion

sputtering.

132
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P*4 ^3 4-1 135
DISCUSSION OF THE PAPERS BY STRAUSSER AND JOHANNES SEN

Thomas : Without resurrecting the argument

about the assignment of the Auger peaks of

elemental silicon, I have to say that two of

the peaks on the low energy side of your

major KLL peak are not really Auger peaks

but rather plasma loss peaks; the same is true of the 107 eV silicon peak you have

labeled as an Auger peak, but there is some controversy over that too. That is the

point that I had to make, but could you also

tell me how you determined the stoichiometric

factor x in your SiOx?

Strausser : As to your comment first of all,

I certainly agree that those are plasma loss

peaks in the silicon KLL spectrum. We simply didn't label them. For the 107 eV peak, we are inclined to accept the interpretation

by Grant and Haas, namely, that it is a doubly ionized silicon Auger transition.

Now as to your question, we have done some quantitative work using normalization factors, which is basically like Joe Morabito's

second technique. He normalizes all the

peaks, applying an empirical correction fac- tor to the major peaks, normalizing them all

to one particular peak and dividing each of

them with a total signal, thus calculating

percentages. Instead of experimentally deriving the normalization factors, we have

calculated them. We use calculated ionization cross sections, a linear correction for

the energy dependence of the CMA transmission, measured and extrapolated fluorescent

yields and escape depths. We combine all of

these into a normalization factor. We can then calculate the stoichiometry , and we have verified the stoichiometry of thermally

grown silicon dioxide, for instance, that it

'is Si02- However, on that particular figure we simply used the oxygen peak heights to determine the stoichiometry of the subdioxide

films using a known Si02 film as a reference

for the suboxides.

Thomas : During the past two or three years,

I have been looking, day in and day out, at "Si02" whatever that may be. We have prepared Si02 by various techniques. If I were

to measure just the ratio of the oxygen peak

to a silicon peak, you would notice that the

ratio would change, depending upon which

technique were used to prepare the oxide.

In other words, you would have a different

value for quartz, dry oxides, or oxides

grown in steam, or spin-on oxides, or whatever it is; the peak ratio can change significantly. So I do not know whether you

take the value two because you always say,

Si02, and particularly how do you get that

value, x = 1.2. Do you measure the value of,

for example, the oxygen peak of Si02 and then measure the ratio of oxygen to some suboxide?

I did not quite follow that.

Strausser : We convinced ourselves that a particular dioxide film really was Si02.

Then we measured the oxygen peak height in

that film and we measured the oxygen peak

heights in the suboxide films and used those

peak heights to determine exactly what the

stoichiometry of those films were. For the

calculations of stoichiometry of Si02 films,

I have used the KLL peak rather than the LW

peak, we find that it is more reliable.

Thomas : Just one more comment. I have never

seen anything like a 7 eV chemical shift in

the KLL Auger peak. For one thing, the ESCA

2p shift for silicon is only about 4 eV,

maybe 5 eV, depending on the author you read.

Can you explain why you see 7 eV?

Johannessen : Well, we do see a 7 eV shift in

the main Si]<LL peak, between elemental silicon and silicon dioxide, and we see both

peaks as a doublet in the suboxide. I can refer you to a recent paper by C. D. Wagner.

A simple explanation is that the Auger shift

is approximately twice that of the shift in

the L2 } 3 level, which is around 2.5 eV.

Thomas : No, why should it be? All the

levels are changing equally. You are talking

about KL2 J 3L2

) 3, why should it be double?

Johannessen : Our value of a 7 eV Auger shift

in an unsaturated silicon oxide agrees well

with similar results obtained by Wagner in

oxidized silicon. In response to your question, I implied that a part of the shift was

due to different chemical shifts of the K and

L levels, although the problem is more complex

than that.

Thomas : I have to disagree with you because

I have never seen anything like that doublet

structure. Also, it is an unfair method of

looking at the stoichiometry because when

you sputter and come to the interface you may

not be removing material uniformly at all.

What you may be seeing is really an Si02

chunk sitting here on the silicon so you may

be getting a silicon signal and a surface

charged Si02 chunk more than anything else.

I have never seen a doublet like that of

equal intensity. You can get the 7 eV shifts

if you have a very large surface charge.

Johannessen : You will see a similar thing if

you look at the natural oxide on silicon.

136
Strausser : I think the important thing is

that this analyzer has enough resolution to resolve that peak and you have to consider

what instrument you are using.

Erickson : In doing ESCA spectroscopy we

quite frequently find that Auger shifts are approximately double the corresponding

shifts one finds in the ESCA spectrum. In fact, for the case of copper oxide, there is no energy difference between metallic copper

and a copper in cuprous oxide, but there are tremendous differences, 3 or 4 eV, in the

Auger binding energies for the two states of

copper. So it is not unexpected to find

larger shifts in the Auger spectra than in

the ESCA spectra. The reason is partly that

in the ESCA spectra you start out with, say

a neutral atom and the final state is a unipositive ion, whereas in the Auger case, you

start out with a unipositive ion and you end

up with a doubly positive ion. You have to consider the energy difference between the

initial state and the final state. There is

a concept known as extra-atomic or atomic

relaxation which can account for these differences in binding energy, or kinetic energy

in the case or Auger electron.

Thomas ; The Auger shifts do not have to be

double the ESCA shifts.

Erickson : It does not have to be double.

In fact, in the cuprous oxide case there is no shift at all in the ESCA binding energy

but there is a 3 or 4 eV shift in the Auger

energy. So there are some mysterious things

which are not completely understood theoretically, but they are certainly experimentally

real. I would agree that the observed shifts

in the KLL peaks for silicon and silicon

dioxide are entirely reasonable. I have seen

shifts with that order of magnitude myself.

A question addressed to either one of the

speakers: In ESCA spectroscopy, one normally

determines a direct N(E) curve and I am always a little suspicious of derivative type

curves for the following reason. If you look

at the ESCA spectrum of silicon 2p for a thermally annealed clean silicon surface and

compare that with an argon bombarded silicon

surface, one finds out that the areas of two peaks are approximately the same. However,

the silicon 2p line for the argon implanted

specimen is broadened by a factor of about 2 or 3. On the other hand, the silicon 2p line

for the dioxide is, to start with, considerably broader than it is for pure silicon so the effect of the 'residual argon in the lattice after bombardment does not have as great an effect on the peak height or on its

width. I am wondering whether or not you

see effects in the Auger spectrum due to residual argon in the lattice, and if you

find any differences in the peak heights depending on whether or not you annealed the

sample after argon bombardment.

Johannessen : We have not done any annealing

of our samples after ion sputtering, and we have not seen any significant effects on either the oxygen or silicon spectra due to argon ion implantation. The only observable

change due to ion sputtering is the removal

of surface contaminants, such as water vapor

and carbon. These contaminants tend to affect the appearance of, for instance, the

LW spectra of silicon nitride. We do not

see any effect on the Si^LL spectra. This

may be due to the larger escape depth, and

hence, the larger sample volume being probed.

I have here a viewgraph (Fig. 6) showing

some recent data on the escape depths of

electrons from silicon and silicon oxide.

The low energy data are taken from a paper

by Lindau and Spicer, and the data on the

high energy side are taken from a recent paper by Flitsch and Raider. If you look at the 1 KeV region, you will see that the es- cape depth for the Si^ix electrons is on the

order of 30 A. The oxygen KLL electrons at roughly 500 eV escape from a depth on the

order of 10 to 15 A. There is a factor of

at least 2 in the escape depth between the

oxygen and silicon KLL transitions.

In the depth profile of the Si02~Si interface we "saw" the effect of the interface on the

Si-KLL transition roughly 30 A before any

changes appeared in the oxygen peak. Therefore, the effect of different escape depths

is born out in our experimental data, and

this also shows that the ion beam does not disturb the sample deeper than some 10 - 15 A.

Morabito : I would also like to comment that

7 eV is not an unreasonable Auger shift.

Secondly, if the shift is due to a surface

charge, then you should also see a shift in

the oxygen peak and this was not seen.

Strausser : If the shift was due to a surface

charge, all the peaks should shift the same

amount and we really did not see that at all,

and we did not see a double peak at the

oxygen position. We always see the oxygen

peak at the same energy in the oxide film.

Harrington : For an escape depth, let us say,

of 20 to 30 A, is not this just the width

over which you are talking about seeing this

transition, and would that not equally well

explain the transition rather than non- stoichiometry?

137
Strausser : That is the escape depth of the

silicon KLL peak. We are talking about the

fall off of the oxygen KLL peak, and the Oj^ll

escape depth is much smaller. Johannessen

was showing the change in the shape of the

^i-KLL transitions through the interface

region, which indicates not only a change in

stoichiometry but also a change in chemical

state of the Si atoms.

Johannessen : We define the metallurgical

interface as the fall off in the oxygen KLL

signal. What we have seen is that the oxygen

atoms apparently always exist in only one state in silicon oxide; they are always

bonded to two silicon atoms as an Si20 quasimolecule

Harrington : But the changes in the silicon

peak were what you are showing for the silicon KLL and that is just the region you are talking about here in thickness, and so you

are interpreting the effect of the penetration or escape depth as stoichiometry changes.

Johannessen : It is true that we average over a depth of the order of the electron

escape depth. In the case of the Si^LL

Auger transition, it sees a much wider interface than that defined by the fall off of

the oxygen peak. It is difficult to exactly

determine the stoichiometry in the interface

region because we deal with a differential

spectrum.

Harrington : Take the situation, where,

ideally, there is just an abrupt interface,

and because you have a penetration depth

(escape depth) you are going to see the silicon and oxygen peaks of the oxide gradually

shift, even though you have an abrupt

interface

Johannessen: That is correct and it would — — o be on the order of 10 A, referred to the

oxygen KLL peak.

Harrington : That would be on the order of

20 to 30 A according to your figures.

Johannessen : Yes, using the silicon KLL

peak, but we are referring to the oxygen KLL

electrons which have an escape depth around

10 A. As a matter of fact, we have seen narrower interface regions in vacuum deposited SiOx , where one would expect an abrupt interface due to the condensation

process on the cold substrate. So we are

able to see differences for different deposition processes.

Morabito : Consider the particular example

of a thin anodic tantalum oxide film on top

of tantalum, and I have done this. If you

now monitor all the tantalum transitions, the

higher energy tantalum transition come

first, you see it coming into the interface

before the lower energy transition. In addition to that, if you have a change in the

Auger yield you are going to have a corre- sponding change in intensity and that could

be in either direction, up or down.

Johannessen : I agree we see a broader interface in the silicon KLL than in the oxygen

KLL.

DiStef ano : After unfolding the effects of

the escape depth, allowing you to see the

silicon before you get there because some electrons are escaping through the thin layer

of Si02, and after unfolding the spatial

variations in the sputter rate, how' much of

that 35 A estimated interface width is left?

How much of that graded junction is real?

Strausser : I think that the last effect is

very, very small. We have looked at these

craters under SEM's and have seen very, very

smooth craters. Secondly, our electron probe

beam is only 5 microns wide, and so it is not sampling a big enough area to have large

variations in sputter yield across it.

DiStefano : Let me paraphrase, after sputtering 1000 A of Si02, how much variation in

thickness would you have? If you have 1%

sputtering variation across that area, you

would have 10 A difference in the interface

position. So it looks like you are very

sensitive to variation in the sputter yield

across the area.

Johannessen : We have been worried about

things like that, and we have seen surface

roughening effects due to the ion beam at very low sputtering rates, less than 10 A/min.

The present data are obtained at approximately

80 A/min using a rastered low energy ion beam.

No surface irregularities were visible even

if we used the TEM and replica technique. We

believe that the provisions I just mentioned

are essential for this type of work.

Now taking all of these effects, escape

depths, ion beam distortion, etc., into ac- count, our estimate of the interface width is on the order of 20-30 A.

138
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

SURFACE COMPOSITIONAL CHANGES WITH ELECTRON

BOMBARDMENT OBSERVED BY AES

Simon Thomas

Semiconductor Analytical Laboratory

Motorola Semiconductor Products Division

Phoenix, Arizona 85008

ABSTRACT

Auger Electron Spectroscopy (AES) is generally considered to be a 'non-destructive' surface analytical technique. However, there are many

instances where the incident electron beam can alter the surface compo- sition of the solid being analyzed. The commonly observed effects of the electron bombardment are the surface decomposition, desorption, out- diffusion, surface migration and co-adsorption. Some typical examples

of these effects, often encountered in the AES analyses of semiconductor

devices are given.

Auger Electron Spectroscopy (AES) has emerged

as a very powerful and popular surface ana- lytical technique for semiconductor device

fabrication and failure analyses. The elemental specificity, sensitivity, depth of analysis and the spatial resolution have con- tributed to its popularity as an analytical

technique for the characterization of a 'practical surface'. In addition to the elemental specificity, the chemical effects often observed in the Auger spectra provide

further insight into the surface composition.

Noble gas ion sputtering is generally coupled

with AES as a means of removing the adsorbed

species from a typical processed surface or for obtaining the depth composition profiles.

Generally, AES is considered to be a non- destructive surface technique. It is assumed

that no significant surface compositional

changes are induced by the incident electron

beam which is typically of energy 1-5 keV and

beam currents of 1-50 uA. However, there are many exceptions to the non-destructive nature

of this technique. The intent of this brief

report is to draw the attention of the numerous Auger spectroscopists , particularly in

the semiconductor industry, to the typical

surface compositional changes which may occur

as a result of prolonged electron bombardment.

The beam effects do not necessarily undermine

the immense capabilities of AES as a surface

analytical technique. Nevertheless one has

to be aware of these effects and try to minimize them, if present. If any deleterious

effects are present they should be reckoned

with in any quantitative analysis. The most

commonly observed effects of the electron

beam bombardment on the surface composition,

are the surface decomposition, desorption,

outdif fusion , surface migration and co- adsorption. Some examples of these are given

below.

SURFACE DECOMPOSITION

In the surface analysis of silicon devices,

the most commonly observed adverse effect is the decomposition of silicon dioxide. With

the prolonged electron bombardment of Si02, one observes the appearance of the Si-L2 3VV Auger peak at ^92 eV which is characteristic

of elemental silicon. Simultaneously, there

is a reduction in the 75 eV Auger peak characteristic of silicon in Si02 and also a re- duction in the oxygen Auger peak height.

The various parameters affecting the decomposition of Si02 have been reported in the literature and will not be discussed here.

There are many other examples of the decomposition of oxides by electron bombardment. An example is given in Figure la where

phosphorus pentoxide present on an aluminum

substrate is reduced to mostly elemental

phosphorus in a matter of minutes (Ep = 3 keV;

Ip = 12 pA) . It is interesting to note the large chemical effects present in the Auger

features of P in P2O5. A large chemical

shift of M.2 eV and a pronounced change in the peak shape are observed from P in P2O5.

These features resemble the chemical effects

from Si and Si02. The decomposition of P2O5

with electron bombardment appears to be matrix dependent as evident from Figure l.b.

In this case P2O5 is present as an impurity

on an oxidized GaAs substrate. No changes

in the Auger features of P2O5 are observed

with prolonged bombardment. Oxides in which

decompositional effects have been observed

in our laboratory include B2O3, Ge02, hydrous

oxides of Al, Ti02 and some cases of bulk

AI2O3. This list is by no means exhaustive.

DESORPTION

Another consequence of the electron bombardment

139
is the desorption of selected surface species.

Electron stimulated desorption (ESD) has been

studied extensively by many authors and com- prehensively reviewed by Madey and Yates (2).

Typical cross sections for the desorption of

adsorbed species are in the range 10~16 - 10~22 cm^. The most conspicuous desorption

effect is observed in the Auger analysis of

adsorbed halogens with desorption cross sections of VLO- -^ cm^. In Figure 2 is shown

the variation of the Auger signal from fluorine adsorbed on an aluminum substrate, as a function of electron bombardment time (E

p

=

3 keV, Ip = 5 uA). The beam current of 5 yA

used in this study is usually considered a moderately low value for typical AES analysis.

However, one can observe the rapid decrease

in the Auger signal, which makes the absolute

peak height measurement somewhat difficult.

Assuming the surface concentration to be uniform one may translate the area of analysis

and obtain a fairly accurate estimate of the

peak height. In addition to the halogens,

the alkali metals show similar desorption effects with a strong dependence on the matrix.

The total desorption cross sections for the

alkali metals appear to be much smaller than

those for the halogens.

OUTDIFFUSION

Semiconductor devices are commonly passivated

with silicon dioxide doped with phosphorus.

In the surface compositional analysis of such

phosphosilicate glass (PSG) using AES, one of

the commonly observed phenomenon is the in- crease in the phosphorus Auger signal with

electron bombardment time. This is due to the outdiffusion of phosphorus from the bulk

to the surface. In all the PSG surfaces typically used in the device fabrication, the

outdiffusion of P has been observed. The

rate of P outdiffusion appears to be dependent

on the PSG deposition temperature and the

stoichiometry , besides the incident electron

beam parameters. Although the mechanism of P outdiffusion is not well understood it is thought that the electric field established

at the surface of the PSG with electron bombardment is the driving force. The phenomenon of the P outdiffusion has to be borne in

mind while attempting to obtain Auger calibration curves for P in PSG. In addition to the P outdiffusion, the surface decomposition

also is observed in the Auger analysis of P in PSG. Phosphorus appears to be present at the surface of the PSG as ?2®5> although this

could not be conclusively proved for the bulk

using the Auger peak features. P2°5 present

at the surface of the PSG is converted to elemental phosphorus with electron bombardment .

Chou et.al. have observed similar outdiffusion of chlorine from SiC>2 grown in chlorine or HC1 and the subsequent desorption

from the surface.

SURFACE MIGRATION

Mobile ions, particularly alkali ions, can migrate along the surface under the influence

of the electron beam. However, this is not a serious impediment in most typical Auger

analyses. Desorption of the alkali ions is a more serious problem than the surface

migration.

CO-ADSORPTION

It is generally observed that the electron

beam assists in the adsorption of oxygen and

carbon from the residual vacuum in the chamber. Although its presence is duly recognized, co-adsorption is the least problematic effect as far as the Auger analysis of

'practical surfaces' encountered in silicon

device technology is concerned. Under typical vacuum of <10-

^ Pa, routinely obtained

in an AES chamber, no significant co-adsorption occurs in the time required to plot an Auger spectrum. The problem, of course, can

be aggravated by the high residual pressures

of reactive gases such as H2O, CO, CO2 or CH4

present in the chamber.

In many cases one may observe one or more of

the above beam effects. If such effects are present , proper care must be taken in minimizing them and in interpreting the Auger

data. The need for caution is acute in the

simultaneous ion sputtering and Auger analysis when the ion beam and the electron beam

are bombarding the surface continuously.

Notwithstanding the likely surface compositional changes which may occur in selected

cases, AES is one of the most valuable surface analysis techniques for semiconductor

processing and device failure analyses.

REFERENCES

1. S. Thomas, J. Appl. Phys . , 45, 161 (1974).

2. T. E. Madey, and J. T. Yates, J. Vac. Sci

and Technol. 8, 525 (1971).

3. N. J. Chou, C. M. Osburn, Y. T. van der

Meulen, and R. Hammer, Appl. Phys. Lett.

22, 380 (1973).

140
0 50 100 150 200 0 50 100 150 200 ENERGY ev ENERGY eV

Figure 1. Effect of electron bombardment on the Auger spectrum from P^O^

.

E = 3 keV, I = 12 yA. (a) ^2^5 P resent on an Al substrate is decomposed

tB mostly elemental form. (b) ^2^5 P resent on oxidized Ga As. No decompositional effect observed.

TIME OF ELECTRON BOMBAROMENT MINUTES

Figure 2. Progressive decay of the Auger signal from

fluorine adsorbed on an Al substrate, as a function of

electron bombardment time. E = 3 keV, I = 5 yA.

141

NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg , Maryland, April 23-24, 1975 (Issued March 1976)

COMBINED SCANNING ELECTRON MICROSCOPY-AUGER

SPECTROSCOPY FOR MICRO-SPOT SURFACE AND IN-DEPTH ANALYSIS OF SILICON AND TRANSISTOR METALLIZATIONS

A. Christou, W. Weisenberger and H. M. Day

Naval Research Laboratory

Washington, D. C. 20375

ABSTRACT

Auger Spectroscopy with a focused electron beam in the scanning electron

microscope (SEM) has provided an extremely powerful diagnostic technique

for silicon surfaces and transistor metallization systems. The focused

electron beam in a SEM allows one to determine the surface topography of the specimen which can be correlated with the Auger spectra of selected

areas. In addition, a scanning Auger image may be obtained by using the amplitudes of a selected Auger peak for intensity modulation of the CRT. The normal modes for micro-spot AES investigation in the SEM include

spot analysis, line scan and Auger electron imaging. Secondary electron

image resolutions of less than 2000 A have been correlated with Auger

spectra of submicron regions. Surface features associated with gold and

aluminum metallizations of microwave power transistors have been analyzed in detail using the AES-SEM technique. In addition the submicron

Auger spatial resolution capability has been applied to the identification of contaminants on silicon surfaces.

INTRODUCTION

The analysis of failed regions on fine geometry devices has previously been extremely

difficult. Energy dispersive x-ray analysis

includes a reaction area about one micron

(at 15 kV) below the surface. Auger analysis only examines about a 10 A depth, however most systems use an analysis area which

has a 50 urn diameter. Even the scanning

Auger microprobe (SAM) uses a 5 ym beam and

has no corresponding high spatial resolution

SEM topography information. Microwave power

transistors are characterized by cells on

the order of 50 to 100 pm with line spacing

on the order of 1 urn.

The utilization of electrostatic lens and

the geometric configuration 1

' 2 of present

day Auger detectors has made the attainment

of submicron size electron beam a formidable

problem. In the present investigation, the

focused electron beam in the scanning electron microscope (SEM) is used as the primary

beam for Auger spectroscopy 3 » 4 > 5 thereby

improving the AES spatial resolution to

.2 urn. The microspot technique therefore

allows the determination of Auger spectra of selected areas in addition to the surface

topography (SEM mode) of a failed region.

Scanning Auger images may also be obtained

by using the amplitude of the Auger peaks

for intensity modulation of the cathode ray

tube. The spatial resolution of . 2 pm to

1

2 um for Auger spectroscopy obtained by this

technique is ideally suited for the analysis

of (1) localized interdif fusion regions in

metallizations, (2) adhesion problems and

(3) process associated contamination. The

instrument requirements for high resolution

Auger spectroscopy of solid state devices in

a commercial scanning electron microscope

will be reviewed in order to understand the usefulness and limitations of the micro-spot

SEM-AES technique.

General Requirements for AES-SEM Instruments

The requirements for AES in a scanning electron microscope are the vacuum chamber, the energy analyzer and the electron optical

column. The vacuum system used to achieve

a clean ultra-high vacuum in our SEM is an electro-ion pump with a mechanical carbon vane pump and cryo-sorption pumping for an oil free 2 stage roughing system 5

. The

electro-ion pump is a four-cell triode type

electrostatic getter-ion pump with a speed

of 1600 1/sec for dry nitrogen. The pump is closely connected to the specimen chamber of the SEM by a 6 inch ultra high vacuum gate

valve. Figure 1 shows a schematic of the entire system. The SEM is an ETEC autoscan

specially made for this application6

. The

entire system is of stainless steel construction. The only non-stainless steel part ex- posed to vacuum is the viton seal at the gate

valve. The ports contain the Auger detector,
ion etching capability, secondary electron

detector, leak valve and one sight port with

an optical microscope. Pressures of 3 x 10-9 torr in the specimen chamber have been

obtained after a mild bakeout atl35°C.

Under normal operational conditions, the

system is operated at pressures between

5 x 10~ 9 torr and 1 x 10-8 torr. A residual

gas analysis of the system was attempted in

order to determine the presence of contaminants. A quadrupole mass spectrometer was installed in the approximate vicinity of the

specimen stage. The constituents due to water had partial pressures of the order of

1 x 10-8 torr, while nitrogen partial pressure was below 3 x 10-9 torr.

The Electron Optics and Auger Detector

The electron optical column is the ETEC

multi-lens system housed in a modular prealigned column. The energy source was a triode electron gun with a high emission

tungsten filament. The spherical aberration

coefficient of the final lens was 3 cm for a working distance of 8 mm. Figure 2 shows

the relationship between probe current and

probe diameter for a fixed working distance

of 19 mm. The curved lines are for the 400

and 500 um aperture. The straight line is

for the optimum aperture obtained by varying

aperture diameter. Since most of the AES

investigations were accomplished with a probe current of 10-7 A to 10-8 A the opti- mum sgatial resolution possible is between

2000 A and 1 um at 20 kV. However, one must

also take into account such parameters as escape depth, atomic number, surface area

and beam energy when determining spatial

resolution. A second factor other than beam

current which must be considered in obtaining the required resolution is recording

time. The signal to noise (S/N) ratio can be increased by utilizing longer recording

times. The recording time of an Auger spectrum can be increased to reduce the beam

current while maintaining the same S/N

ratio.

The Auger detector used is the cylindrical

mirror type 5 designed to minimize the work--

ing distance of the scanning electron microscope to 19 mm. The analyzer energy reso- lution was measured to be .5% at 3 kV beam

voltage with 10% transmission. Using a sputter ion gun and a multiplexing unit for

monitoring peak to peak amplitudes, up to six Auger peaks can be monitored while the

specimen surface is etched.

AES-SEM Modes and Analyses

a. Spot Analysis . The normal analytical

modes include spot analysis, line scan and

1

scanning Auger imaging. Each of these analyses can be combined with sputter etching in

order to obtain an in-depth analysis. The

Auger image can be recorded for a specified

array of points or for a specified number of

line scans. The amplitudes of the Auger

peaks are then recorded or displayed as intensity or amplitude modulation of the recording cathode ray tube.

The spot analysis will now be applied to the

failure analysis of a microwave power tran- sistor in order to illustrate the unique

application of this instrument as a diagnostic tool. Figure 4 illustrates a region

typical of a failed microwave power transistor from pulsed RF accelerated life tests,

showing the presence of a film material between the metallizations. Micro-spot AES-SEM

analysis was applied to the reacted regions

between the metallizations in order to identify the reaction products (oxide overcoat

and aluminum reaction)

Application of the energy dispersive x-ray

analysis technique (EDXA) on the area of

interest was attempted before the initial

AES analysis. A large 1740 eV Si peak was observed (figure 3) along with the 1490 eV aluminum line. The silicon detected could

have been present in the reaction region or

it may be from the substrate silicon. In addition, the EDXA technique cannot differentiate between aluminum in A120 3 and free

aluminum or between free silicon and silicon

in Si0 2

. The AES spectrum with the SEM electron beam situated on the reacted region is

shown in figure 4. The position of the

aluminum KLL Auger transition is indicated

at 1407 eV. Also shown is an aluminum oxide

peak at 57 eV. Therefore, the surface of

the reacted zone consists essentially of

A1 20 3 , oxygen and carbon (from handling and

testing). Free silicon, up to the sensitivity limitations of the instrument which is

1.0 percent per monolayer was not detected.

The lack of detection of silicon is shown

by the absence of the 92 eV LMM and 1619 eV transitions. Auger analysis of the unreacted

emitter metallization shows the position of

the Al (KLL) transition at 1396 eV as expected

for pure aluminum''. In addition, the position of the Al LMM transition was located at

69 eV instead of the 57 eV peak found for the

aluminum reacted zone. The peak shift and

change in shape of the KLL peak for aluminum

clearly indicated that the film region of

figure 4 was the result of the formation of

A1 20 3

.

b. Auger Imaging. A problem related to semiconductor device reliability is the

thermal dissolution of silicon into aluminum

enhanced by microcracks or flaws in the
aluminum thin film. At 450°C the diffusion

process is extremely rapid and the solubility requirements are satisfied within a few

minutes. However, microcracks in the film

accelerate the diffusion process and often

result in the migration of silicon to the

top surface. Figure 5 shows the example of

silicon diffusing through aluminum after a one hour anneal at 450°C. The surface of

the interdif fused aluminum film was scanned

with the analyzer set on the 92 keV KLL peak

maximum and with a frame time of approximately .35 h. The signal was amplitude modulated in order to obtain the Auger image

shown in figure 5. Figure 5 indicates resolution of submicron details in the scanning

Auger mode at a working distance of 23 mm. A feature of interest on the silicon stripe

can now be analyzed by moving the specimen

to the center of the display stopping the

primary beam scan and measuring the Auger

spectrum in the usual manner.

c. Auger Line Scan Analysis . Analyses

given thus far have reviewed the spot and

imaging techniques . A third Auger technique

which can be accomplished successfully in an

SEM is Auger line scan analysis, for example,

of an extruded silicon particle oh the surface of a microwave transistor. The image

of figure 6 was obtained with the analyzer

set on the 92 eV Si Auger peak and shows the

results of a single scan of the electron

beam across the particle. The position of

the particle correlates with the maximum of

the Si line scan indicating that the particle is essentially silicon. Also shown in

figure 6 is the spot Auger analysis of the

particle identifying the 92 eV silicon peak, > the 270 eV carbon peak and the 1619 KLL silicon peak.

d. Micro-spot Sputter Profiles . The AESSEM technique when combined with electronic

multiplexing 2 ' 5 has been applied to obtain

micro-spot Auger sputter profiles of reacted

thin film structures. Figure 7 shows localized gold-silicon recrystallized regions of

reacted gold-tungsten films on a silicon

substrate. Analysis of the globular regions

formed at 450°C were found to have a 3:2

surface composition of Au-Si as shown in

figure 7a. The interdif fused region of AuSi was approximately 1500 A thick (assuming

a sputtering rate of 60 A/min. The migration

of tungsten into the Au-Si structures was

observed during the initial 25 minutes of

sputtering. After sputtering for 100 minutes, the gold was completely removed, with

silicon remaining probably as a WSi2 structure 8

. The areas adjacent to the globular

regions were found to be gold with trace amounts of silicon out-diffused from the

145

silicon substrate (figure 7b). The silicon

distribution was found to be uniform throughout the gold and tungsten films with the surface composition varying between 2 and 10

percent

.

In summary, the general approaches for high

spatial resolution Auger spectroscopy in an SEM have been reviewed. A system consisting

of a commercial SEM with an ultra-high vacuum

capability, integrated with an Auger electron

spectrometer and sputter capability has been

demonstrated. A secondary electron image

resolution of approximately 2000 A and Auger

spectra of features less than 1 u has been

attained. Data taken on aluminum contacts of

semiconductor devices has been reviewed in

order to examine the capabilities of the instrument .

REFERENCES

1. Chang, C. C. , Surface Science 25 , 53

(1971).

2. Weber, R. L., Research Devetopment 2^3 (10), 22 (1972).

3. MacDonald, N. C. , Auger Electron Spectroscopy for Scanning Electron Microscopy, Scanning Electron Microscopy 1971,

0. Johari and I. Corvin, Eds., pp. 89-

96, (IIT Research Institute, Chicago,

Illinois, 1971).

4. Brandis, E. K., and Hoover, R. A., Proc.

7th Natl. Conf. on Electron Probe

Analysis, 1972, Paper 84.

5. Christou, A., Auger Spectroscopy of

Solid Surfaces in a Dry Pumped High

Resolution SEM, Scanning Electron Microscopy 1975, 0. Johari and I. Corvin, Eds.,

pp. 149-156 (IIT Research Institute,

Chicago, Illinois, 1975).

6. ETEC Autoscan SEM, ETEC Corporation,

Hayward, California 94545.

7. Palmberg, P. W. , Handbook of Auger

Spectra (Physical Electronics Industries.

Eden Prairie, Minnesota, 1972).

8. Sinha, A. K. , Thin Solid Films 2_0, 115

(1974).
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146
PROBE DIAMETER

Figure 2. The relationship between probe current and

probe diameter for AES-SEM system.

Figure 3. EDXA analysis of a failed microwave power transistor showing the detection

of both a large 1740 eV Si peak along with

the 1490 eV aluminum peak. Also shown is a schematic comparing the X-ray and Auger

electron analysis regions.

147

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149
DISCUSSION OF THE PAPER

Harrington ; Although you gain in spatial

resolution what you do feel you loose in

sensitivity with the Auger technique at the

small beam diameters?

Christou : I do not have a quantitative

answer for that, but you certainly degrade

the signal to noise ratio significantly.

The best way would be to computer interface

your system so that you could observe peaks

hidden in the noise.

Madey : I would, just like to issue a word of

caution about beam induced damage and problems of the sort that you have described

here. I think it is particularly acute when

you have such a small beam and a high current. For the situation you described with

10~9 amps and a 0.1 micron beam diameter you

have a current density that is of the order

of 10 amps per square centimeter. Now for

electron stimulated desorption and decomposition, typical cross sections are of the

order of 10~20 to 10~22 square centimeters

for decomposition of surface oxides, desorption of oxygen and other species from

surfaces. Under these conditions the average lifetime of the oxygen on the surface

could be of the order of 1 to 100 seconds

which is of the same order of magnitude as you need to get a workable signal to noise

ratio. So it is quite possible then you can

get appreciable beam decomposition and damage, particularly in compounds, insulators

and semiconductors.

Christou : Yes, I agree and it is for that reason that we have been mainly applying

this technique to interdif fusion phenomena

and thin films.

150
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg , Maryland, April 23-24, 1975 (Issued March 1976)

APPLICATIONS OF X-RAY PHOTOELECTRON SPECTROSCOPY

(ESCA) TO MIS DEVICES*

Frank J. Grunthaner

Jet Propulsion Laboratory

California Institute of Technology

Pasadena, California 91103

Summary

X-ray photoelectron spectroscopy (XPS, or ESCA) is described as both an analytical and

experimental tool for the characterization

of metal-insulator-semiconductor (MIS

devices. A brief description is given of general problems which lend themselves to the application of this technique. Key ex- perimental methods in XPS are discussed.

Several illustrative applications are con- sidered including characterization of the Si/Si02 interface, surface corrosion in

hermetically sealed packages, mobile ion instabilities, organic surface contamination,

etchant residues and heavy-metal impurities. New developments in XPS are discussed and

some limitations of the technique are assessed .

Introduction

resistance. In each of these modes, several

failure mechanisms can be detailed which are dependent on the chemistry of the oxide, that

of the defects in the oxide, and that of residual species determined by the process. The study and characterization of the process

requires chemical and physical information

resulting from the probing of well-defined

regions of the device. The study of the surfaces of devices can yield such information.

X-Ray Photoelectron Spectroscopy

XPS has seen considerable development since

the first monograph on the phenomenon appeared

in 1967. 1 Work in this field has also been

catalogued under the more general description

of electron spectroscopy for chemical analysis (ESCA) . A number of recent reviews have

appeared 2-5 and can be consulted for experimental detail.

In the last three to four years, major ad- vances have been made in the technologies of surface spectroscopy. These spectroscopies

perturb the surface of a sample with

photons, ions or electrons, and monitor the

relaxation processes which results. This

yields a variety of data concerning the physical parameters of the surface. Several

of these methods give estimates of the elemental composition of samples. One of these

techniques, XPS, also gives significant

chemical information about the surface.

The production of an MIS device is basically

an exercise in chemical processing technology. The fabrication of high reliability

devices is dependent on achieving an understanding of the interaction of processing

chemistry with device physics. In CMOS technology, for example, several dominant

failure modes can be identified including:

oxide breakdown, threshold voltage shifts,

resistive and open connections, excess

leakage and shorts, and changes in channel

Figure 1 describes the physical process. A

soft x-ray source is used to irradiate the

surface of a sample with photons of a characteristic energy (e.g., AS, Koij 2 )- Tne absorption of the photon can lead to two excitation cases: 1) x-ray absorption where

an electron is excited from a core-level to the Fermi level, and 2) electron emission in

which the electron is excited from the bound

core state into the vacuum continuum. Either

of these processes result in the excitation

of the absorbing atom. The system can undergo relaxation through electron or photon

emission or through more complicated radiationless processes. The photon emission process gives rise to an x-ray quantum (x-ray

fluorescence) . The electron emission event

is termed auto-ionization and gives rise to an Auger electron. Auger electron spectroscopy (AES) typically uses high-energy electrons to create the core-level holes required

to stimulate this electron emission.

In the case of photon absorption followed by

direct electron emission, the energy of the

ejected electron is given by:

hv = E + EF

(k) + qV + W (1) kin s spec

This paper presents the results of one phase of research performed at the Jet Pro- pulsion Laboratory, sponsored by the National Aeronautics and Space Administration

under Contract NAS 7-100. where hv is the photon energy, E the

151
kinetic energy of the photoelectron, V9 the

sample surface potential, Wg pec the work

function of the spectrometer, and E^(k) is

the total final energy of the absorber after

the ejection of an electron from the k^

orbital as a photoelectron, referenced to the Fermi level.

Since Wg p ec has a constant value for a particular spectrometer, Vg has a constant

value under controlled conditions, and hv is

well known, the kinetic energy of the photoelectron is directly related to the total

energy with which it was bound in the adsorber. This relationship is generally

written as

E DC = hv - E (2) BE kin

where Egg is the binding energy of the

photoelectron.

E (k) is a function of the orbital as well

as the element, and consequently a number of

photoelectron lines can be resolved for each

element. This is illustrated in figure 2,

where the ordinate is given in units of

atomic number and the abscissa is given in

binding energy (eV) of the appropriate

photoelectron. Each of the elements has an XPS line well separated from the others, and

peak positions can be readily determined

with precisions of better than 0.1 eV. The

corresponding plot of Auger electron (AES)

lines plotted in units of apparent binding

energy (eV) is given in figure 3. At high

kinetic energies (low binding energies)

these lines are also quite distinct. In practice, both AES and XPS lines are used to characterize the presence of an element.

The intensity of a photoelectron peak is

determined by the photon flux, the effective

sample area, the spectrometer acceptance

angle, the photoelectric cross section, the

mean-free-path of the photoelectron in the

sample, and the number density of emitting

atoms. 6 Individual photoelectron lines have

been empirically scaled to a common element

(fluorine) and the result plotted in figure

4. Here, relative intensity is plotted vs atomic number for several photoelectron

lines.

XPS spectra are formulated through determination of the kinetic energy distribution of

the photoelectrons emitted from the surface

of the sample. The experiment is surfacesensitive because the measured photoelectron

line represents those electrons escaping the

solid without undergoing inelastic scattering. This escape depth determines the

probing region of the experiment and is a

strongly varying function of the electron

kinetic energy and the density of delocalized

valence electrons in the material. Metals

have much shorter photoelectron escape depths

than do insulators. The kinetic energy dependence of the escape depth of the photoelectrons is diagrammed in figure 5. The

position of the curve along the length axis

is determined by the material characteristics

as discussed above.

The surface or bulk sensitivity of XPS and

AES is largely determined by the kinetic

energy of the escaping electron and not by

the method of excitation.

The analytical aspects of the XPS spectrum

are illustrated in figure 6 which shows a

0 - 1000 eV binding energy scan of the kinetic energy distribution. These spectra

correspond to a 63 A film of Si02 , thermally

grown on Si. The ordinate in this spectrum

is the number of electrons expressed in

arbitrary units. On the abscissa is plotted

the binding energy (eV) increasing from

right to left. The lines at 100 and 150 eV are the silicon 2p and 2s lines, respectively.

The line at 285 eV corresponds to the carbon

Is level, which the peak at 530 eV is 01s

level. The peak to the far left at 985 eV arises from the KLL oxygen Auger transition.

The photoelectron lines show widths of 0.58 - 1.45 eV, while Auger electron lines range

from 2 to 27 eV wide.

In addition to the analytical aspects of XPS

spectra, the line positions are shifted by

the chemical state of the element. This

chemical shift parameter is the most important information available in the experiment.

The position of XPS lines was found to have

an "error" of i 5 eV in early experiments. 1 Upon closer examination, line positions were

found to be quite reproducible for a particular compound, but varied for different com- pounds of the same element. This chemical

shift has a magnitude of ± 5 eV over the

range of chemical states observed for a given element. Recent advances in instrumentation have reduced line widths to 0.55 eV from 1.0 eV, resulting in a major increase in

chemical definition. This chemical shift is

illustrated in the insets of figure 6. The

right inset is a 10 eV expansion of the silicon 2p region. In the Si 2p spectra, two distinct lines are observed. The peak at 102.5 eV corresponds to the silicon of the

silicon dioxide while that at 99 eV is due

to the silicon substrate beneath the Si02

films

In summary, the XPS experiment gives analytic information in terms of quantitative
and qualitative analysis. Quantitative information requires calibration, and precisions in the ± 1% range are easily

achieved. In qualitative terms, sensitivities of about 1 ppm have recently been demonstrated. 7 The chemical shift is the most

important aspect of the experiment and permits assignment of the chemical state of

various surface species. The observation

depth is typical of surface spectroscopies

and ranges from 3 - 4 A to about 125 A.

Experimental Considerations

The execution of the XPS experiment is diagrammed in figure 7. An x-ray source is

required to irradiate (excite) the sample

surface. The kinetic energy distribution of" the emitted electrons is determined by means

of an electron monochromator . This sequence

requires that source, sample and analyzer be

contained in a high vacuum system.

In practice, however, this basic experiment

imposes several serious constraints on the

equipment. The most important of these are listed in table I. An ultra-high-vacuum

system is mandated by the surface nature of

the experiment. Unless sample limited pressures of 10-9 torr and less are achieved,

the surface can become recoated during analysis. The residual gases present in the

spectrometer must be as non-reactive as possible, e.g., He, A, N2, etc., with low

sticking coefficients.

Since photoelectron spectra consist of sharp

lines rising out of a significant scattered

electron background, the highest possible

signal-to-background ratio must be achieved.

The background level is determined by the

S/B ratio of the x-ray source, the surface

features of the sample, the presence of

secondary electrons in the sample chamber,

and the electron background of the analyzer/

detector. High S/B ratios in the x-ray

source require the use of soft x-ray monochromators. The use of monochromated

sources gives rise to collimated x-ray

beams', which reduce the secondary electron

flux in the sample chamber.

The analyzer, x-ray source and detector must

be treated as a system and optimized for the

maximum throughput; that is, for each photon

impinging on the sample surface, the largest

possible fraction of electrons must reach

the detector.

A sample-mounting system is required which

will permit variation in the electron acceptance angle of the analyzer in relation to the sample surface. This variation permits

depth analysis in the range 2 - 120 A without destruction of the sample. 9 Control of

secondary electrons in the sample chamber

enables a series of electrical experiments

which will be described later. Variable

temperature capability permits capture of

reaction intermediates as well as study of

"in situ" reactions, thermal degradation,

etc

.

The actual XPS experimental setup is illustrated for the JPL electron spectrometer in

figures 8-10. Figure 8 gives a simplified

diagram of the basic ESCA spectrometer. The

soft x-ray source consists of an aluminum

target bombarded with 15 keV electrons.

These photons are focused by the spherically

bent quartz-crystal monochromator on the

sample surface in a separate chamber. The

x-ray beam is collimated to a spot size of

1x5 mm. Electrons enter the energy ana- lyzer through the 4-element electron lens,

which retards and focuses the electrons.

The analyzer is a spherical capacitor operated at a fixed energy transmission. The

electrons are then directed against the

multielement detector and the data then read

into the multichannel analyzer.

A physical representation of the sample

chamber of the spectrometer is given in

figure 9. The x-ray source and analyzer

are indicated in perspective. The sample

is mounted on a sample rod which permits

variation of the electron take-off angle

seen by the analyzer. This permits the

emphasis of either surface or bulk features

in the ESCA spectrum. A low-energy electronirradiating device (flood gun) is mounted in

order to control charging effects and thus

the surface potential. It provides a diffuse

beam of electrons whose average energy can be

varied from 0 - 10 eV above the vacuum level

of the gun. Emission current can be varied

over three orders of magnitude. An electron

probe is being installed for direct measurement of the surface potential. The probe is

provided with high-resolution optics at low

kinetic energies and with capability for x-y

scanning along the surface. The sample support is insulated from ground and therefore

can be connected through a sensitive ammeter

to the ground in order to detect currents

from the electron probe.

The general surface analytical installation

is diagrammed in figure 10. The key component is the high-resolution electron spectrometer (HP-5950A) . The electron probe

discussed above is also planned to operate

as a scanning electron gun to stimulate

Auger electron spectra. A sample reaction

area is provided which is separate from the
analyzer chamber and which contains equipment for secondary ion mass spectroscopy,

electron-stimulated desorption, ion milling,

metal deposition, and gas-solid reactions.

A sample-preparation area is attached to the

inlet chamber of the spectrometer consisting

of a dry box with introduction lock. The

dry box is under 5 psi pressure of nitrogen,

which is dried over LN2 . A high-temperature

furnace is provided for controlled growth of

oxides on silicon. The furnace has an introduction lock permitting direct transfer

of samples to spectrometer without aqueous

or carbonaceous contamination. A plasmacleaning facility is interlocked to the dry

box to permit removal of trace organic contamination.

Sample Preparation

For reproducible results in XPS, considerable attention must be given to samplemounting procedures. In this specific case, where the Si02 /Si interface is the object of

study, reproducible electrical contacts must

be fabricated to the silicon substrate. The

problems of ohmic contact and that of physical mounting were solved by using lowmelting-point solders, such as InSn and

GeAu. The sample and the mounting configuration are indicated in figure 11. A gold

(Au) platen is prepared by washing with TCE,

acetone, then several rinses with high re- sistivity (20 megohm) water. The platen is

bonded to the rough underside of the oxidized silicon wafer by heating on a hot

stage in the nitrogen atmosphere of the dry

box. A gold window is placed over the top

surface of the sample. The area illuminated

by x-rays is indicated in the center of the

window zone and is somewhat less than 1 x 5 mm (0.3 x 4 mm). Since the detector of

the HP-5950A spectrometer is positionally

sensitive, the long dimension can be reduced

to approximately 1 mm, giving 0.3 mm2 as the

minimum analytical area.

During the experiments described in the

following sections, the spectrometer was operated at a pressure of 6 x lO

- ^ torr. The ambient gas consisted primarily of H2 , He with lower levels of N2 and N. H2 0 and

A were minor components, but some CO was

present

Applications

Silicon-Silicon Dioxide Interface

In CMOS technology, the fundamental device

characteristics are determined by the oxidesemiconductor interface. Our initial experiments indicated that the escape depth of

photoelectrons in Si02 was consistent with

X e = 40 A. This suggested that one could

study the elemental silicon surface when o covered by thin layers of Si0 2 up to 100 A

thick. Other experiments 9 have indicated

that reproducible pinhole-free thermal oxides

can be grown on silicon with thicknesses in

the range of 15 - 200 X.. These oxides,

grown in our laboratory, show extremely uniform thicknesses, and give excellent device

characteristics

Figure 12 reproduces the silicon 2p photoelectron spectrum of a freshly etched p-type

silicon crystal. The spectrum comprises

10 eV, and the 2p spin-orbit doublet of

silicon is clearly resolved. The spin-orbit

splitting of silicon has a magnitude of about

0.6 eV. An instrumental system resolution of

0.55 eV is indicated by this spectrum.

In order to demonstrate the chemical shift

between silicon in the elemental state and

silicon in the +4 oxidation state (Si02 ),

the spectra of figure 13 show both cases. The upper spectrum arises from the 2p doublet

of elemental silicon. This sample is identical to that of the lower spectrum except

that, in this case, the Si02 has been stripped

with H2 0/HF and a monolayer of fluoride is

present on the sample surface. The lower

spectrum is due to a wet 40 A oxide grown on silicon 100. In this spectrum, the Si 2p

line due to the substrate occurs at nearly

the same energy as in the elemental case. The peak at higher binding energy (104.3 eV)

is due to the oxidized silicon of Si02

. This

shift of about 4.0 eV is characteristic of

the 2p binding energy differences between Si0 and Si 4+

.

There is reason to believe that the line

width, line shape, and spin-orbit splitting

of Si1*"1" is similar if not identical to that

of Si°. The broader structure of the oxide

peak is indicative of a distribution of

chemically inequivalent states. Although

this distribution can be mathematically

determined from the data,*" it cannot be

directly observed through higher instrumental resolution. It is possible, however, to experimentally modify this distribution of

states by applying electrical stress to the

oxide film.

In the XPS experiment, the metal-oxide interface of an MIS device can be modelled by the

vacuum/oxide boundary. Electrical contact

to the oxide can be accomplished by means of

the vacuum level. The XPS experiment tends

to produce a positive surface charge in the case of an insulator. An x-ray photon interacts with the surface of the sample creating

154
an electron hole pair. If the electron is

emitted into the vacuum, a hole is left behind. Therefore, a positive charge can

build up in the thin surface layer interrogated in the XPS experiment 100 A) . The

steady-state surface charge reflects the

balance of the depletion current (photoemission) against the replenishment currents

(vacuum secondaries, tunneling currents,

etc.). Since the vacuum secondaries are minimized through x-ray collimation, and

low x-ray flux densities, a significant

positive surface charge can be developed in

Si02. In very thin oxides this potential

is ultimately limited by electrons tunneling

from the Si interface. 11

A corresponding negative potential can be

induced by increasing the current density

of low-energy electrons to the sample surface from the flood gun. Therefore, a bias

of either polarity can be applied dynamically in the course of the XPS experiment.

Figure 14 illustrates the effect of bias on

Si 2p spectra of a 44 A Si02 layer on silicon (100) . The positive bias spectrum is

given as the lower curve. The negative bias

spectrum (upper curve of figure 15) was re- corded while bombarding the sample surface

with electrons of an average energy of

8.6 eV above the vacuum level. A change in

line shape of the elemental line is observed,

as is a change in the relative position of

the oxide peak.

Previous work of just two to three years ago

identified the escape depth of photoelectrons in Si0 2 at 10 - 15 X (A £

). 2 Figures

15-a and 15-b give the Si 2p spectra of a 113 & film of Si02 over Si. In figure 15a,

the large difference in relative intensities

of the oxide and substrate peak are evident,

while figure 15b shows the line shape of

the substrate silicon line at higher gain

(20 X). Its binding energy of 99.06 eV is

in good agreement with that of the Si° peaks

of figures 12 and 13.

The Si1* state (Si02 ) is clearly resolved

from that of Si°. The determination of

silicon to oxygen ratio can be directly calculated by comparing the oxidized silicon

peak intensity with the intensity of the

oxygen Is peak. The magnitude of excess

silicon, and reduced silicon oxides can now

be directly measured, within the sensitivity of the experiment. For these thin

oxides, no ion milling is required to see

the interface and hence no ion-stimulated

redistributions need be considered in interpreting the data.

Migratory-Gold-Resistive-Short

The technology of metallization in integratedcircuit devices is based on either aluminum

or gold. The presence of moisture within the

device package has long been known to leak to electrolyte corrosion of aluminum. This

corrosion results in a well-documented failure mode. The chemistry of gold suggested

that metallization schemes using this noble

metal would be free from susceptibility to such electrochemical attack.

Recently, electrolytic corrosion has been

observed and described phenomenologically as the migratory-gold-resistive-short (MGRS)

mode. 12 The failure is seen as the abrupt

decrease of resistance between two or more

adjacent metal stripes. The shorting current are on the order of milliamperes . SEM

micrographs of failed parts reveal that fernlike or dendritic gold deposits have formed

a bridge between metal stripes.

The dendritic deposits appear to progress

from the negatively-biased stripe (cathode)

toward the positive one. The general form

or morphology of the growth is quite specific

and similar to that of deposits found in

electrochemical silver migration. 13 This

well-known phenomenon results from the surface transport of silver which is driven by

large electric fields. This surface transport is strongly dependent on the presence

of moisture and on the nature of the substrate surface.

The MGRS failure is observed on the surface

of integrated circuits using Ti/W/Au metallization which has been covered with a layer

of e-beam-evaporated quartz. Failure analysis of affected parts showed significant

levels of water in the purportedly hermetically sealed package.

The presence of water made possible a corrosion cell. This implied that the rate of

attrition of parts is determined by the

amount of time adequate potential is applied

across adjacent metal stripes. The major

difficulty, however, was that dendritic

growth requires as few as 10 10 atoms of gold

to bridge a stripe.

Since gold cannot be simply dissolved in

aqueous solutions, we studied the surface of

failed and reference parts to determine the

nature of the impurities residing atop the

passivating glass layer. The presence of

significant levels of halogens was observed

by means of XPS.

155
This is an example of surface corrosion

chemistry, and three points are relevant to this discussion: 1) positionally sensitive

detection, 2) limiting sensitivity, and

3) decomposition of sample system. The

first of these points is illustrated in

figures 16 and 17. Figures 16 shows a 0 - 1000 eV binding energy scan of the top of

the chip, together with the bonding pads,

wire leads, and some parts of the package

base. This is determined by the relative

intensities of the gold 4f photoelectron

line at about 85 eV and of the oxygen Is line at 530 eV. The gold line is approximately 25% more intense than the oxygen Is line. In figure 17, the detector is adjusted to scan the top of the chip, and the

oxygen/gold intensity ratio has changed by a factor of 22.

The analysis of the surface of the glassivated substrate by XPS indicates that significant quantities of the halogens (CI , Br~, I") are present together with hydroand f luorocarbons and several trace metals,

such as Na+ , K+ and Hg. Of the halogens,

Br was most abundant while CI and I had

relatively equal concentrations. The

bromine 3d region of the photoelectron spectrum is given in figure 18. The structure

arises from at least 2 different chemical

states of the element . The iodine 3d region

of the spectrum is plotted in figure 19. The

spin orbit doublet is clearly defined in this

spectrum.

Over the range of failed and unfailed parts

which have been examined by XPS, the concentration of I has varied over four orders of

magnitude. Levels of iodine as low as 1 ppm

can be readily detected. This represents a sensitivity limit about three orders of

magnitude lower than was previously observed. 2 This is related to the relatively

high photoelectric cross-section of the

iodine 3d electrons for 1.5 keV x rays, and

to the magnitude of the spin-orbit splitting.

This separation of two lines with an appropriate intensity relationship facilitates

the use of sophisticated noise-removal techniques for data manipulation.* 1*

Interestingly, many of these impurity atoms were not observed in electron-stimulated

Auger and x-ray spectra. This suggests that

for many problems, electron-stimulated desorption is a serious experimental difficulty. Monochromatized x-ray excitation

offers a minimal source of secondary electrons for such desorptions.

Mobile Ion Instability

The role of mobile ions in causing instabilities in MOS devices is well known and several reviews have appeared. 15-19 The previous work has demonstrated that sodium is

the most common source of such instabilities.

Sodium contamination in thermally grown Si02

is often introduced by the sodium in the

oxidizing or annealing ambient. 20 Most of

the sodium present in the ambient can be

introduced through the walls of the furnace

tube since this tube is usually fabricated

of silica, which permits extremely high

diffusion rates of sodium at the required

operating temperatures. Additional furnace

tube liners are often used to reduce this source of contamination. The processing

steps following high-temperature oxidation

are also potential sources of sodium contamination, since most solvents—organic and

aqueous—show significant sodium levels if

not controlled. The effective concentration

of sodium in the bulk arising from surface

contamination of the oxide is directly

determined by the maximum temperature to which the structure is exposed. Finally,

sodium contamination can be introduced

during metallization unless extreme precautions are taken.

It has been established that the rate of

removal of sodium from the Si02 /Si interface

is much faster than that from the Si02 /metal

or Si02 /vacuum interface. The mobility of

Na ions in the "bulk" of the Si02 is quite

high (at least 4 * 10~ 13 cm2 /V sec at room

temperature) . The rate-limiting step during

forward drift appears to be the "emission"

of ions from "traps" at the Si02 /metal interface. The trapping at the Si02 /Si interface

is much weaker and is strongly dependent on the conditions of oxidation and annealing.

The existence of two forms of traps for Na+

at the Si02 /Si interface has been suggested

where one of these becomes saturated at high

ionic contamination levels. This patch model

has been developed to explain variations at the metal/oxide interface as well. 21

Recently, methods have been developed to getter sodium ions in Si02 or to neutralize

Na+ at the Si/Si02 interface. Phosphosilicate glass (PSG) is widely used to trap

or getter Na+moving from the metal/Si02 boundary to the Si02 interface. Recent

experiments 22 have indicated that the

addition of HC1 in the high-temperature

oxdiation step modifies the silicon interface and permits charge neutralization of

156
Na as it approaches the Si boundary. These

methods basically involve a modification of

the chemistry of trapping defect sites within the oxide.

Unfortunately, these traps created through

process variation are highly susceptible to high-energy gamma- and electron-radiation.

In radiation fields, the sodium is again released, leading to device failure. We chose

to initiate a study of mobile ions in thin

SiC>2 by means of XPS in order to understand

the chemistry of the defect-site mobile-ion

interaction.

Conversations with device physicists indicated that there existed two fundamentally

different forms of sodium in silicon oxides.

One of these is termed positive, has a +1

charge, and moves when subjected to appropriate fields. The second form is termed

neutral sodium because it is positionally

trapped and cannot translate in response to

dc fields.

These observations, together with the "patch"

model of specific traps suggested a coordination chemistry of sodium which might lead

to well-resolved states in the XPS spectrum.

In a previous section, it was observed that

a positive or negative bias could be applied

to the surface of an insulating sample in

the course of the XPS experiment. The com- bination of applied bias and elevated tem- peratures was employed to pull the Na+ in

the oxide out to the oxide/vacuum interface,

or alternately push the Na"*" into the oxide

silicon interface. The effects on the silicon, oxygen and sodium spectra have been

catalogued and interpreted. 23 Figure 20

gives the Na Is spectrum of an as-processed

device. Electronic measurements of this

sample indicated a mobile-ion concentration

of about 10^ atoms/cm2

. This spectrum is

the result of 51 hours of accumulation time.

In figure 21, the Na Is spectrum of a different area of the silicon wafer is given.

In this case, however, the Si02 film has

been etched in H2 0/HF. The accumulation

time for this spectrum is identical to that

of figure 20. A composite of the peak positions observed in this experiment is reproduced as figure 22.

Five species of sodium are discernible and

the corresponding peaks are numbered in the

figure. With positive bias at 525°K, peaks

3 and 4 diminish in intensity, while peaks

1 and 2 increase. Application of negative

bias results in an increase in the intensity

of 3 and 4 with a decrease in 1 and 2 . No change is seen in peak 5 as a function of

bias. In a particular sample its intensity

appears to scale as the flood gun current; however, this effect has not been firmly established. Alternating positive and negative

bias fails to transfer more than 30% of the

intensity of peaks 3 and 4 to peaks 1 and 2.

Finally, as a function of increased intensity

of peaks 3 and 4 at negative bias, a splitting is observed in the apparent vacuum

level of the Si02 under the x-ray beam. The

splitting is of the order of 0.7 to 1.0 V and is quite reproducible.

A number of additional states were observed

in some samples but were not adequately re- producible to warrant including them in this

composite. The possibility exists, however,

that future work may resolve as many as 4 additional states of Na in Si02

•

Significant changes in the spectra of the

oxygen Is and silicon 2p regions were also

observed and correlated with ion movement. 23

Organic Surface Contamination

The fabrication of an MIS device can be re- duced to a sequence of chemical processes

which are implemented under clean conditions

at atmospheric pressure or at high vacuum. Throughout these process steps, the wafer is

exposed to a number of organic compounds

present in the gas phase. As a consequence,

all wafers have a surface coating of carbo- naceous material which is generally several

atomic layers thick.

Rinsing with organic solvents leaves this

overlayer unaffected. Oxidizing and re- ducing cleaning treatments modify the chemistry of the layer but do not remove it.

There is mounting evidence 24 that this overlayer persists after oxidation at 1000°C in

dry 02•

The presence of these overlayers can lead to bond failure and anomalous contact resistance in metallization systems. Tentative

overlayers of appropriate chemical characteristics can modify etchant performance.

Residual carbon doping at the Si/Si02 interface during gate-oxide growth may effect

device performance.

Unfortunately, thin overlayers of carbon are difficult to analyze. Electron- and ionstimulated surface analyses cause rapid

desorption of the layer in vacuum, and can induce chemical changes in its structure.

Since XPS involves the lowest magnitude of

electron flux to the sample surface, we have

157
begun a study of organic contamination in

MIS device processing.

Figure 23 illustrates the presence of such

films. This 0 - 1000 eV binding energy scan

is that of a kovar package interior. Gold,

copper, chromium, iron, cobalt, oxygen,

chlorine, fluorine and a number of other

elements are readily observed. The carbon

line at about 285 eV represents a major

component of this surface.

Simple element detection is not an adequate

use of XPS. The importance of this experiment is the ability to identify chemical

parameters. This is illustrated in figures

24 and 25. These represent scans of the

surface of a GaAs sample on which a thin

oxide has been grown. Experiments in a solar-cell research program at JPL have indicated that oxides grown in different con- ditions of carbon contamination had markedly

superior operating parameters. The sample

giving rise to figure 24 was significantly

superior to that of figure 25. In both

spectra, a number of aliphatic carbon states

can be assigned. Similarly, carbonyl or carbon-oxygen chemistry is found in both

samples. In the superior device an amide

carbon signal is observed. The presence of

amines in an epoxy used to attach the wafer

to a metal substrate had resulted in the incorporation of some of these compounds on

the surface of the GaAs. This observation

permits the design of a process to inject

this species at the surface. If this had

not been observed directly, the process

would have to be developed through random

optimization of conditions.

Figure 25 gives a final illustration of the

tenaciousness of surface adsorbed carbon.

This silicon sample had been etched in hot

H2 0, HF under dry N2 and directly introduced

to the spectrometer. This 0 - 1000 eV spectrum is now dominated by the silicon 2s and

2p lines. The fluorine Is line is clearly

observed as is the fluorine KLL Auger structure. Note, however, the intensity of the C

Is is still quite significant. In these ex- periments, silicones, freons, and simple

halocarbons can be readily detected. Proteins and lipids resulting from human handling can also be observed. In our work,

only cleaning in an 02 plasma resulted in

removal of the carbon overlayer.

Etchant Residues

Yet another of the fundamental chemical steps

in device processing involves the etching of

metals and oxides in conjunction with photoresist masking. Although the chemistry of

metal and oxide dissolution is well understood for bulk materials, little work has

been done to characterize the effect of

these solution techniques on the device surface.

One of the basic difficulties with liquid/

solid reactions concerns the possibility of

etchant residues deposited on the wafer surface. Often the presence of the surface

alters the solubility of the end products of

the etch and consequently such films are exceedingly difficult to remove. Proper re- moval requires the ability to directly monitor their presence. In figure 26, several

peaks are observed at high binding energy

which can be attributed to the presence of two or more forms of cerium oxide.

This sample was prepared from a wafer which

had a chromium/gold metallization over the

gate oxide. The metal contact was etched

with KI3 solution to remove the gold, followed by eerie ammonium sulfate solution to dissolve the chromium. This sample was then

treated with hot H2 0/HF to remove the Si02 film. After these processing steps, little

tri-iodide is observed on the surface, but

the cerium remains.

Such residues, if undetected, can then be

driven into the device by subsequent processing.

Figure 27 shows a narrow energy scan of the

cerium region showing the splitting of

states in the individual components of the

cerium 3d doublet. Exposure of the wafer

to chemical solutions of appropriate pH can virtually eliminate this oxide residue.

Heavy Metal Impurities

In the region of 920 and 950 eV in the spectrum of figure 28, components of the copper

2p doublet can be detected. Most of the

samples studied in our laboratory are fabricated on single-crystal silicon substrates

which have been polished by the copper-ion

polishing method. This polishing system involves exposure of the wafer to cupric ion

solutions. After polishing, these wafers

are pre-oxidized to a depth of several thousand Angstroms. This oxide is then stripped

to remove any heavy metal contamination as well as crystallographic damage introduced

by the polishing mechanism.

Our studies indicate that copper persists at the Si/Si02 interface in spite of these precautions. Figure 28 illustrates this situation with the copper 2p region of a freshly

etched wafer. The copper 2p doublet is

158
clearly resolved. A scan o of this sample

before removal of the 50 A Si02 film indicated that the copper was barely detectible

through the oxide. The intensity ratios of

the Cu 2p line before and after oxide removal

indicate that the copper is located at the

Si/Si02 interface, and is probably imbedded

in the elemental silicon. Repeated preoxidations slowly reduce the amount of

copper present.

Three types of copper can be discerned over

the range of wafers studied, and there are indications that the different species have

unique effects on the electrical characteristics of the device.

New Developments

threshold condition exists. Let the value

of Vfg for threshold be V

Q ; then, by scanning

Vfg from some bias setting such that Vfg

scans from Vfg < V0 to Vf > VQ

, the current

through the sample must abruptly rise at Vfg = VQ to a saturation value. This scan- ning is accomplished in a time shorter than

the relaxation time of the surface charge

generated during the fixed bias. The first

derivative of the current with respect to Vfg gives a maximum determining VQ

. In fig- ure 29, the cathode work function W

c

, the

electron affinity of the oxide Xox» t'ie Si02 /Si barrier energy, ()>g, and the silicon

bulk potential Vg are all constant. Therefore, changes in V

Q equal changes in V

g and

provide a hard energy reference for XPS of

insulators

.

One of the basic problems facing the application of electron-emission spectroscopies

to the study of semiconductor and insulator

samples is that of referencing the observed

data. As was developed previously, the

binding energy observable is a function of

the work function and surface potential. In insulator samples the work function and surface potential are not fixed, but, rather,

are a function of the x-ray flux density,

the emission currents, and the secondary

electron currents.

A new method has been devised which enables

measurement of hard reproducible values for

the work function and surface potential of

the sample during the XPS experiment . In order to illustrate this approach, figure 29

shows the relation of the energy levels of

the flood gun/vacuum/Si02 /Si system. All

XPS measurements are referenced to the Fermi

level E-p of the spectrometer which is fixed

to Ej, of the silicon substrate (a known bias

may be applied to the substrate) . The surface potential V

g of the sample determines

the fields in the sample hear the surface,

affecting the energy shifts in the XPS lines.

The magnitude of Vg is given by the sum

V = F w + V

s ox ss

where FQX is the oxide field , w is the oxide

thickness, and Vssis the surface potential

at the silicon interface.

In this experiment, the flood gun cathode is

biased to ground by a known voltage Vfg.

When the vacuum levels are such that

Efg > Eox , some of the electrons emitted

from the flood gun reach the sample surface,

charging the surface more negatively and

causing a current through the sample. When

Efg < Eox , these electrons are totally re- flected from the surface. When Efg = EQx , a

1

The current from the flood gun through the

sample is monitored with the experimental

arrangement shown in figure 30. The sample

is connected to ground through an ammeter

with a lock- in amplifier. A reference

oscillation is applied to the cathode of the

flood gun and phase-sensitively detected at the lock-in amplifier. The first harmonic

of the ac signal gives the first derivative

of the current.

Since the flood gun gives a broad diverging

source of electrons, maxima appear in the scan corresponding to both the x-ray irradiated and unirradiated region of the oxide as well as the gold mask region at the periphery

of the sample window. The maximum due to the gold mask is constant and equal to the

difference in work functions of the gold and

flood gun cathode. The irradiated region of

the oxide is shifted to more positive potentials by the positive charging effect of the

x-rays

.

Figure 31 gives a representative indication

of the variation in surface potential (of

the Si02 under the x-ray beam) as a function

of different flood-gun parameters. A range

of approximately 4 eV was observed in the case of a 30 A dry oxide thermally grown in

Si. An analogous shift was observed in the

relative positions of the Si 2p peaks for

Si02 and for Si, as a function of different

flood-gun parameters. Additional experimental detail concerning these experiments and

those planned with the electron probe have

been reported elsewhere.

Since this experiment can be accomplished in

a time scale of 10-8 to 10-9 seconds, the

ability to determine hard reference energies

permits energy lock-in and should lead to much narrower line widths. Also, the routine
intercomparison of data between laboratories 2. would be greatly facilitated.

Recent advances in sample preparation make 3. possible the fabrication of clean active

surface adsorbers. Placing these adsorbers

in contact with solutions and gaseous atmospheres should facilitate the sampling of 4. impurities at sub-part per billion levels.

Conclusions

XPS is a powerful tool for the study of the 5. surface and interface chemistry of MIS device processing. It offers reasonable sen- sitivity together with minimum interferences 6. for analytical capabilities. The chemical

shift observed in this spectroscopy is its

most important feature. The physics of the

experiment can be manipulated to give con- 7. siderable electronic/chemical information

about the insulator/semiconductor interface.

Since the observational depth of the method

is sharply defined, this chemical information is often lost in the application of con- ventional techniques. Consequently, this

data is often of critical utility in the 8. solution of processing problems.

Although technical advances on the part of 9. instrument manufacturers have made possible

promising new applications of XPS, several

difficulties remain. The most fundamental

difficulties are those of sample area and 10. time. Current state-of-the-art limitations

on sample size are about 0.1 x 0.3 mm. This

restriction is imposed by the physical dif- 11. ficulty of making high fluence x-ray sources

and focussing soft x ray. Electronstimulated methods, though more destructive 12. than XPS, offer a considerably narrower field

of view. This limitation is directly re- lated to the second fundamental difficulty:

time. Low sodium levels reported here re- 13.

quired observation times on the order of

50 hours. The corresponding SIMS experiment

would require less than one hour's accumu- 14. lation time.

15. References

1. Siegbahn, K. , Nordling, C. , Fahlman, A.,

Nordberg, R. , Hamrin, K. , Hedman, J., 16. Johansson, G. , Bergmark, T. , Karlsson,

S. , Lindgren, I. , Lindberg, B. , Electron Spectroscopy for Chemical

Analysis - Atomic, Molecular, and 17. Solid-State Studies by Means of Electron Spectroscopy (Almqvist & Wiksells

Boktryekeri AB, Stockholm, Sweden,

1967)

Brundle, C. R. , J. Vac. Sci. Technol.

11, 212 (1974).

Bremser, W. , X-Ray Photoelectron Spectroscopy, Topics in Current Chemistry,

p. 36, (Springer-Verlag, Berlin, 1973).

Fadley, C. S., Theoretical Aspects of

Electron Spectroscopy, NATO Conf. on Electron Emission, Ghent, Belgium,

1973.

Hollander, J. M., and Shirley, D. A., Ann. Rev. Nucl. Sci. 20, 435 (1970).

Henke, B. L. , X-Ray Optics and X-Ray

Microanalysis, pp. 157-172, (Academic

Press, New York, 1963).

Grunthaner, F. J., Griswold, T. W. , and

Clendening, P. J., Migratory Gold Resistive Shorts: Chemical Aspects of a Failure Mechanism, Proc. 13th Annual

Reliability Physics Conf., Las Vegas,

Nevada, April 1-3, 1975, pp. 99-106.

Fadley, C. S., J. of Electron Spectroscopy 5, 725 (1974).

Kriegler, R. J., Cheng, Y. C. , and

Colton, D. R. , J. Electrochem. Soc. 119, 388-392 (1972).

Grunthaner, F. J. , Klein, J. , Barton, J

(to be published)

Williams, R.,and Woods, M. H. , J. Appl.

Phys. 44, 1026 (1973).

Shumka, A. and Piety, Proc. 13th Annual

Reliability Physics Conf., Las Vegas,

Nevada, April 1-3, 1975.

Butts, A., and Coxe, C. D. , Silver,

Chapt. 34, (Van Nostrand Company, 1967)

Grunthaner, F. J., (to be published).

Kriegler, R. J. , Proc. 12th Annual

Reliability Physics Conf., Las Vegas,

Nevada, April 2-4, 1974, pp. 250-258.

Kerr, D. R. , Proc. 9th Annual Reliability Physics Conf., Las Vegas, Nevada,

1971, pp. 1-8.

Szedon, J. R., and Handy, R. M. , J. Vac. Sci. Tech. 6, 1-12 (1969).

160
18. Snow, E. H., and Deal, B. E. , Trans.

Metal. Soc. AIME 242, 512-523 (1968).

19. Hofstein, S. R. , Solid-State Electron.

10, 657-570 (1967).

20. Kriegler, R. J., Proa. 2nd International Symposium on Silicon, Materials and

Technology, Electrochem. Soc. Meeting,

Chicago, Illinois, 1973, pp. 363-375.

21. DiStefano, T. H. , J. Appl. Phys. 4<4_, 527 (1973).

22. Kriegler, R. J., Cheng, Y. C. , and

Colton, D. R. , J. Electrochem. Soc.

119 , 388-392 (1972); or see Grunthaner,

F. J. and Maserjian, J., Proc. 13th

Annual Reliability Physics Conf., Las Vegas, Nevada, 1975, p. 15-26.

23. Grunthaner, F. J., Lewis, B. F.

,

Griswold, T. W., and Maserjain, J.,

Differential Charging in X-Ray Photoelectron Spectroscopy, ACS Meeting,

Honolulu, Hawaii, June 12, 1975.

24. Clendening, P. J., Maserjain, J., and

Grunthaner, F. J., Chemical Structures

of Carbon Contamination at Si/Si02

Interfaces, submitted to J. Electrochem.

Soc.

161
TABLE I

EXPERIMENTAL CONSIDERATIONS

• Ultra High Vacuum

• "Clean" Residual Gas Background

• High Signal/Background Ratio

• Highest Possible Throughput

X-Ray • Analyzer • Detector

• Rotatable Samples

• Collimated X-Ray Source

• Electron Control

• Variable Temperature Capability

162
ELECTBOW EVENTS STIMULATED BY j RAY PHPTW ADSOHPTIOH

IIMlm Bmliiio* tmm llcctna klnlN

Figure 1. Electron events stimulated by

x-ray photon absorption.

PRINCIPAL AUGER ELECTRON ENERGIES

Binding Energy (eV)

Figure 3. Principal Auger electron energies

plotted by atomic number. Abscissa is given

at apparent binding energy (i,e., electron

kinetic energy minus 1483.6 eV).

PRINCIPAL PHOTO ELECTRON ENERGIES

U^r»fcJ--r jp *>J. "

I _ Lt_ l~r~ Tj1—i i

Binding Energy (eV)

Figure 2. Principal photoelectron energies

plotted by atomic number versus binding

energy for Al Ka^, 2 (1483.6 eV) excitation.

Atomic Number,

Z

Figure 4. Relative intensities of the

photoelectron lines normalized to sodium

line intensity =1.0 .

163.
Figure 7. Simplified diagram of photoelectron experiment.

164
X-RAY PHOTOELECTRON SPECTROSCOPY

Figure 8. Simplified diagram of optics of

x-ray photoelectron spectrometer.

Figure 9. Representation of excitation

sources and sample arrangement in sample

chart ber

.

Introduction

. Lock J

Controlled Growth System Si/SlO?

Figure 10. Block diagram of JPL XPS facility. Figure 11. Sample preparation diagram

giving x-ray illumination area.

165
Figure 12. Silicon 2p spectra of etched

silicon substrate.

t

107 105 103 101 99 97 — ENERGY, EV

T 1 1 1 1 1 1 1 1-

* lii... | 107 10} 10} 101 » 97 ENEKCV, EV

Figure 13. Silicon 2p snectra of freshly

etched silicon and a 42 A film of Si02 on Si. Figure 14. Silicon 2p spectra of 44 A Si02 on Si under positive and negative (upper

curve) bias.

166
Figure 15a. Silicon 2p spectra of a 113 A

film of Si0 2 on Si.

107 :

-

- , \ 97

^ EN ERGY, eV

Figure 15b. Spectrum of figure 15a but with

20 X gain.

167
168
SN0U13313

•H

CO u

c 3 rt H XI CD cu •H CU 4J

CO

aj c

CO )_i cO •pi •H

M-l M-l +j *o o CO o QJ *H cu

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CO

J_l ^

4J •u CO o O cO QJ

<u o cu 01 >2 "J Cu o CX i—

1

CO •H CO M-4 CO

,—

i

t. o

CO *H

r-| CO rH CO cu •r-1

4_l 4_)

Cfl c CO •H •H *H Q CU S3 to CO *"U

t-H Q r; B a T3 p Q

iH B QJ (J •H CO J2 Q m CO u O fa + M-l 4-1 0) CO <u 0) iH O 53 rH

•H 14-1 ^ U C/l CH u O CU 0) O CU

01 o

ft o CJ n CU

S3 <• cu rH cfl CO d > 4-1 *H u O M-l rH i—

1

D cni CU CN O CN CO CN p, H

CU a) CO aj c CU 0) M i—

i

M M o u

3 Pu 3 3 0 CO

C>0 e 4-1 oa- 00 •H

•H cO CO •rl CO •H 43

CO (3 Z Pm 4-»

S NO Mi 031 3

'

169
170
171
Figure 29. Band diagram showing electronic levels of

Si/Si0 2 /vacuum system with electron probe.

0 1 I I I 1 I i— 10 8 a 4 2 0 -2

-«! FLOOD GUN VOLTAGE

Figure 30. Diagram of bias experiment and Figure 31. Surface potential variation as appropriate electronics. a function of flood gun voltage setting for

a 30 A Si02 film on silicon.

172
DISCUSSION OF THE PAPER

Morabito : You skipped over the gold migration problem somewhat. Could you go into more detail?

Grunthaner : The gold migration problem can be treated from two points of view. The

physical cause is the failure of the part's

hermeticity. This resulted in the injection

of water into the package during a thermal

shock test. The chemical cause of the problem is the simultaneous presence of the

halide ion for gold ion complexation and a thin film of water on the surface of the

part. The existence of water and applied

potential would have resulted in slow corrosion of the metal stripes. This would have

been of little long-term concern. The presence of the halogens permitted gold transport. The glassivated coating has a multitude of open cracks and incomplete coverages

in the region of the gold. The combination

of strong complexation of the gold ion and

high current densities resulted in dendritic

growth. The morphology of the dendrite is

related to the strength of the gold complex,

since this, together with the actual concen- tration of gold ions, determines the effective concentration of gold at the depositing

electrode

.

Initially, we had felt that an extremely

clean electrode surface was required for

dendritic growth. Subsequent experiments

indicated that high-current density and low source ion concentrations were the limiting

requirements. The presence of iodine was

due to residual etchant which was strongly

adsorbed to the surface of the remaining

gold. Bromine and chlorine were also found,

although the sources of these ions have not been clearly established. The chlorine may

have been present from human handling or may

have had a significant concentration in the

water entering the package.

Morabito : Was chlorine more of a problem

than iodine?

Grunthaner : The gold-chloro-complex is much

more stable than the iodo-complex. In large

gold ion concentrations, the chloro-complex

would be most effective in generation of

dendritic growth. In these lesser concentrations the iodo-complex provides transport,

and reasonable concentrations. The polarized iodine on the gold surface also assisted

in the dissolution of the gold stripe. The

chlorine will be active but the iodine is probably the more important immediate problem. Bromine and bromate were also observed.

The bromine case would be intermediate between chloride and iodide.

The primary question to our program was the

possibility that there existed a long-term

failure mechanism, i.e., parts failing after

one or more years of operation. This fail- ure mechanism was observed over periods of

hundreds of hours with multilayer coverage

of H2O over the surface. Was failure at ex- tended times possible at one- or two-layer

coverage? The situation is resolved into

ionic and hydrogenic currents and current

study is addressing this problem.

The data available at this point are adequate

to suggest several possible corrective ac- tions to minimize MGRS failures: 1) Reduce

the water content of the package to several

hundred ppm (by volume) or less. 2) Reduce

the impurity levels at the electrode and

substrate surfaces, so as to eliminate

transport. 3) Getter reactive impurities by

chemical reaction, thereby modifying their

chemistry, or by use of clathrates to immobilize them. 4) Treat the surface of the

device to produce a hydrophobic surface,

such as by use of methyl silane or other

silicone derivatives.

173

NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg , Maryland, April 23-24, 1975 (Issued March 1976)

CHOOSING BETWEEN ESCA AND AUGER FOR SURFACE ANALYSIS

Gary E. McGuire

Materials Characterization Laboratory

Texas Instruments Incorporated

Dallas, Texas 75222

ABSTRACT

Although photoelectron and Auger spectroscopy were developed over similar

time frames and required similar hardware development, until recently the

two techniques have remained as separate tools for surface analysis.

Auger spectroscopy was readily accepted by the semiconductor industry for

materials characterization. Interest in photoelectron spectroscopy did

not develop as quickly nor did it become as widespread as Auger

spectroscopy

.

The two techniques being quite complementary can be used very well

within the same laboratory if not within the same vacuum system. A num- ber of examples will be discussed for which one would choose by preference ESCA over Auger. Although the choice may not be apparent beforehand, ESCA can be the more valuable characterization technique in certain

aspects of polymer surfaces, electroplating and thin film metallization

systems for semiconductor devices.

INTRODUCTION

If one takes into consideration all the com- mercial and non-commercial designs of electron spectrometers there are two basic differences that continue to distinguish the Auger electron (AES) spectrometer from the X-ray photoelectron (ESCA) spectrometer,

resolution and excitation source. Auger

spectrometers emphasize high transmission

and use an electron beam for excitation while

ESCA spectrometers emphasize high resolution

and use soft X-rays for excitation. The soft

X-ray source irradiates a much larger sample

area with less intensity than an electron

beam. Soft X-rays cause less damage to the

sample, allowing longer times for data ac- quisition without altering the surface. The

high resolution capabilities of the typical

ESCA spectrometer and simplicity of the

photoelectron spectra permits the identification of chemical shifts.

The two factors, resolution and excitation,

ultimately govern the choice of which technique to use for a particular problem. The

next few examples will be used to demonstrate

occasions where these factors favor the use

of ESCA. Because of the speed of data ac- quisition and excellent spatial resolution

of Auger spectroscopy, it has been readily

adapted for analysis of semiconductor surfaces. The examples presented here are not intended to downgrade the capabilities of Auger spectroscopy but are an effort to show

how one would choose between the two techniques intelligently. These examples are slanted toward ESCA because we feel it has

not been exploited as much in problem solving

for semiconductor materials.

Chemical Identification

A feature that has always been one of the major strengths of ESCA has been the ability

to observe chemical shifts in the spectra

corresponding to the chemical states of the

system being studied. For example in Figure

1 the separation between the peak positions

for Si and Si02 is approximately 4.5 eV. The line widths are sufficiently small that

the peaks are well resolved and easily distinguishable. The corresponding Auger chemical shift for Si and Si02 is much larger at

13 eV. The two chemical states may be easily

identified in the Auger derivative spectrum

but the transitions cannot be completely

separated. Frequently it is possible to en- hance certain features using integration

techniques; however the integrated Si-Si02

Auger transitions overlap extensively

(Figure 2)

.

In general photoelectron line widths, ex- pressed as full width at half maximum (FWHM)

,

are usually narrower and less complicated

than the corresponding Auger line widths. Most Auger series consist of a family of curves with one or two dominant features. Even in examples where the AES shift is much

175
larger, the chemical shift in the ESCA spectrum continues to be more easily distinguished. This may be enhanced by the higher

resolution capability of the typical ESCA

spectrometer (0.02%) over that of the typical AES spectrometer (0.5%).

With the numerous cited examples of chemical

shifts in the literature, one may get the

impression that every chemical combination

produces a unique combination of photoelectron binding energies that are readily distinguishable. Although the combinations may

be unique, the magnitude of the shifts may

be so similar the peaks overlap extensively.

These then may not be reduced to meaningful

data except by comparing elemental ratios

which could also be accomplished from the

corresponding AES spectra.

Frequently the simplest method of studying a process that malfunctions is to compare

"good" and "bad" samples in order to identify

the presence of anything that may be detrimental or unusual. Numerous useful examples

of this type may be found throughout the

literature. Typical AES spectra of goldplated surfaces which displayed "good" and

"bad" plating characteristics are presented

in Figure 3. The AES spectra reveal the

presence of large deposits of K, C, N and Sn

associated with "poor" plating characteristics. The same information concerning the

elemental composition can just as easily be

obtained by ESCA. In addition, ESCA gives

some insight into the failure mechanism.

From the AES data, it is not apparent whether

the presence of K, C and N on the gold plated

surface is due to adsorption of the gold salt

KAu(CN)2 on the surface or due to coprecipitation of the end product, K+ and CN~. By

comparing the binding energies of gold 4f lines (Figure 4) from evaporated Au, KAu(CN)2,

a commercial gold salt, and the gold plated

surface, we get an indication of the failure

mechanism. The binding energies of the gold

plated surface agree with those of evaporated

gold while those of the commercial gold salt

and KAu(CN)2 agree. From this figure one

could make the observation that the gold

plated surface displayed the same chemical

behavior as evaporated gold. The potassium,

carbon and nitrogen present on the surface

are end products of the plating process that are trapped in the gold. The similarity in

the binding energies of the commercial gold

salt and KAu(CN)2 suggests a common chemical

form for the two. Correlations such as this

may be used to obtain substitute chemicals

for expensive commercial formulations.

Excitation Source Influences

It is generally accepted that the influence

of an X-ray source is less severe than that

of an electron beam. Electron induced desorption and decomposition can alter the

sample sufficiently to make data from some samples meaningless. (Bremsstrahlung from

the X-ray source may produce the same re- sults but to a smaller extent.) It is very

difficult to obtain AES data from rough insulating surfaces such as printed circuit

board because of charging induced by the

primary beam. Examination of the same types

of materials by ESCA are straightforward and

simple. In Figure 5 is a representative

ESCA spectrum of printed circuit board which

we were unable to analyze by Auger. The

analysis picked up the presence of Pb and Na

on the surface even in areas away from the

Pb-Sn plate.

If electron induced desorption or migration

occurs then the Auger spectrum may not distinguish differences in samples with widely

varying surface properties. In the following example, we were unable to detect any

differences in surface composition by AES

while the ESCA data displayed some significant variations (Figure 6) . The samples

examined were clean tin oxide, tin oxide

treated in a N2 atmosphere and tin oxide

treated in an 62 atmosphere. The untreated

surface exhibits a clean tin oxide spectrum

as observed by the Sn3d3/2~3d5/2 doublet.

Treatment in a nitrogen atmosphere results

in a small amount of Na appearing on the

surface. The tin oxide surface is masked by

a thin layer of Na when an O2 atmosphere is

used. The presence or absence of Na was verified by the Na KLL Auger transition that occurs in the same energy region of the ESCA

spectrum as the Sn3d doublet. The important

point we are trying to make is that examination of the same samples by AES does not reveal the presence of Na at all while ESCA

indicates significant differences that cor- relate with the known physical properties of

these films. We also observe a reversal in

the physical properties of these films to the original state upon examination by AES

(electron beam exposure) which has been cor- related with the absence of Na.

Metal Contact Systems

ESCA may prove to be one of the most useful

techniques available for examination of

metallurgical contacts and thin film metallization systems, particularly heavy metals.

176
Most of the heavy metals, Au, Pt, Ta, W,

used in semiconductor schemes have very in- tense 4f transitions with very narrow FWHM. The AES spectra for these same metals have

complex structures with a number of observable transitions, none of which are particularly intense. The high energy Auger tran- sitions which are the most distinctive of

these metals have broad FWHM and require a 5 to 10 keV primary beam in order to optimize

cross sections.

Using ESCA one could examine both implanted

and alloyed contacts. In the case of implanted heavy metal contacts, the region of

o maximum concentration is approximately 200 A below the surface. An example of this type

is shown in Figure 7 of an Au-implant in

silicon. In this example the surface has

been chemically etched to the region of maximum concentration. Since the sample was exposed to air after etching, a transition

due to Si02 from residual oxide formation

appears in the spectrum accompanying tran- sitions due to Si and Au. Although the implanted dose is rather large (6.8 x 10-^

ions/cm^), it would be possible to detect

doses of several orders of magnitude smaller.

The same results may be obtained by ion

sputtering as opposed to chemical etching.

Since the signal comes from a large area it

is necessary to use a mask in order to analyze a uniformly sputtered region. This re- sults in a reduction of signal by a factor

of three. Since the dose required to implant metal contacts is within the detection

limits of ESCA and since the implant is to such shallow depths, ESCA is useful in the

study of doping profiles and compound

formation.

The alloyed Pt contact to Si is presented as an example of the usefulness of ESCA in the

study of compound formation. This system is

commonly used for integrated circuits and

silicon Schottky barrier diodes. When a thin film of platinum is deposited on a silicon substrate and the couple sintered,

platinum silicide is formed. The stoichiometry has been identified as PtSi.^" After

the system is sintered, any unreacted Pt is

removed by chemical etching with aqua regia.

Contact resistance measurements as well as the etching characteristics indicate the

presence of an additional high resistivity

layer at the Pt-PtSi interface.

Rand and Roberts examined this system by

Auger spectroscopy and identified the resistive layer as SiC>2. When we examined the

same system by ESCA,-^ we came up with a different interpretation. In Figure 8 are the

ESCA spectra for the three different regions

of the Pt-PtSi system, the surface, interface

and PtSi region. From these spectra can be

identified three different valence states of

silicon (99.8, 102.3, 103.3 eV) and three

distinct valence states of platinum (71.0,

71.4, 72.4 eV) . The lowest binding energy

Si2p peak was identified as being due to elemental silicon. The presence of free silicon is observed in each of the three spectra

and acts as a reference point. The presence

of elemental silicon on each of these sur- faces may be due to a combination of rapid

silicon diffusion and reduction of silicon

compounds due to Bremsstrahlung radiation

from the X-ray source. Neither of the re- maining two Si2p transitions can be identified

as due to Si02 since the Si2p binding energy

in Si02 is 104.3 eV. The Si2p transition at 103.3 eV appears to be associated with the

Pt4f 7 / 2 72.4 eV transition while the 102.3 eV Si2p transition appears to be associated with

the 71.4 eV Pt4f7/2 transition.

The Pt4f7/2 binding energy of 71.0 eV was identified as platinum metal. This assignment agrees with other literature values.^

The line shape of this Pt4f -j /2 transition is

slightly broadened due to the presence of a second chemical state of Pt with a binding

energy slightly higher than Pt metal. When

the surface is chemically etched to the

Pt-PtSi interface, the 72.4 eV Pt and 103.3

eV Si peaks dominate the spectrum. Since

Si02 has a Si2p binding energy of 104.3 eV

its presence at the interface is ruled out. From our experimental results, we have observed that the binding energies of Pt and Si

at the interface region displayed associated

shifts. This fact has been taken as an indication of compound formation involving Pt, Si

and 0. Comparing the intensities of these

transitions with those from pure Pt and Si

targets indicates a stoichiometry of PtSi04.

Examination of the platinum silicide region

is easily achieved by ion etching beyond the

interface region. The binding energies displayed by Pt and Si in this region are 71.4

and 102.3 eV respectively.

By the use of ESCA it was possible to identify

the layers of the metallization system as well

as the compounds being formed. Careful examination of the intensities of the transitions

may also be useful in determining the stoichiometry of the system. Other analytical

techniques are available for structural identification when sufficient material is available but these metal systems may only be

several hundred angstroms thick.

77
SUMMARY

REFERENCES

ESCA has proven useful for many applied

semiconductor research problems. Except for

spatial resolution and speed, it offers the same advantages as AES as a surface characterization technique. In addition it may be

used when sample charging, decomposition and

electron induced migration are potential

problems. It also offers the advantage of

high resolution with resultant chemical

shift information. The simplicity of the

photoelectron transitions are easily reduced

to quantitative information.

In the future one can expect to see ESCA

used to obtain "chemical in-depth profiles"^

as opposed to the more common elemental profiles. ESCA will also prove to be useful in

the study of diffusion mechanisms^ with the

added dimension of elucidating compound

formation.

1. Muta, H. and Shinoda, D. , J. Appl. Phys.

43, 2913 (1972); Poate, J. M. and

Tisone, T. C. , Appl. Phys. Lett. 24, 391

(1974).

2. Rand, M. J. and Roberts, J. F. , Appl.

Phys. Lett. 24, 49 (1974).

3. Danyluk, S. and McGuire, G. E. , J. Appl.

Phys. 45, 5141 (1974).

4. Bearden, J. A. and Burr, A. F. , Rev. Mod.

Phys. 39, 125 (1967).

5. McGuire, G. E. , unpublished results.

6. Danyluk, S. , McGuire, G. E. , Koliwad,

K. M. and Yang, M. G. , Thin Solid Films

25, 483 (1975).

178
SILICON SLICE WITH RESIDUAL OX I DE

800

co 400

110

Si 2p

Si0 2 ,l043eV Si.99.8eV

Figure 1. X-ray photoelectron spectrum of a silicon slice with a residual oxide coating.

The Si(2p) transitions of Si and Si02 display

a chemical shift of 4.5 eV.

106 102 98 BINDING ENERGY (eV)

Figure 2. The characteristic minima at 78 and 91 eV for Si0 2 and Si respectively can be observed in the derivative Auger spectrum.

The integrated curve is superimposed over the

derivative curve.

Gold Plated Samples

Au "~ /Au C source #2, good plating characteristics

Y Au Ij N Sn q

| source #1, substrate #2

[Ail K Iq poor plating characteristic!

t^. " Cu Cu /Au ,fNSn o

source #1, substrate #1 Au poor plating characteristics K

i 1

C

I 1 I 1 I 1 l 200 400 600 800

Electron Energy (ev)

65 10 15 W 85 90 95 ELECTRON ENERGY (eV)

Figure 3. Auger spectra of gold-plated

surfaces. The presence of potassium, carbon

and nitrogen was identified with Au-plated

surfaces with poor bonding characteristics.

179
Tin Oxide

Sn3d3/2

and Na Auqer

Sn3d 5/2

500 495 490 485 Binding Energy (eV)

300 200 BINDING ENERGY (eV)

Figure 5. X-ray photoelectron spectrum of printed

circuit board. Lead and sodium were found on the

surface and associated with leakage currents causing the printed circuit to fail.

Figure 6. X-ray photoelectron spectrum of tin

oxide surfaces. Treatment of the tin oxide surface in an O2 or N2 atmosphere leaves a thin

coating of Na. The sodium is induced to migrate

away from areas exposed to an electron beam as in

Auger analysis.

180
Au Implant in Silicon

U5 105 95

~ 85 ~

Electron Binding Energy. eV

Figure 7. X-ray photoelectron spectrum of Au implant in silicon. The surface has been chemically

etched to the region of maximum concentration.

Figure 8. X-ray photoelectron spectrum of Pt contact to silicon. Three spectra are

shown: (a) platinum surface, (b) sample

that has been chemical etched of excess

platinum, and (c) sample that has been chemically etched then ion etched to the PtSi

region.

EVAPORATED PLATINUM UN Si (111) SINTERED 630°C

MO 90 80 BINDING ENERGY (eV)

181
DISCUSSION OF THE PAPER

Grunthaner : I would like to comment about

your statement concerning the line widths of

the gold as opposed to silicon. The silicon

lines in the spectra are probably the narrowest lines that one observes at all, although the presence of the silicon 2p doublet, when broadened by the instrument and

that sort of thing does give rise to a some- what broader line. The silicon lines that we have seen in the spectrometer are of the

order of about 520 millivolts wide. The

narrowest gold lines we have observed are of

the order of about 680 millivolts wider.

McGuire : Most people get the impression

looking at the commercial brochures that the

gold and the silver give the narrowest lines,

but sometimes these elements are merely the

most convenient, they are good conductors

and the separation between the doublet is

quite large.

Morabito : Because of the difference in the

escape depth between the ESCA photoelectron

as compared to the Auger electron, it is

possible that, for the last example you

showed, the layer within 5 or 10 A of surface could be Si02. You are sampling a much

deeper region with ESCA, so below you may

have platinum silicide but above that it is

possible you have Si02-

McGuire : I would like to make a comment

concerning the escape depth. I have seen both Auger and ESCA papers that have come up with different escape depths but I have a hard time believing that an electron, whether

it is due to an Auger process or a photoelectron process, will behave differently as far as its coming out of a given material.

I do not think they should differ.

Morabito: That depends on the kinetic

energy of the electron.

McGuire : Yes, so I would say electrons from

either process with the same kinetic energy

should have identical escape depths.

Participant : I would like to point out that

in stripping off the platinum and looking

again whether you have not really done anything to the surface, which presumably would

have the oxide or silicate on it, you could

be changing the Fermi level of the sample

with respect to the vacuum, which may give

you a shift in the lines.

McGuire: Well, in the spectrum I showed we were using the elemental silicon as a reference point. That is the reason we used the

data we got from the McPherson instrument,

i.e. , it was very convenient. We had a lot more difficulty with the Hewlett-Packard

instrument because we did not always see the

presence of the elemental silicon.

182
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

SILICON DEVICE APPLICATIONS USING A COMBINED ESCA/AUGER ANALYSIS SYSTEM

L. E. Davis and G. E. Riach

Physical Electronics Industries, Inc. Eden Prairie, Minnesota 55343

ABSTRACT

Silicon device applications are described which emphasize the advantages

of measuring both X-ray induced photoelectrons and electron beam stimulated Auger electrons as a function of depth in the same analysis system

The measurements include depth-composition profiles and chemical binding

information for nichrome thin films and boron predeposits on silicon.

INTRODUCTION

ESCA and Auger electron spectroscopy have

much to offer in the study of semiconductor

devices. 1 In conventional Auger systems,

electron beam excitation allows localized

surface analysis and rapid signal readout.

ESCA, on the other hand, offers a means of characterizing chemical bonds at the surface.

When used in conjunction with sputter-etching, both techniques can be used to obtain

data as a function of depth. The applications to follow illustrate how a combined

ESCA/Auger electron spectrometer can be used

to give a more complete characterization of thin films than either technique alone.

EXPERIMENTAL APPARATUS AND METHOD

The schematic diagram of a commercial

electron energy analyzer optimized in design

for high transmission and energy resolution

is shown in Figure 1. In the ESCA analysis

mode (via pulse counting) the essential advantages of the analyzer are: 1) selective

area analysis within an area 2mm in diameter

and 2) adjustable energy resolution. In the

conventional Auger analysis mode (via the

lock-in amplifier) the high transmission

(approximately 7%) of the analyzer allows

real-time measurement of the Auger electron

peak-to-peak heights as a function of sputter-etching time. A complete description of the electron spectrometer functions and

specifications have been reported by

Palmberg

.

2

The analysis was performed in an ultrahigh vacuum system capable of reaching pressures

less than lO- -^ Torr. The electron spectrometer, X-ray source and sputter ion gun

are pre-aligned so that when the specimen is positioned for optimum (photoelectron or Auger electron) signal, the ESCA/Auger analysis area and ion beam coincide. In this

manner, X-ray photoelectron and electron

beam excited Auger electron measurements

could be made after ion beam sputtering

without moving the specimen. The typical

analysis procedure was to record the Auger

electron peak-to-peak [dN(E)/dE] amplitudes

simultaneously with argon ion beam sputtering and, thus, obtain a depth-composition

profile. 3 High energy resolution ESCA analysis was performed at selected depths by

turning off the ion gun at the desired depth,

making the ESCA measurement, and then re- suming the Auger depth profile measurement.

The two types of data are then easily

correlated

.

APPLICATIONS

Oxidation of a Nichrome Film

As indicated in the Auger electron depthcomposition profile of Figure 2, the nichrome

film after heating in air for 30 seconds at 450°C is composed of a metallic oxide layer

at the surface. A chromium rich layer dominates near the surface while nickel has

formed a distinct layer at the metal/silicon

interface. To gain a better understanding

of the effect this heating process had on the chemical bonding, high energy resolution

ESCA measurements of chromium and nickel were made at the surface and at depths A, B,

C and D indicated on the depth profile.

This data is shown in Figure 3.

The chromium and nickel 2p3/2 photoelectron

lines are shifted to higher binding energies

(approximately 2.5eV for chromium and 3.0eV

for nickel) at the surface. These shifts

are consistent with ^203^ and NiO 5 oxidation

states. Slightly below the surface, at depth

A, the Ni 2pj/2 line nas shifted to a binding

energy consistent with elemental nickel. The

chromium, however, remains in oxide form

until depth B where both oxide and metallic

183
states are apparent. At depths C and D,

both chromium and nickel are predominantly

metallic

In conclusion, it can be seen that heating

this nichrome film in air produced a surface

oxide layer about 100A thick consisting

mainly of (^203. Within ISA of the surface

a small amount of NiO was formed. Deeper

than this within the structure, however, the

nickel remained entirely in the metallic

state.

Boron Predeposit on Silicon

In the doping of silicon with boron, the

first processing step is the deposition of a shallow layer of boron on a silicon wafer.

The predeposit is produced by allowing a boron gas (e.g., a borane) to react with the

silicon surface at elevated temperatures.

As Figure 4 indicates, a glass surface layer

is formed during the predeposit process.

The next step in doping is to chemically re- move the boron glass, leaving only a boron

rich silicon surface. The remaining boron

then is diffused at high temperatures into

the bulk silicon to the desired depth and

concentration level. In some cases the predeposition is carried out at a high enough

temperature that the diffusion occurs simultaneously with the predeposit.

In order to achieve reproducible dopant

levels, it is of considerable aid to the

process engineer to know the properties of

the predeposit before glass removal. The

main properties of interest are (1) the

thickness of the glass layer, (2) the boron

concentration as a function of depth and (3)

the chemical state of boron and other elements as a function of depth.

The depth-composition profile of Figure 4 illustrates that by monitoring the Auger

signals of Si, B and oxygen during sputteretching, the predeposit glass layer and boron

density can be rapidly characterized as a function of depth. In the example, the boron

is entirely within the glass overlayer. The

glass removal and diffusion steps have not been carried out on this specimen at this

point. At the depths A, B, C and D indicated on the plot, the profile was interrupted so that high energy resolution ESCA measurements

could be made to determine the chemical state

of the important components of the predeposit

layer

ESCA spectra of the Si 2p and B Is photoelectron lines at the depths indicated in Figure

4 are shown in Figure 5. The Si 2p binding

energies (103eV) at depths A and B are

consistent with Si02. At depths C and D the

elemental silicon lines (at 97.8eV) are also

present. The boron photoelectrons at 193eV

do not change in binding energy as a function

of depth. B Is electrons at this binding

energy belong to the oxide state of boron.

The depth-composition profile in Figure 6 shows the results of a combined boron predeposit and diffusion carried out at 1050°C

on polycrystalline silicon. Note that here

the boron is concentrated at the glass/silicon interface. The ESCA spectrum in Figure

7 obtained at depth A shows that two B Is photoelectron lines exist at the glass/silicon interface. The photoelectrons at 193eV

correspond to the boron oxide state, the same binding energy as the B Is electrons

shown in Figure 5. The lower binding energy

electrons at 186. 2eV belong to elemental

boron.

These examples point out the usefulness of

joint ESCA and Auger analysis for complete

characterization of thin films on silicon

surfaces. Both composition and chemical

binding information as a function of depth

are obtainable without changing specimen

orientation. In addition, the examples show

that the information obtained is in a form

that can be applied directly for definitive

process control and evaluation of silicon

devices

Acknowledgements

The authors gratefully acknowledge Drs. N. C. MacDonald and P. W. Palmberg for suggesting

the application topics and supplying background information about specimen processing.

REFERENCES

1. McGuire, G. E. , Choosing between ESCA

and Auger for Surface Analysis, ARPA/NBS

Workshop IV, Surface Analysis for Silicon

Devices, A. G. Lieberman, Ed., NBS Spec.

Publ. 400-23 (January 1976).

2. Palmberg, P. W. , Vac. Sci. Technol. 12,

379 (1975).

3. Weber, R. F. , Research/Development 23 , 22 (1972).

4. Kim, K. S., Purdue University, private

communicat ion

5. Fiermans, L, Hoogewijs, R. , and Vennik,

J., Surf. Sci. 47, 1 (1975).

184
£^ PRECISION ANALYZER

SPECIMEN

\ \ SPUTTER ION GUN

Figure 1. Schematic diagram of the ESCA/

Auger Electron Spectrometer (Physical Electronics Industries, Inc., Model 548).

Figure 2. Auger electron depth-composition

profile of a thin nichrome film heated

(450°C for 30 seconds) in air.

150A

.' («) (|) © (6)

Cr

Si

Ni

10 15 20 25 30 ETCHING TIME, MINUTES

35 40

Figure 3. High resolution ESCA spectra of

Cr 2p3/ 2 and Ni 2p 3 / 2 photoelectron lines as

a function of depth.

578 576 574 572 858 856 854 852 850 BINDING ENERGY, eV

185
Figure 4. Auger electron depth-composition Figure 5. Silicon 2p and boron Is photoprofile of a boron predeposit on single electron spectra as a function of depth,

crystal silicon.

Figure 6. Auger electron depth-composition Figure 7. Boron Is photoelectron lines at profile of boron diffused into polycrystal- the glass/silicon interface,

line silicon.

186
DISCUSSION OF THE PAPER

Thomas ; Did you say that the boron had been

oxidized by the glass structure?

Davis : No, I did not mean to imply that. My

interpretation is that the boron in the first

predeposit example exists as boron oxide,

probably as a borosilicate. The Auger electron peak, shape of boron in the glass layer

(not shown) was quite different from the elemental boron peak shape. This, of course, is

additional evidence that the boron in Figure

5 is in oxide form.

187

NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

PHOTODEPOPULATION TECHNIQUE FOR THE STUDY OF ELECTRONIC TRAPS IN INSULATORS*

T. H. DiStefano

IBM Watson Research Center Yorktown Heights, New York 10598

and

J. M. Franz IBM System Products Division

Essex Junction, Vermont 05452

ABSTRACT

A new type of photodepopulation spectroscopy has been used to obtain

the energy distribution of electrons trapped in an insulator. The

change in the trapped charge produced by monochromatic light is detected by a silicon surface effect device, which serves as one of the electrodes. The result is the spectral photoyield for trap ionization,

free of any complication of extraneous currents. This photodepopulation technique was applied to the Si02-Si 3N4 system to determine the

trap distribution responsible for charge storage in the MNOS transistor.

The results, although somewhat complicated by photoinj ection from the

electrodes, show an ionization energy of 3.0 eV. Generally, the technique is applicable to the very sensitive spectroscopic measurement of deep traps in wide bandgap insulators.

INTRODUCTION

A new photodepopulation technique-'- has been

developed and used to examine spectroscopically the impurities and trapped electrons

in insulating thin films. Essentially, the

technique involves a measurement of the re- laxation of the charge stored in an insulator that is exposed to monochromatic light. From the discharge of the traps produced by

illumination from a scanned monochromator

source, the spectral yield for photoemission

from the traps is obtained directly. The

charge in the insulator is detected by means

of one of the electrodes, which is a field

effect device that serves as an electrometer

Either of two modes of detection, small signal capacitance or surface conductance, is satisfactory for measuring the field at the

electrode from which the net charge is determined. The photodepopulation technique

has an advantage in that it samples only the density of occupied trap levels, and it is unaffected by extraneous sources of current.

It also has an advantage over a photocapacitance technique^ used by Harari and Royce to study traps at the SiO2~Al203 interface; our

* Work supported in part by the Defense Advanced Research Projects Agency and monitored by AFCRL under contract number

F19628-74-C-0077.

photodepopulation technique is not influenced

by a redistribution of the electrons among

the various trap levels in the system caused

by light-* or by thermal excitation. However,

as with any photoionization technique for thin film insulators, our technique is limited to a window equal to the electrode barrier energy; for photon energies greater than

this barrier, the photodepopulation results

are complicated by injection from the electrodes .

We have used the photocapacitance technique

to determine the distribution of the specific

traps, near the Si02~Si3N4 interface, which

are filled and emptied by electron tunneling

during the normal operation of an MNOS device

Electron traps in the Si3N4 itself have been

well characterized by several techniques including thermally stimulated currents^ »^ and

I-V measurements k ' ^ . Both techniques are able to determine the binding energy of traps

which are characteristic of the bulk of Si^N4

Binding energies are found to lie in a range from about 0.5 eV to 0.9 eV, as measured by

thermally stimulated currents and from about

0.5 eV to 1.5 eV determined from the I-V dependence. However, neither of these techniques is capable of being used to characterize the deep trapping levels associated with

charge storage near the Si02~Si3N4 interface,

since they are obscured by the shallow traps

in the conductivity measurements. It is the

189
deep traps which are of specific interest

since these states near the interface of SiC>2 and Si3N4 are responsible for the charge

storage in MNOS memory transistors. Models

for switching8 and charge retention^ in

the MNOS transistor would benefit by a determination of the deep trapping states. The

spectral photoyield was obtained for traps

filled to various charge levels by tunneling.

From the photoemission spectra and thresholds, information is derived about the energy

distribution of the deep traps in the Si02~

Si3N^ structure.

EXPERIMENT

The capacitive mode of charge detection was

used for the photodepopulation study of trapped electrons near the Si02~Si3NA interface.

To some extent this particular technique is

a variation of photocapacitance spectroscopy,

which has been used to study traps in the depletion region of a semiconductor junction.

During the measurements, a bias voltage is

applied to the MNOS sandwich to maintain a null electric field at the silicon surface,

as is shown in Fig. 1. The zero surface

field condition was sensed by a capacitance

measurement at 100 kHz and maintained by a feedback loop in which the capacitance measurement is used to control the sample bias.

The capacitance is measured by a phase sen- sitive detector in order to reject leakage

currents and the photocurrent . In this

"flat-band" condition, the bias voltage is

directly related to the net charge density

stored near the Si02~Si3N4 interface,

P =

e e V

0 1 fb

(1)

where p is the net charge density, e± is the

dc dielectric constant of the Si3N4, Vf^ is

the bias voltage necessary to maintain the

"flat-band" condition, and d is the thickness

of the Si3N^ layer. It was assumed that the

stored charge lies near the Si02-Si3N^ interface; the assumption is reasonable in this

case, where the charge was introduced either

by tunneling from the silicon or by photoemission from the silicon.

The photoresponse of the system was obtained

by shining monochromatic light through the

semitransparent aluminum electrode and onto

the traps near the interface. The light source was a Bausch and Lomb 0.5 meter monochromator with a Hg-Xe high pressure arc

source. The emergent light was de-focused

to cover uniformly the area of the sample

electrode. The light intensity was measured

with an RCA 1P42 photodiode covered by a layer of sodium salycilate phosphor; the

photodiode was calibrated against an Eppley

thermopile. Measurements were obtained for

photon energies from 2 eV to 6 eV, which is

the upper practical limit of the light source. A filter was used at energies below

3.5 eV to eliminate second order light in a range over which the sample was extremely

sensitive. All of the spectra were obtained

at room temperature. The measurements were performed on samples which are characteristic

of an MNOS transistor. The sandwich struco ture comprised a silicon substrate, 25 A of

Si02 Q formed on the silicon by oxidation,

535 A of Si3N4 formed by chemical vapor deposition at 875°C, and finally 125 A of

aluminum to form a semitransparent electrode.

Samples were formed on (100) 2Q-cm silicon

substrates of both n and p type. The semitransparent electrodes, 0.060" in diameter, were deposited by evaporation in a vacuum of

about 2 x 10~6 torr. The samples exhibited

typical hysteresis effects, with a maximum

total stored charge of about 8 x 10-^ electrons/cm^

Light of sufficient energy Kw excited electrons from the deep traps into the conduction

band of the Si3N4. For the condition shown

in Fig. 1, excited electrons are collected

by the aluminum electrode; re-trapping of

the electrons in the Si3N4 was not found to be a significant factor at room temperature

and at fields above 10^ V/cm. Because the

photoexcitation of electrons from traps re- sults in a discharge of the insulator, the

flat-band voltage shifts with exposure to light. The photoyield Y(Kuj) is then obtained

as

Y(Ku>)

Vl AVfb

del (Kid) At

(2)

hw

where I(Koj) is the photon flux at energy Koj and e is the electronic charge. The measurements were made in steps of finite interval

At, between 10 and 60 seconds. The interval

At was chosen to be sufficiently brief so that the total spectrum was not apprecially

distorted by a depletion of the trapped

charge near the end of the measurements se- quence. The discharge measurements were performed in order, from low to high photon

energy

Photodepopulation spectra were obtained from

the decay of Vfj-, , shown in Fig. 2, produced

by a sequence of exposures to light of increasing photon energy. In this case, the

sample was exposed for 15 sec. at each photon

energy, and the charge measurement was made

between the intervals of exposure. Measurements were taken in the dark because any incident light would cause a shift in Vf^ due

190
to a small induced surface photovoltage.

Some portion of the initial trapped charge

is depleted during the course of the measurement, leading to a distortion of the

spectra from this data. 3 Throughout the set

of measurements reported in Fig. 2, the electric field in the Si3N^ was always greater

than 10^ V/cm, which is sufficient to ionize

shallow traps in the bulk of the Si3N^ and

to assure the collection of electrons injected into the conduction band of the Si3N4.

RESULTS

Photoyield spectra obtained from Vfj, discharge curves are shown in Fig. 3 for an initial charge density of 7.8 x lO-^ 2 e /cm2

. Similar spectra were obtained for initial

conditions of 6.6 and 5.1 x lO-*- 2 e /cm2

. The spectra display a parabolic rise above a threshold of about 3.0 eV. The yield reaches

a saturation level at a photon energy of

about 4.0 eV and then falls off. For photon

energies within the 4.25 eV window, the

spectra are relatively simple, characteristic

of emission from a single trap level with an ionization energy of about 3 eV. At photon

energies above the 4.25 eV window, the yield

drops rapidly to a negative value because of

a filling of the traps by photoinj ection

from the silicon substrate. Because of this

photoinj ection , the photoyield spectrum is

valid only for photon energies in the window

below 4.25 eV. The curves above 4.25 eV are shown as a dashed line in the figure.

Below 4.25 eV, the yield spectrum shows no structure other than the parabolic edge near threshold for all values of net negative

charge; the shape of the yield curve is not significantly influenced by the total charge

density. At photon energies near threshold,

a plot of the half power of yield vs photon

energy fits the data quite well, as in Fig.

4, to determine a threshold for ionization

of about 3.0 eV. It is interesting that the

threshold increases from 3.00 eV to only

3.05 eV when the net negative charge is

changed from 5.1 x 10 12 e~/cm2 to 7.8 x 10 12

e /cm2. This suggests that the distribution

of traps is not filled in sequence from low

to high energy, but rather the traps are filled uniformly at all energies of the distribution.

Some information on the photodepopulation

process can be obtained from the decay of

Vfb or stored charge in response to light of

different photon energies. Fig. 5 shows

time dependent discharge curves for 3.5 eV,

4.0 eV, and 4.5 eV. Both the 3.5 eV and 4.0

eV discharge curves decay to the same saturation value, corresponding to a drop in net stored charge from 7.8 x 10 12 e /cm2 to about

3.0 x 10-*-^ e /cm2

. Presumably, saturation

corresponds to the point at which all electrons accessible to ionization by a photon of

energy no) have been removed. Since the satu- ration level is the same for 3.5 and 4.0 eV photon energy, one might conclude either that no significant trap levels lie between 3.5

and 4.0 eV or that any traps lying between

3.5 and 4.0 eV are ionized by a multiple ex- citation. The evidence from the photodepopulation spectra indicates that the former possibility is reasonable; the photoyield saturates at an energy of about 3.8 eV, indicating

that few traps appear between 3.8 eV and the

4.25 eV window. The discharge curve is quite

different for light of 4.5 eV, which lies

above window of photon energy. Here, the Vf]-, decays rapidly to a saturation level which is

higher than that at 3.5 eV or 4.0 eV. This

larger net charge at saturation results from

the operation of two competing processes, the

filling of traps by photoinj ection and the

emptying of them by photoionization.

There is some question about the effect of

charge carrier injection at the Al-Si3N4 interface upon the "window" available for photodepopulation spectroscopy. The photoinjection of holes into the Si3N4 would compete

with photodepopulation in the discharge of

trapped electrons. However, the process of

photoinj ection of holes or electrons from a metal electrode into Si3N4 is very inefficient, and is not thought to be a significant

factor in the spectra presented for Si02~

Si3N4- To test this hypothesis, we measured

the accumulated positive charge photoinj ected

into the insulator from the Al electrode, af- ter all of the accessible electron traps had

been emptied. During the interval of exposure

to light, the gate electrode was maintained

at +10 V, and the silicon surface field was

not nulled. No evidence was found for any

significant photoinj ection of holes from the

aluminum, so that the Al-Si3N4 contact does not seem to limit the available "window".

Injection from the aluminum may, however,

contribute an insignificant component to the

photodepopulation spectra.

The information derived from the photoionization study adds to our knowledge of the Si02~

Si3N4 interface, which is shown schematically

in Fig. 6. The electronic levels are shown

for the specific MNOS structure comprising

Si-Si02-Si3N4~Al. As found in the photoionization study, the electron traps lie in a band beginning 3.0 eV below the Si3N4 con- duction band. The total width of this trap

band is probably less than 0.5 eV. However,

other trap bands may exist below the 4.25 eV limit of the photoionization technique.

Another technique would be necessary to probe

191
the deep traps, with a binding energy greater

than 4.25 eV. No evidence is found in the

photoionization measurements for electrons

in the shallow traps, seen in TSC and I-V measurements of 0.5 eV to 1.5 eV binding

energy, which are characteristic of bulk

Si3N4. Presumably, the shallow traps

throughout the Si3N4 are ionized by the high

electric field present during the measurement. As a result, the traps responsible

for charge storage in MNOS structures are

not the shallow bulk traps, but rather, they

lie at 3.0 eV or more binding energy. With

the generally accepted 1.1 eV discontinuity

in conduction bands between the Si02 and the

Si3N^, the top of the trap distribution is

0.15 eV above the silicon valence band.

This trap energy is somewhat lower than

might be expected. However, the trap is

only approximate because of the uncertainty

in the 1.1 eV step at the Si02~Si3N4 interface. A downward adjustment of the 1.1 eV step would raise the trap distribution with

respect to the silicon.

DISCUSSION OF THE TECHNIQUE

Photopopulation spectroscopy is a simple and

powerful technique for examining deep traps

in insulating thin films. The spectra can be used to determine the photoionization

thresholds of traps within an energy "window",

which is equal to the electrode contact barrier. A principal advantage of the technique, a high sensitivity for examining very

small numbers of traps, results from a measurement of a change in total charge rather

than a current. Extraneous sources of current do not obscure the photodepopulation

spectra for photon energies within the

"window". The technique should be particularly useful in examining traps introduced

into MOS structures by ion implantation, impurities, or radiation damage.

REFERENCES

1. DiStefano, T. H. , and Franz, J. M. , to be published.

2. Harari, E. , and Royce, B. S. H. , J. Appl. Phys. T2, 106 (1973).

3. DiMaria, D. J., and Feigl, F. J., Phys.

Rev. B9, 1874 (1974).

4. Kendall, E. J. M. , Phys. Stat. Sol. 32., 763 (1969).

5. Franz, J. M. , Fitzgerald, D. , and

DiStefano, T. H. , Charge Trapping in

NMOS Capacitors, Extended Abstracts of

the Meeting of the Electrochemical

Society, 145th Meeting, San Francisco,

May 1974.

6. Kendall, E. J. M. , Canadian J. of Phys.

46, 2509 (1968).

7. Swaroop, B., and Shaffer, P. S., J. Phys.

D3, 803 (1970).

8. Lundstrom, K., and Svensson, M.,, IEEE

Trans. Electron devices ED-19 , 826

(1972).

9„ Lundkvist, L. , Lundstrom, I., and

Svensson, C. , Solid State Electronics

16, 811 (1973).

10. Deal, B. E. , Fleming, J. J., and Castron,

P. L. , J. Electrochem. Soc. 115 , 300

(1968).

192
Figure 1. Schematic representation of the

photocapacitance technique which was used to determine the energy distribution of electrons trapped in a Si02~Si3Nit system. A semitransparent aluminum electrode is maintained at a potential V^, for which the

electric field at the silicon surface is zero. The zero field condition is sensed by

a high frequency capacitance measurement and

maintained by a feedback loop. Photons of

energy -ttw excite electrons from traps in the

system to cause a change in the which

can be related to the photoemission yield

for the process.

Figure 2. Spectral dependence of the photon

induced discharge of the trapped electrons

for three different levels of initial charge.

The discharge curves show the decay of the

flat-band voltage caused by 15 second ex- posures to light of increasing photon energy.

5

-

• 7.8 x IO

l2 e"/cm 2 x6.6 x 10

5. I x I0 1

12

4 5

1 1 i

i

h h- J\ \

i i

\

i i i 3

Figure 3. The

yield for photo

Si0 2 -Si 3 Ni+ inte

photon energies

a component due

photoinjection

This injection

in the number o photon energies

spectral dependence of the

emission from traps near the

rface. The photoyield at above hw = 4.25 eV contains

to the filling of traps by

from the silicon electrode,

current causes a net increase

f trapped electrons for

above 4.7 eV.

193
2.8 30 3.2 3.4 3.6 PHOTON ENERGY (eV)

Figure 4. The spectral dependence of photoyield. Near the threshold for photoexcitation, the yield is nearly parabolic with

photon energy above extrapolated thresholds

of 3.00, 3.01, and 3.05 eV for initial

charge densities of 5.1, 6.6 and 7.8 x 10 12 e~ I cm2

.

10V

Figure 5. The decay in Vj^ produced by ex- posure to photons of three different energies. The flat-band voltage relaxes to a saturation value of 3.9 V for a photon en- ergy of 4.25 eV and below. For a photon en- ergy above the 4.25 eV "window", the saturation charge level is somewhat higher and

the behavior more complex because of the

operation of two competing processes: photoionization of traps and trap refilling by

photoinjection from the silicon.

1000

Si SiO4.25 eV

Si,N, 3

IV 4

TIME (sec)

-475 eV

Figure 6. The energy level diagram for the

Si02~Si3Ni+ system showing the energy of the

electron traps near the interface. This

assignment assumes a 1.1 eV discontinuity

in the conduction bands of Si02 -Si3Ni+

. a. (Ref. [12]: B. E. Deal, P. J. Fleming,

and P. L. Castro, J. Eleotrochem. Soo. 115

300 (1968)).

194
DISCUSSION OF THE PAPER is to measure the charge and then differentiate that.

Participant : How will you determine the trap

density?

DiStefano: The density of occupied traps is

determined by working back to get the photoyield as a function of photon energy. From

an assumed ionization cross section you can

get an approximate value for the trap density.

The point is that you can determine the photoemission yield from a measurement of the

photodepopulation without actually measuring

that current directly. I might also point

out that often you cannot measure the photocurrent from the traps directly because the

currents involved are small and obscured by

other currents in the system. The only way

you can detect the trap depopulation current

Participant : This is specifically for a nitride film?

DiStefano : No, this is for the MNOS systems

we have examined. You can, I guess, measure

the current directly in the case of a large

amount of charge trapped in the insulator.

You can look at photocurrent and get a photoconductivity by a beaching of those traps.

But, for the case of a very small number of

traps in insulators like Si02, the discharge

time is so long and the total charge density

is so small it is very hard to detect the

photocurrent directly. Also, with the photocapacitance technique, we measure only the current from the deep traps of interest,

without the complication of extraneous

currents

.

195

NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg , Maryland, April 23-24, 1975 (Issued March 1976)

PH0T0EMISSI0N AND PHOTOVOLTAIC IMAGING OF SEMICONDUCTOR SURFACES*

T. H. Di Stefano

IBM Watson Research Center

Yorktown Heights, New York 10598

ABSTRACT

The new electro-optical techniques of scanned internal photoemission and

scanned surface photovoltage can be used to obtain images of non-uniformities and defects at semiconductor surfaces or interfaces. The spatial

resolution of either method is primarily determined by the wavelength of

the laser used as a collimated light source. The images obtained by the

photoemission technique are replicas of the local contact barrier at the interface between two materials. Typically, photoemission images show

details due to variations of stoichiometry or local concentration of impurities. The photovoltage technique, on the other hand, displays re- combination centers such as precipitates, dislocation, stacking faults,

cracks, and gross damage near a semiconductor surface. Each of these

electro-optical imaging techniques provides information about interfaces

that was previously inaccessible to observation.

INTRODUCTION

Focused beams of various types of ionizing

radiation have been used in the past [1-20]

to probe semiconductor surfaces and devices,

generally in the region near a p-n junction.

The signal induced in a junction by a scanned

optical beam has been found to provide in- formation [3-10] on resistivity fluctuations

in the material, on gross defects in the device, and on surface irregularities or inversion layers near the junction. Generally,

in this type of measurement, a beam of electrons or photons is scanned across an area

of a semiconductor to produce hole-electron

pairs in the material. An irregularity in the material will produce a fluctuation in a sample current of some sort measured when

the beam scans across the defect. In the

standard configuration, a p-n junction collects current produced within a diffusion

length of the junction. Any change in either

the generation or recombination rate near the

junction is detected as a fluctuation of the

collected current. The beam induced current

is displayed on a rastered CRT to produce an image of the defects and inhomogeneities

.

Imaging techniques of the electron beam induced current [11-20] (EBIC) or photon induced current type [4-10] are useful in identifying defects in individual semiconductor devices, but they cannot easily be

* Work supported in part by the Defense

Advanced Research Projects Agency and monitored by AFCRL under contract number

F 19628-74-C-0077.

applied to the study of semiconductor surfaces

or interfaces on which no p-n junction exists.

The necessity of a p-n junction to collect

the beam induced current makes non-destructive

evaluation of material surfaces at various

stages of processing virtually impossible.

Recently, several scanned beam techniques,

including scanned internal photoemission

[21-24] (SIP), scanned surface photovoltage

[25, 26] (SSP) , and the corresponding elec- tron beam analogs [27, 28] have been developed

to probe free semiconductor surfaces. These new techniques provide images of defects near

surfaces or interfaces without the necessity

of a p-n junction to detect the signal current. As a result, they can be used to examine non-destructively the defects which occur

on semiconductors at various stages of material processing. Each of the techniques has

characteristic advantages and capabilities.

The e-beam techniques provide the best reso- lution, better than about 1000 A, [27] but

they are complicated by beam induced charging

of any insulating film on the surface. [28] On the other hand, the optical techniques

have a resolution capability of on the order

of one micron, but they do not charge the surface insulating or passivating film.

Hence, the optical techniques can also be

used to examine charge non-uniformity in the

insulating film on the surface without disturbing the system. Additionally, the optical

scanning apparatus can be made to be cleaner

than existing e-beam systems because it does not require a complex vacuum system.

197
Of the optical techniques, SIP determines

small inhomogeneities in the contact barrier

$ between two materials, one of which is

typically an insulator such as Si02. The

SIP technique is useful in imaging changes

in 3> which are due to contamination, interface chemical reaction, impurity segregation

at the interface, or surface crystallographic

defects. In contrast, the SSP technique is

sensitive to local recombination centers

that are within about one micron of the surface and to charge inhomogeneities in the

insulating overlayer. Typical recombination

centers seen in SSP images include crystallographic defects, precipitates, and gross

surface damage; these defects are usually

introduced during growth, cutting, polishing,

and high temperature processing of the semiconductor material. The SSP technique has a unique advantage in that no direct contact

to the sample is required if the photovoltage measurement is made by capacitive coupling. Since physical contact is unnecessary,

the SSP technique is clean and completely

non-destructive.

OPTICAL SCANNING SYSTEM

The instrumentation for the two scanned

optical techniques is quite similar in that

a small spot is rastered on the sample surface while a signal, either photoemission or photovoltage , is detected and displayed on a CRT. A cw laser was used for the light source in order to obtain a measurable signal from a small, nearly diffraction limited

light spot. Reasonable signals were obtained

with 1.0 mW of light incident on the sample,

without any noticeable thermal damage to the

sample at the 105 W/ cm2 power density at the

focused spot. The spot was swept by the

sawtooth motion of a galvanometer driven

mirror, which was scanned at a 20 Hz rate.

In the other direction, the beam was swept

slowly in order to complete a 600 line

single frame in 30 sec. The frame rate,

which is limited by the rate of the fast scan mirror, could easily be increased an order of magnitude by increasing the fast scan rate to 200 Hz. A further increase, at the expense of a reduced spacial resolution,

could be obtained with an accoustic deflection system for the fast scan direction.

The total resolution capability of the optical scanning technique is limited by three

equally important factors: the divergence

of the light output from the laser, the diffraction limit of the objective lens, and

the bandwidth of the detector. At best, the

spacial resolution (of the SIP measurement)

is estimated to be about 0.43 microns, with

the dominant factors being lens diffraction o at A = 3250 A and laser divergence.

The simple optical system shown in figure 1 was used to scan a light spot over a square

area of the sample surface. Radiation from

a laser source was focused onto a 50 micron

pinhole in order to eliminate the light that

diverged by more than 0.5 mrad. The laser source was chosen to be either an RCA He-Cd

laser operating with 3mW at 3250 A or a Spectra Physics He-Ne laser with 5 mW at 6328 A. Light emergent from the pinhole was projected a distance of one meter onto a series of two front surface mirrors, which

rotate through angles 0 and <j>. Light from

the two scanned mirrors completely filled the

aperture of the objective lens for the full

range of the scanned angles. Then, the

focused spot of light from the objective lens

scanned an area of the sample as the two orthogonal mirrors were rotated. The size of

the scanned area, and to some extent, the

resolution of the system, were determined by

the objective lens. The focal length of the

lenses used in the measurements range from

0.54 cm to 20 cm. Reflecting objective

lenses were used for the high resolution

measurements so that correction of the optics

for the wavelength of each laser was unnecessary

The fast scan mirror was driven at about 20 Hz by a rotary galvanometer, G— 306, available

from General Scanning, Inc. Care was taken

to balance the mirror and rotor in order to minimize parasitic oscillation of the mirror

assembly. This parasitic oscillation or vibration causes deflection of the light

perpendicular to the direction being scanned

by that particular mirror. Although faster

galvanometer scanners are available, the

relatively conservative rate of 20 Hz was chosen to minimize the problems caused by

vibration and to optimize resolution.

The problem of positioning and focusing the

sample, something which is particularly difficult in the case of UV light, was solved by

detecting light reflected back into the system from the sample. Initially, light from

the pinhole goes through the scan system

which focuses it onto the sample. When the

system is in focus, the light reflected from

the sample is focused by the system back onto

the pinhole, independent of the mirror positions 9 and <|>. However, if the system is out

of focus , the reflected light falls outside

the pinhole where it strikes a fluorescent

screen. The system is brought into focus by

adjusting the sample position along the optical axis until all of the reflected light is

refocused back onto the pinhole.

In order to locate the sample in the beam and

to examine the surface conditions during the

198
measurements, a reflectivity image of the

sample can be obtained and displayed in the

coordinate frame of the CRT. A portion of

the light reflected from the sample is directed onto a photocell by a beamsplitter.

The light falling on the photocell comes to focus at a point which is independent of the scan angles 6 and <j>. Current from the photocell, when displayed as Z-axis modulation of

a rastered CRT, forms an image of the sample

at the wavelength of the laser source and in

the same coordinate frame as the SSP or the

SIP image of the same area.

Both the photoemission and the photovoltaic

signals are detected and displayed on a CRT screen by the circuitry shown schematically

in Fig. 2. The signal current is detected

by a Kiethley 18000 picoammeter which has

been modified for a fast response; the modification involves a reduction of the feedback capacitance in each of the two amplifier

sections, while maintaining the stability of

the overall system. The signal in the SIP

measurement is the photocurrent , while the

signal in the SSP measurement is the capacitive current induced in the probe electrode

by the surface photovoltage. The amplitude

and dc level of the signal are adjusted and

applied to the Z-axis of the CRT, along with

a retrace blanking pulse derived from the

fast scanner. The X and Y axes of the CRT are driven by voltages proportional to those

directly across the drive motor armatures.

Images of the Z-axis modulated display of

the SSP or the SIP signal are recorded

photographically

.

The measurement apparatus, shown in figure

3, includes the scan mirrors, the galvanometer motors, the reflecting microscope objective, the sample positioning stage, and the

picoammeter. Light, incident on the mirrors

from the right , is deflected onto the two orthogonally mounted scanning mirrors by a stationary positioning mirror. From the

fast scanned mirror, the beam enters the re- flecting objective lens (Beck 36/. 5) which

focuses the light onto the sample. Because

of the curvature of the focal plane of this

objective, the overall resolution is noticeably degraded for beam deflections greater

than ±0.5°. In order to extend the useable

field of view, it is necessary to use a re- fractive piano-objective which is corrected

for flatness of field at the laser wavelength.

The sample is held in position and focused

by an X-Y-Z adjustable hot stage. Signals

from the sample are detected by the picoammeter located behind the hot stage in figure

3. The entire assembly is secured to a two meter optical bench.

SCANNED INTERNAL PHOTOEMISSION

Scanned internal photoemission (SIP) is a relatively recent technique [21-24] for probing interfaces with a beam of light to produce

electronically an image or replica of the

interface contact barrier. Typical images

display a modulation of the internal photocurrent induced by the scanned beam which is

due to a lateral inhomogeniety of the contact

barrier. The technique is somewhat analogous

to scanning electron microscopy (SEM) , except

that in this case a light beam is used to excite electron emission into a dielectric instead of into vacuum. As in the SEM technique, the emitted current is displayed on the Z-axis of a CRT.

The mechanism of SIP is internal photoemission

on a local scale, as shown in figure 4. A scanned light beam of energy 'hw penetrates a transparent [29] or semi-transparent [30]

electrode and is absorbed in the opposite

electrode where it excites electrons into the

conduction band. Some of the electrons have

sufficient momentum normal to the interface

so that they are able to surmount the barrier

$ and reach the opposite electrode. Near

threshold, the quantum yield for photoemission is

Y (-nu

where y is a constant and a is an empirically

determined power, typically in the range 2 to

3. A small percentage reduction in 0 leads

to a relatively larger increase in yield,

d[ln ¥] =

(fito - $)

d[ln $]

The sensitivity of the SIP technique depends

upon the rapid increase of photocurrent produced by a small decrease of $ below the photon energy <hw. For a reasonable image con- trast, the laser and any Schottky barrier re- duction are chosen so that /fiw is several

tenths of an eV above $.

The Si-Si02 interface is an interesting

candidate for examination by SIP imaging, as can be seen in figure 5 and 6. Here, the

Si-Si02 interfaces have been uniformly coated

with about 4 x 10^ Na/cm2 by electrodeposition in order to reduce the interface barrier^l from 4.25 eV down to a level that is

accessible to the He-Cd laser operating at 3.81 eV. Actually, the sodium acts as a "staining agent," increasing the local photocurrent in those places where ions accumulate

on the interface. Since the sodium coverage

is influenced by defects, it was found to en- hance the SIP contrast around defects. Fig.

5 shows an interesting pattern of dark spots

199
seen in the emission from an Si-SiC>2 interface

grown on vLO20 P/cm3 doped (100) Si. These

spots, which are about 10 microns in diameter, have been identified as phosphorous

that has segregated onto the Si-Si02 interface during high temperature growth [32, 33]

of the 1000 A thick Si02- Apparently,

sodium that does reach the Si-Si02 interface

is rendered ineffective by chemical combination with the phosphorous lying within a few

atomic layers of the interface [33]. These

phosphorus islands, observed for the first

time by the SIP imaging technique, were

found to be correlated with the phosphorous

doping density in the silicon substrate.

Another type of defect on the Si-Si02 interface is seen in figure 6 as a line, about 30 microns long, with bright spots at either

end. The defect, which lies along the <110>

crystallographic direction, is thought to be

a stacking fault that has been decorated by

sodium ions. The contrast of the stacking

fault is enhanced by the sodium decoration.

The SIP technique was also used to study

metal-insulator interfaces such as Nb205~Bi,

which is of interest because of the fast

switching found in this structure. The SIP

image in figure 7 represents photocurrent

emitted from a 200 A Bi electrode into

Nb205 • The light source was a He-Ne laser

at 1.96 eV, which is close to threshold

where the technique is most sensitive. In

this system, the contrast mechanism is not simple because of the large number of traps

in the ^205. ' S'

SCANNED SURFACE PHOTOVOLTAGE

Photovoltage images show defects that cause minority carrier recombination on a semiconductor surface. The SSP images have also

been found to show any non-uniform charge in

the passivating layer on a semiconductor.

An SSP image is formed by displaying the

first or second derivative of the photovoltage induced by a scanned beam of light on a semiconducting surface. The travelling

light spot produces a comet shaped cloud of

minority carriers which are captured on the

slightly depleted surface, as represented in

figure 8. Carriers from the moving cloud cover the surface to produce a photovoltage.

When the beam passes over a defect, some of

the charge cloud is lost to recombination,

leading to a dip in the photovoltage. This

dip is detected by a capacitively coupled

electrode near the surface, and displayed on

a CRT. Since the charge cloud is quite

large, on the order of the diffusion length,

the resolution of images produced directly

from the photovoltage is rather poor. The

high resolution [25] of the SSP technique is

due to the use derivative of

image. The di

signal in the

where the char

rithmic singul

derivative is

the capacitive

tions of the s

of the first or second

the photovoltage to form the

f f erentiation accentuates the

region around the light spot,

ge cloud density has a logaarity. In practice, the first

obtained directly by measuring

current induced by fluctuaurface photovoltage.

By using conduction of minority carriers

along the slightly depleted surface, it is

possible to obtain SSP images at some distance from the electrode. The local photovoltage is coupled to the entire surface by

the slightly conducting surface. In some cases , it is necessary to enhance the surface

conductivity by generating extra minority

carriers by a dc blanket illumination. The

resolution of remotely detected images was

found to be reduced if the surface conductivity is too low. Also, in order to develop a photovoltage signal, the surface must be depleted, either by a blanket electrode (the

electrode can be an aquous solution, as was shown by Lile and Davis [26]) , or by a surface treatment such as that produced by dry

oxidation of p-type silicon.

The signal S' is proportional to a convolution

of the excess local recombination rate R with

an experimental sampling function, which is

approximately equal to the x-derivative of

the charge cloud density n. Mathematically,

dn(x-x ,y-y ) (x - R(x,y) •V dx

where G is the total generation rate, S is tne effective depletion width, Ea is the permittivity of the semiconductor, E^ is that of

the dielectric, and t 0 is the surface minority

carrier lifetime. Ideally, dn/dx would be a singly differentiated Dirac delta function

and the signal S' would then be proportional

to dR/dx. Because of the finite size of the

light spot, however, the resolution is re- duced. In a coordinate system centered on the moving spot, a modified diffusion equation

[25] is solved to obtain the charge density n in the region around the scanned spot. The

charge density

n = A exp (Kr)

where

of

V T o\l/2

4 ? )

'

and where KQ is a modified Bessel function

the second kind, and D is the surface diffusion coefficient. The sampling function

dn/dx is plotted in figure 9 for several scan velocities, with the reasonable parameters

200
D = 15 cm2 /sec and x 0 = 0.01 sec. The form

of dn/dx deviates little from 1/r near the

light center at r = 0 . Based on this simple

theory, the resolution at 1/e is approximately 2.7 times the light spot diameter for

singly differentiated SSP and nearly 1.0 for

the doubly differentiated technique.

Several examples of SSP images, obtained

from (100) silicon surfaces, display the

resolution, sensitivity, and comparative advantages of this relatively simple technique.

An SSP image of several clusters of emergent

dislocations which penetrate the (100) surface of a sample are shown in Fig. 10. The

dislocations were produced by mechanically

damaging the reverse side of a silicon wafer,

and then annealing it in dry He at 800 °C for

30 min. It is known^^ that such damage will

cause dislocations to propagate through to the front surface during high temperature

processing. Surface defects were also found

in SSP images of nominally defect free silicon, an example of which is shown in figure

11. Here, a semitransparent blanket electrode was used in order to maximize resolution. The image shows scratches, some as long as 1000 microns, which are not parallel

to a low order crystallographic direction.

Presumably, these scratches were introduced

during the polishing process. Another type

of defect, commonly seen in device grade

silicon, occurs throughout the area of fig- ure 12. These small spots are thought to be crystallographic defects associated with

precipitates of oxide or other foreign material. The density of these defects was

found to be correlated with oxygen concentration, suggesting that the defect may be

associated with an oxide precipitate. The

density was found to be as high as 10^ / cm2 (within several microns of the surface) . A rather uncommon surface defect, displayed in

figure 13, was found on a device-grade silicon wafer. It appears to be similar to the

surface defects, identified as microsplits

by Schwuttke [35], which are small cracks

along the <110> direction in the silicon

surface, with stacking faults at either end.

Schwuttke has found that microsplits like

that in figure 13 are introduced during the saw cutting of the silicon.

It is interesting to compare the SSP techniquwith the useful but destructive technique of

selective etching. A sample was prepared

[36] for this purpose by polishing a piece

of saw cut silicon at a 2° bevel off normal

to the (100) surface. The sample geometry

is represented in figure 14. A Sirtl etch was applied to half of the surface to reveal

the saw damage. The sample was then oxidized

and the series of SSP images in figures 14

A

1 - D' were measured. For comparison, figures

14 A - D are micrographs of the etched surface

at points of polish depth equal to that of

the equivalent SSP image. Both the etch pattern and the non-destructive SSP technique

appear to provide the same type of information. However, it is thought that the SSP

measurement is sensitive to sub-surface

structure such as precipitates which may not appear in an etch pattern.

CONCLUSION

Each of the scanned optical techniques,

scanned internal photoemission and scanned

surface photovoltage , has unique capabilities

for examining semiconductor surfaces and interfaces. Photoemission imaging is useful in

determining interface reaction or contamination. However, it is not amenable to wide

use in process control or material evaluation

because of the difficulty of the measurement

and the requirement of a contacting counterelectrode. On the other hand, photovoltage

imaging is a simple, fast, non-contacting,

non-destructive technique that offers reasonalbe resolution, about half a micron for a double differentiated measurement at a wavelength of 3250 A. Since scan speed does not degrade resolution, there is no inherent

physical limit on the frame repetition rate. The photovoltage technique has been used

routinely to detect stacking faults, dislocations, precipitates, gross surface damage,

and local insulator charging.

The photovoltage imaging technique SSP has

great potential for use in semiconductor

material evaluation and in process control in

the production of semiconductor devices. By

observing the defects introduced during the

various stages of production, it is possible

to identify and control the factors that

introduce defects. In the particular case of

silicon integrated circuits, both the production and the reliability of the devices is

critically dependent upon crystalline perfection of the material which can be monitored

by SSP imaging. The use of Photovoltage

imaging is somewhat less useful as a research

tool than the electron beam analog, as developed by Bottoms [28] , because the special

resolution is considerably lower.

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33. 16. Foss, N. A., J. Appl. Phys. 41, 823

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19. Gonzales, J. A., Scanning Electron

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(ITT Research Institute, Chicago,

Illinois, 1974).

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Fujikawa, K. , J. Appl. Phys. 46, 2288

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Scanning Surface Photovoltage Study of

Defects in Silicon, Proc. 13th Annual

Reliability Physics Conf. Las Vegas,

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The bevel polished sample was prepared

by J. W. Philbrick, IBM Hopewell Junction,

New York.

202
SCANNING

MIRRORS BEAM

SPLITTER

LASER

Figure 1. Optical system and measurement apparatus for the scanned optical measurements.

Laser radiation is focused through a 50 urn pinhole and then deflected by two scanning

mirrors rotating about perpendicular axes. The deflected light fills the aperture of a microscope objective lens for all deflection angles 6 and <f> . The system is brought to focus by adjusting the position of the sample along the optical axis to minimize the size

of the light spot reflected onto the fluorescent screen on the back of the 50-um pinhole.

A portion of the reflected beam is diverted and brought to a stationary focus on a photocell

in order to detect a reflectivity image of the scanned area.

LASER

X-Y-Z PHOTO

EMISSION

DISPLAY

YX

Z

X-Y-Z REFLECTIVITY DISPLAY Y

X

Z

Figure 2. A schematic representation of the scanned optical system showing the instrumentation required to obtain SIP or SSP images. The x and y inputs to the CRT are driven

by the voltage applied to the mirror galvanometer motors.

203
Figure 3. The optical scanning apparatus showing the scan mirrors, galvanometer drivers,

microscope reflecting objective, sample positioning stage, and the preamplifier head.

\W\VA\

Figure 4. A representation of the processes involved in scanned internal photoemission

A small reduction in the effective interface barrier allows an enhanced photocurrent to flow.

204
Figure 5. Scanning internal photoemission

images of three samples of Si-SiC^ interfaces covered with 4 * 10 12 sodium atoms per

square centimeter, and measured at a photon

energy of 3250 A. The light areas indicate

a high photoyield produced by the presence

of sodium on the interface. The dark areas

are thought to be due to regions of phosphorus rich layers near the Si-Si02 interface. Phosphorus segregates at that interface during the formation of the Si02 layers,

produced by the oxidation of 10 20 P/cm^

doped silicon.

. 6

.

Figure 6. Scanned internal photoemission

measurements of an Si-Si02 interface which

has been uniformly coated with 4 x 10 12 Na/cm2

. The Na is found to decorate a

30 pm, long defect which lies in the <100>

direction on the silicon surface.

Figure 7. Scanned internal photoemission

image of a Bi-Nb20 interface, measured at

A = 6328 K. The vertical striae are due to

a dc level shift of leakage currents.

205
S'(t)

Figure 8. The rudiments of the scanned surface photovoltage measurement. A spot of light

is swept across a semiconductor surface, generating a comet shaped cloud of minority

carriers in a surface depletion region. The surface potential, capacitively detected by a remote electrode, fluctuates when the spot traverses a surface recombination center.

Figure 9. The spacial derivative of the surface minority carrier density, along a line in

the direction of travel of a 1 pra spot of light. The derivative is relatively insensitive

to the scan velocity v for the values shown in the figure. The quantity [dn/dx| is the

sampling function in the singly differentiated scanned surface photovoltage measurement.

206
Figure 10. Scanned surface photovoltage

image of several clusters of emergent dislocations through the (100) surface of a silicon wafer. The dislocations were generated by sandblasting a spot on the reverse

side. During subsequent annealing and oxidation steps, the dislocation loops generated at the damage sites grow through the

wafer, eventually penetrating the front

surface.

Figure 12. Scanned surface photovoltage

image of a (100) silicon surface showing

small, randomly distributed defects.

Figure 11. Scanned surface photovoltage

image of surface damage produced during th

cutting and polishing of a silicon wafer.

The damage lines do not follow the <110>

direction on the (100) surface.

i

Figure 13. Scanned surface photovoltage

image of a defect, apparently a microsplit

which was introduced during saw cutting of

the (100) surface. The microsplit is

thought to be flanked by stacking fault

arrays in the <110> direction on the

surface

.

207
A A'

Figure 14. Scanned surface photovoltage images (A', B', C', D') of saw damage in a silicon

wafer which has been polished at a 2° angle. For comparison, micrographs (A, B, C, D,) of

an adjacent area of the surface were made after etching. Both sets of images, taken on a highly polished area of the beveled surface, show the transition from the damaged region to the undamaged region occuring at about the same depth, as is observed on an etched portion

of the bevel.

208
DISCUSSION OF THE PAPER

Feigel : Photovoltaic imaging is a non- contacting technique, right?

DiStef ano : That is correct. The only

contact that is made to the sample is the

capacitive contact. The sample is simply a sandwich between capacitive electrodes.

Feigel : Then is it relatively easy to engineer?

DiStef ano : Nothing is hidden, everything

was shown on the figure.

Feigel : I meant there are no super critical

dimensions or dimensional stability problems you have to worry about?

DiStef ano : Focussing is a problem but, as I mentioned, we use the back reflected light

to focus the sample. Focussing is a problem

in the ultraviolet because you cannot see

the spot. Other than that, there is nothing

intrinsically difficult about the technique.

Thomas : Can you actually identify the

collecting species? You said it is actually

copper.

DiStef ano : In some of the cases we intentionally put copper into the system.

It is an inference that the precipitate is

copper, since there is copper in this sample.

Thomas : We also have found copper, but

there are other metals that can actually

decorate the defects. Would you say there

is some similarity between this technique

and the EBIC technique?

DiStef ano : There are some similarities.

With the electron beam induced current technique you need a junction, and you look at junction current. There is no problem with

resolution, but it is necessary to have a surface junction on the sample. One of the

beauties of the photovoltage technique is that it is nondestructive, so you can look

at an area, process that area, and then go back and examine what defects were caused by

the processing. We are studying various

things such as nitride passivation and the

defects introduced by the mechanical stress. We are looking at deep etching and oxide

isolation in silicon. So this technique is

ideal for following a piece of silicon

through a stage of processing steps.

Participant : Can you use the amplitude of

the surface photovoltage to obtain a threedimensional representation of your data

similar to the Y amplitude images?

DiStef ano : We can do that, but we do not,

because that tends to make it harder to identify the position of the defect. We do use a series of photographs in different colors to look at the depth of the trap. There is one thing that I have not mentioned, namely you

can also bend the bands of the silicon surface and submerge a trap or a surface defect

below the energy level so you no longer see

it. You can take a series of images as a function of the Fermi level at the surface as defect disappears to obtain an X, Y, and

energy representation of the defect.

209

NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices

, held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

ELECTRON BEAM INDUCED IMAGING OF SILICON SURFACES*

W. R. Bottoms

Department of Electrical Engineering

Princeton University

Princeton, New Jersey 08540

Modern electron optics can provide a focused

beam of less than 100 angstrom units in diameter which may be swept across a sample

in order to address any portion within the

"field of view." This allows one to input

energy to a well defined region and, by collecting the signals associated with the various energy relaxation processes which occur, an analysis of the structure of the sample

with high spatial resolution can be undertaken. Figure 1 illustrates the primary

signal channels which are activated by the

electron beam with the principle type of

analysis associated with each beam. The en- ergy dissipated within the sample in terms

of electron-hole pair production can be detected in an external circuit and images,

associated with any process which alters the

current measured in the external circuit,

can be formed. The contrast mechanisms

which are responsible for the images formed

are listed in Table l

1

. It is often difficult to determine which mechanism or which

combination of mechanisms are responsible

for the formation of a given image and, in an effort to deconvolute the various contributions to contrast, data is gathered from

as many different sources as possible. Both

C-V and G-V measurements are conducted to determine as fully as possible the effect of

the incident beam and the various sample

processing steps on the distribution of internal charge.

Electron beam induced imaging is carried out

as a function of the incident beam voltage

which can be related to the depth within the

sample by the calculated range of the incident electron beam2

, Figure 2. Potential

barriers of known depth within the sample

also provide depth resolution since they

contribute to the signal current by separating excess carriers generated by the incident beam. In addition, images are formed

as a function of bias across the sample and

processing history (principally the bias

thermal stressing used to determine the num- ber of mobile ions present and move them

from one region of the sample to another)

.

Work supported by ARPA Contract No. F19678-

72-C-0298.

21

The electron beam induced images of "bare"

silicon surfaces and internal interfaces in

MOS samples are analogous to the corresponding

images which can be obtained using electromagnetic radiation as the exciting probe with

two very important differences. The optical

probe can be chosen so that the incoming beam

does not see portions of the sample which are

of little interest in a particular image 3 (i.e., the silicon-silicon dioxide interface

can be probed with an optical beam of energy

such that a signal is stimulated by internal

photoemission at the interface but the oxide

is transparent to the probing beam) . In the case of electron beam excitation, it is possible to obtain images of interface structure

associated with the excitation of secondary

electrons across the potential barrier between

the silicon and the oxide, as in internal

photoemission, but this signal includes cur- rents associated with all other interactions

of the beam with the sample 1*. The various

current components are convoluted and indeed,

under most conditions, another imaging mechanism will dominate resulting in the reverse

of the contrast obtained from the internal

photoemission mechanism 5

. The result of this

propensity of the electron to undergo energy

exchange interactions which appear continuous

rather than discrete (insofar as the scale of

energy interactions of the experiment is con- cerned) limits the technique. Although the

electron beam interacts much more strongly

with the system under investigation and

therefore is able to "see" events which

could not be seen with optical or X-ray

probes, this very density of information

makes it difficult to deconvolute the contributions from various processes. The second

important difference between the optical probe

and the electron beam probe is that it is

possible to generate an incident electron

beam with a diameter more than an order of

magnitude smaller than that of the optical

probe and therefore, under many imaging conditions, provide a much higher spatial reso- lution with a similar improvement in the

depth of field.

An example of the first of the contrast mechanisms listed in Table 1 is shown in Figure

3. This is the normal image which is derived

from the scanning electron microscope (SEM)

and it is a map of the topography of the

1
sample as well as the variations in the

plane of the secondary emission and reflection coefficients. An appropriate adjustment of incident beam parameters can give

the situation where electrons of low energy

are excited across the barrier between the

gate metallization and the insulator in a MOS sample thus providing an image of the

variations in the barrier height between the

metal and the insulator. An example of such

an image is shown in Figure 4. The structure has been attributed to local variations

in barrier height due to mobile ion aggregation at the gate which was generated by

bias thermal stressing of the sample. Figure 5 is an example of an image of the SiSi02 interface. Variations in barrier

height due to aggregation of Na ions at the

interface are responsible for the structures

in this image.

Figure 6 gives an example of image structure

due to lateral nonuniformities in the charge

stored in the insulator. In this case charge was stored selectively in some parts

of the sample by the electron beam and the

regions of bright and dark intensity with

respect to the background showed that both

positive and negative charge can be stored

in the sample with the incident beam. This

stored charge is an ever present artifact of

the electron beam imaging technique and must

be kept in mind when interpreting images

formed in this manner.

CONCLUSIONS

The use of electrom beam induced images has

been shown to be a powerful tool for the

observation of subsurface and surface variations of electronic structure in silicon.

Spatial resolution of the order of 1000 A

has been obtained for structure at an interface located several thousand angstroms below

the surface. Structure located closer to the

entrance can be imaged with even higher res- olution since there is less spread of the incident beam. The electron beam can produce

the analog of scanning internal photoemission

and surface photovoltage images with better

spatial resolution than is obtained for similar images induced by electromagnetic radiation. The major disadvantages of electron

beam techniques are of cost and of the difficulty in deconvoluting the information obtained. We are, in most cases, only able to present models for the structural images

which are consistent with our experiments.

We are not able, at present, to make a unique

determination of the mechanism responsible

for a given image and therefore we cannot

uniquely identify the structure. The ability

to obtain images of defects on "bare" silicon

has been demonstrated. This has important

implications for developing technologies

which place more rigid specifications on the

starting semiconductor material such as high

voltage power devices and large area CCD

devices

When defect structures in a silicon sample

cause local variations in recombination

rates we can detect a signal due only to the

change in the capacitance of the sample induced by the spatial variation of the

charge. 1 This gives rise to the electron

induced analog of the surface photovoltage

mechanism5

. Figure 7 shows stacking fault

defects in a Si (111) crystal where the

image formation has been attributed to this

analog of the surface photovoltage technique.

This image is of the Si-Si02 interface in an MOS capacitor. Figure 8 is an image of a "bare" silicon wafer which was mounted in

the microscope with spring tension holding

it between two flat metal plates. The natural oxide covering the silicon and the

metal plates is of sufficient thickness to prevent large dc current flow across the

sample and the ac signal is coupled out through the capacitor between the silicon

and the top metal electrode. In this con- figuration the image is again attributed to the analog of the surface photovoltage mechanism.

References

1. Bottoms, W. R., and Roitman, P., Critical

Reviews in Solid State Sciences (in

press)

2. Everhart, T. E.,and Hoff, P. H. , J. Appl.

Phys. 42, 5837 (1971).

3. DiStefano, T. H. , Appl. Phys. Lett. 19_, 280 (1971).

4. Bottoms, W. R., and Guterman, D., J. Vac. Sci. Tech. 11, 965 (1974).

5. Bottoms, W. R. , Guterman, D. , and

Roitman, P., J. Vac. Sci. Tech. 12_, 134

(1975).

6. DiStefano, T. H. , Photoemission and

Photovoltaic Imaging of Semiconductor

Surfaces, ARPA/NBS Workshop IV, Surface

Analysis for Silicon Devices,

A. G. Lieberman, Ed., NBS Spec. Publ.

400-23 (March 1976)

212
TABLE 1

SOURCES OF CONTRAST FOR ELECTRON BEAM INDUCED CURRENT

IMAGES OF INTERNAL INTERFACES IN MIS STRUCTURES

* Variations in secondary emission and reflection coefficients at the sample/vacuum interface.

* Variations in the barrier height at the metal/insulator and

insulator/semiconductor interfaces.

* Variation of the electric field within the insulating region

resulting from interface barrier height changes.

* Local variations in the recombination/generation rates.

* Lateral nonuniformities in the charge stored within the sample.

* Local field variations resulting from topographic structure

at the interface.

* Interface topography resulting in variation of the total energy

dissipation within a given region.

213
SCANNING ELECTRON MICROSCOPE

MODES OF OPERATION

INTERNAL MICROSTRUCTURE

Figure 1, A schematic diagram illustrating the primary signal channels activated by an electron beam. The principle information contained in each signal is noted.

214
2 3 4 5 6 7 8

RANGE (1000 A)

Figure 2. The rate of energy loss for electron beam penetration of Al-SiO -Si systems

calculated using the Everhart-Hof f function.

Figure 3. A conventional secondary electron

image of a region on a MOS capacitor where a "self-healing" dielectric breakdown has occurred. The bright areas correspond to gate

metallization not evaporated during breakdown.

215
Figure 4. Interface image of the metal/

insulator interface taken with a 2 Kev primary beam after BTS processing to drive

mobile ions to the gate.

Figure 5. Interface image of the insulator/

semiconductor interface taken on the same capacitor as in Figure 4 after additional

BTS processing to drive the mobile ions to the insulator/semiconductor interface. This

image was formed with a primary beam energy

of 6 Kev.

itor with the same imaging conditions as in

Figure 4 after it had been subjected to non- uniform electron beam irradiation. Image

detail due to the resulting nonunif ormity

in stored charge is indicated.

216
Figure 7. An image of the insulator/semiconductor interface of an MOS capacitor

taken with a 30 Kev primary beam. The

defects visible with definite crystallographic orientation are attributed to stacking faults.

Figure 8. Electron beam induced images of a "bare" (111) oriented, 150 fi-cm, p-type

silicon wafer which has been mechanically polished. a) analog of the surface photovoltage

image showing residual polishing damage on the wafer. b) conventional secondary electron

image of the same area of the sample showing the topography of the surface.

217
DISCUSSION OF THE PAPER

Participant : Do you see exactly the same

sort of phenomena in Si02 that Tom DiStefano

had mentioned earlier regarding the use of

photoemission and photovoltaic imaging?

Bottoms : Yes, we do see exactly the same

things. For instance, there seems to be a correlation between those stacking fault defects that occur after processing and oxygen

content as well as some metal content in the

material. There is no disagreement in our

data. Tom has more data that is tied to a

production line than I do . I do not have a production line.

Participant : Do you share his inferences?

Bottoms : I am not sure of that. They are consistent with everything we know. We have

samples not from IBM processes, but where

the chemical processes were altered and the

alteration in the images related to the chemical process alterations in the way consistent with everything Tom was talking about. We only have consistency but not specificity.

218
NBS Special Publication 400-23, ARPA/NBS Workshop IV, Surface Analysis for Silicon Devices,

held at NBS, Gaithersburg, Maryland, April 23-24, 1975 (Issued March 1976)

A COMPARISON OF THE TECHNIQUES FOR SILICON SURFACE ANALYSIS*

Charles A. Evans, Jr.

Materials Research Laboratory

University of Illinois

Urbana, Illinois 61801

This paper will summarize the five chemical

analytical techniques that are most widely

used to characterize surfaces and thin films

in the silicon technology area. These are Auger Electron Spectroscopy (AES), Electron

Spectroscopy for Chemical Analysis (ESCA)

,

Ion Scattering Spectrometry (ISS), MeV on Backscattering Spectrometry (BS), and Sec- ondary Ion Mass Spectrometry (SIMS) commonly

referred to as the ion probe.

Because of space limitations, concepts rather

than specifics will be discussed. As such,

there will be exceptions to these generalizations. Sensitivities may be worse or may

be better by a factor of two or three in cer- tain circumstances. The most important con- clusion to make is that every technique has

its advantages and its limitations in a spe- cific analytical feature. More importantly,

the solution to a particular analytical problem may require the application of more than

one of these techniques. So it is important

to see how these techniques interdepend, interrelate, and compliment each other.

The first question that is generally asked

about a technique is "What is its elemental

sensitivity?" In Table 1 the five techniques are characterized in four features of elemental sensitivity:

1. Coverage - All of the techniques can detect from lithium to uranium (i.e.,

almost full periodic coverage). In addition the SIMS technique can detect hydrogen. This is a definite advantage.

In terms of silicon technology, it could

be very important to determine the role

of hydrogen which may be coming from CVD processes, what is happening to it and

how does it enhance or reduce diffusion,

reactions, etc., in a thin film.

2. Specificity - This is the ability of the technique to resolve different elements,

either as it is limited by the intrinsic nature of the technique or in terms of spectral interferences which can be en- countered and reduce the ability to see

a given element. Auger electron spectrometry has a reasonable specificity;

intrinsically it can detect every element

but hydrogen and helium. However, specif

i

interference situations are encountered.

Backscattering spectrometry specificity is definitely dependent upon the mass of the

element. At low mass this technique provides good resolution, however, resolution

decreases as one goes to higher mass. It becomes difficult to tell tungsten from

gold in certain circumstances. ESCA has

good specificity much as the AES technique

Ion scattering spectrometry exhibits much

the same situation as MeV backscattering.

As long as the probing particle mass is of the order of the target particle, good res olution is obtained. However, as the target atom mass greatly exceeds that of the probing ion, spectral resolution decreases

The SIMS technique has very good specificity. In particular, it gives isotopic detection capability which might be useful

in self-diffusion studies.

3. Sensitivity - Sensitivity variations are very important in considering a technique

since this affects the ability to relate

the analytical signal to the concentration

of that element. The two electron techniques, AES and ESCA, have only a factor

of 10 variation from the least sensitive

to the most sensitive element with perhaps

a little bit more variation in ESCA. In BS there is a pronounced, but predictable,

variation in sensitivity with atomic number. From bismuth to oxygen, a factor of about a hundred in sensitivity is lost. Over this same range, ISS suffers a sen- sitivity variation of about a factor of

five. The SIMS technique has the largest

variation in elemental sensitivity. Four

to five orders of magnitude in variation

of sensitivity is encountered between the least sensitive and the most sensitive

element

.

4. Detection limits - Auger electron

spectrometry sensitivity is about 10 3 atomic fraction. Large variations in detection limits are encountered in

backscattering spectrometry between

lithium and uranium. The ESCA technique

has average detection limits between 10~ i

^Supported in part by National Science Foundation Grants OTS-74-05745 and DMR-72-03026

.

219
and 10 atomic fraction. Ion scattering

spectrometry has an average detection

limit of 10~.J vary from 10

SIMS_detection limits

to 10 8 atomic fraction.

Figure 1 summarizes the detection limits in an interference free situation for the different techniques. This figure shows that

there is a periodic dependence of detection

limit on atomic number in the SIMS technique.

It should be noted that these detection limits are for positive secondary ion spectrometry. Negative ion spectrometry with Cs+

ion bombardment provides a reversal of these curves such that many of the less sensitive

elements become the most sensitive. Backscattering spectrometry has a smooth curve which provides detection limits approximately

10 k in the case of bismuth. Auger electron

spectrometry is essentially atomic number

independent. Thus AES detection limits are best related as a broad band across the periodic table with about 10 3 to 10 2 atomic

fraction for detection limits. ESCA exhibits

a periodicity whereas ion scattering has a smooth curve much the same shape as high energy backscattering but with somewhat poorer

detection limits.

There are two important features to these

data. Doping levels in semiconductors are at the 10 3 atom fraction (

r»10 19 atoms/cm 3

) and

go to lower concentrations. Therefore, if

the analyst is to look at doping levels there

is really only one technique that is able to provide elemental identification and chemical

concentration - the SIMS technique. Secondly, it is commonly assumed SIMS is quite subject to spectral interferences. However, the

technique has tremendous sensitivity for many

of the elements, and interferences do not become a problem above the 10 3 level. So the

possibility for spectral interferences in

SIMS does not generally occur until the con- centration is below the detection limits of

the other techniques. More recently, high

mass resolution SIMS instruments have become

available, and even in interference situations very good detection limits can be re- alized. As one goes to high mass resolution,

these doping level detection limits can still

be realized.

Another important aspect of any analytical

technique is the matrix effects that are en- countered, and these are summarized in Table

2. As has been discussed during the whole

conference, AES is subject to matrix or chemical shifts. These generally occur only

for a few elements, and they have become

quite useful in characterizing certain types

of compounds. Backscattering is a nuclear

process and exhibits no matrix effects. In

ESCA chemical shifts are the justification

for the technique. So chemical effects and

matrix effects are very important in ESCA and

are the key to the technique. ISS does exhibit

some matrix effects. This dependence on the

matrix results from the matrix influence on ion neutralization. Differences in intensities

are encountered for the same concentration when

going from one matrix to another because of

probing ion neutralization. In SIMS, matrix

effects are quite often found. Researchers are trying to use them to understand compound presence or absence in materials, but at the present time they are generally considered a nuisance

Since the amount of sample for analysis may

be limited, the alteration or consumption of

material during analysis can be important. As seen in Table 3, the only truly nondestructive

technique for depth profiling is high energy

backscattering. It is non consumptive in all

modes. However, radiation damage of organic

material can result. If channeling experiments

are done to determine radiation damage during

an implantation, interstitial versus lattice

locations in silicon, etc., there can be alteration of the sample with respect to these aspects upon extended exposure of the sample to the probing beam. The remainder of the techniques break down into two areas. The first area includes those techniques which are non- destructive during the analytical process,

i.e., the AES and ESCA techniques, but must use ion etching in order to perform depth profiles. Again, there must be concern about

radiation damage; electron impact desorption

can be quite important with AES. ESCA is less

subject to radiation damage in the analytical

mode when sensitive materials are examined.

In the second area, ISS and SIMS consume sample material in all modes. Even though helium

ions may be used for low sputtering rates during the analytical step, there will still be some material sputtered, and one must be aware

of this at all times

Now, combining all of the above factors, how

does one make these techniques quantitative?

As seen in Table 4, none of the techniques except backscattering are quantitative on an absolute basis. Only for the backscattering

technique can a sample be analyzed using no standards. From the data, quantitation in

both concentration and in the depth scale can

be provided. There is no dependence on the

sputtering rates. The use of standards in BS can improve quantitation from about 5% to about the 2% level. Standards are required by

all the other techniques: AES, ESCA, ISS, and

SIMS. Obviously, one would hope to remove the

requirement for standards and start using some

sort of mechanistic or fundamental approach in

220
order to quantitate the other four techniques.

Quantitation procedures are just beginning in

Auger. The author is not familiar with any

quantitation or mechanistic approach to quantitation in ESCA. In the ISS technique,

known scattering cross sections for two different elements can be used to determine relative concentration by the application of

these known scattering cross sections. In SIMS, some work is being done in terms of

mechanistic approaches. The most successful

process gives ± 200% (that is, a factor of 2)

in almost every case and in many cases gives

quantitation in the + 20% range.

As was implied above, another important point

of quantitation is the ability to quantitatively locate a particular subsurface feature

within the sample. Only the backscattering

technique can provide this on an absolute

basis. As an example, Fig. 2 shows the energy

loss observed versus the thickness as calibrated by a variety of other techniques for a tantalum oxide matrix [1]. The points shown

are actual data points, but the curves are not best fits. These are theoretical curves - two with a depth correction and one assuming

just a straight linear approximation. Two different densities were used since the density is required. As is seen, the theoretical calculation fits very well with the data

points for a variety of tantalum oxide films

as thick as 5,000 angstroms. This is a very,

very important aspect of the backscattering

technique. It does provide absolute calibration of the thickness of the layers in a multilayer film. The as-deposited thickness

of a film may be known, but the annealed

structure can only be characterized by one

of these analytical techniques. Only backscattering spectrometry can provide this

information without need for sputtering rate

calibration.

It is very important in actual device

characterization and in debugging a product

line to be able to perform chemical analyses

with high lateral resolution (< 10 urn) . Those

techniques that provide microanalysis are the

AES technique with a microf ocussed primary

electron beam, the SIMS technique with a microf ocussed ion beam, i.e., the ion microprobe, or SIMS with a stigmatic mass

spectrometer, i.e., the ion microscope.

"Milli-analysis" is now available with ISS using 0.1 mm diameter probing ion beams.

Macroscopic area analysis is provided by the

backscattering technique since a one millimeter probing beam is generally used. There

is one instrument available with a one to ten micrometer beam for backscattering [2]. It is a highly specialized instrument and

very time consuming in its operation. ESCA

usually illuminates rather large areas. People do work with masks to reduce X-ray

illumination areas in order to obtain local

analysis, but it still remains a macroscopic

technique. Table 5 summarizes the lateral

analytical features of these five surface

analytical techniques.

Examination of the two microanaly tical techniques reveals that there is a very intimate

interrelationship between the detection limit

that can be realized and the probe, or analytical, area being used as illustrated in Fig.

3. The detection limits for Auger electron

spectrometry decrease as a smaller probe is

used. This results from a smaller exciting

current available at the smaller diameters.

In this instance, the fall-off begins at about 10 pm diameter. Unfortunately going to larger probes and hence larger currents does

not provide detection limits better than about

0.1% atomic. Such an effect is observed since

the sensitivity is detector limited in the

AES technique. However, with the ion probe

technique, since material is consumed and the

technique is not detector limited, the volume

consumed determines the number of ions that are available for analysis. Thus there is an interrelationship between the realizable detection limit and the probe area. Given the

assumptions shown in the box of Fig. 3, the

larger the analyzed area, the better are the

detection limits for SIMS. The attainable

limits will vary depending on the ion yield

of the different elements. For the sensitive

elements such as aluminum and iron, the ion

probe technique (for all probe sizes) is

much more sensitive than Auger. However, as elements with lower yields are analyzed, the

ion probe may be more or less sensitive than

Auger depending on the available analytical

area. Before one can choose Auger or SIMS

for a microanalysis, the relative sensitivity

of the sought-for element and the analytical

area must be considered.

As has been discussed, all of these techniques

provide an in-depth analytical capability.

Backscattering spectrometry samples all depths

simultaneously and, therefore, gives an in- stant readout of all depths as well as chemical identification. The other techniques - AES, ESCA, ISS and SIMS require ion sputtering with the analytical excitation in order

to realize a depth profile. When using the

sputtering techniques one must be concerned

about a variety of artifacts which can distort the analytical information. One must

be concerned about how the sputtering process

distorts the chemical information provided by

ESCA. Is the ion beam going to damage the

sample so that it is no longer in the same chemical state it was before? Under

221
reasonable conditions, the sputtering process

can provide remarkable results. We have been

able to detect with both AES and SIMS a 5 angstrom layer of antimony under 8,000 angstroms of amorphous silicon. Even after diffusion we were able to detect the antimony.

Sputtering in either amorphous or in single

crystal situations is not a particular problem. It is only when polycrystalline , com- pound or multi-phase systems are sputtered

that one can run into problems. Depth reso- lution in the backscattering technique is

controlled by detector resolution and energy

straggling in the sample. Whereas in all

the techniques employing sputtering, the

controlling parameters are the flatness of

the crater, differential sputtering, "knockon" of atoms into deeper layers, and the

need to calibrate the sputtering rate. This

latter point is quite important because some- how time must be converted into distance.

Table 6 summarizes the concepts and concerns

to be considered in depth profiling.

If the analyst is required to depth profile

a large volume of thin films, he must be

concerned about analytical time and system

turnabout. Table 7 compares the five

techniques in the time required to depth profile a 5000A film. Backscattering is very

fast. All accessible depths are sampled

simultaneously, and the vacuum requirements

are minimal. It is not a monolayer surface

analytical technique but really a technique

for chemical and compositional analysis of

multilayer thin films. The time for a 5,000

angstrom film is 10 or 15 minutes. Since

the analytical time is independent of thickness, the "same time would be required for a 1,000 or 10,000 angstrom film. Backscattering spectrometry is the fastest technique

for depth profiling thick films (or thin

films depending on your definition) . The

SIMS technique is also generally quite fast.

The use of sample carousel systems or vacuum • interlocks provides fast sample turnaround

times. Approximately 10 to 15 minutes would

be required for this 5,000 angstrom film since

sputtering rates of 300 to 500 angstroms per

minute are attainable. However, to realize

this type of efficiency, a computer or multichannel analyzer data system must be employed.

AES is generally slow in its turnaround time

since vacuum compatibility with the analytical process is required and the spectrometer

must be calibrated for every sample. The

sputtering rates employed in AES can approach

those of SIMS. However the analyst must be

somewhat careful if he wishes to monitor

subtle changes in concentration because it

takes a finite amount of time to measure

Auger peak height or SIMS ion intensity.

Thus depth resolution can be reduced if

sputtering rates are too large. The ESCA

and ISS techniques are very slow in their

turnaround time because of the low signal

to noise encountered, which implies a very

slow sputtering rate. They are not amenable

to the depth profiling of very thick films.

They are, of all of these techniques, the true surface analytical techniques. They

are not intended for, or amenable to, thick

film depth profiling.

Several general conclusions (Table 8) can be

made about these techniques. Auger is a good

technique with its sensitivity to low Z elements, minimum matrix effects, and microanalytical capability. Backscattering, most

importantly, is fast as well as being quantitative in concentration and depth. This is

extremely important in the characterization

of thin films. Unfortunately, it has very

poor sensitivity for low Z elements and has

poor lateral resolution. The most outstanding

feature of ESCA is its chemical information.

The outer monolayer analysis capability of ISS is its distinguishing feature. SIMS has

ppm detection limits for almost all the elements, provides isotopic resolution, is fast

and is a microanalytical technique. As far

as disadvantages are concerned, ESCA has poor

detection limits, poor lateral resolution and

very slow profiling time. ISS has low sensitivity, poor lateral resolution and slow profiling. The limitations of SIMS are the

troubles with quantitation and the matrix

effects encountered.

The most important conclusion from the above

discussion is that complete characterization

of a system may well require more than one of

the above techniques . Joint research on re- actions and transport in thin films with M-A. Nicolet and J. W. Mayer of Caltech has evolved

a multi-technique approach to the characterization of thin film systems as illustrated in

Fig. 4.

Every sample is subjected to a backscattering

analysis. A large number of samples per unit

time can be characterized in this manner. If we wish to do additional in-depth profiling,

we then go to AES because of its low Z element

capability. This is to supplement the BS analysis to provide the major and minor element overview and depth scale calibration.

If we must go to low concentrations for impurities and dopants, we then go to SIMS. If they were available to either laboratory, we would use ISS for outer monolayer analysis

and ESCA for chemical information. We do not have these two techniques but they are in the

figure for prospective. If we require lateral

analysis after the backscattering analysis,

scanning Auger microscopy or the ion

222
microprobe is used. We find more and more

it has become necessary to begin to look at structural information because of transport,

precipitation, etc., when we are studying

reactions. So structural information is provided by using optical microscopy for

morphology or an SEM (with X-ray detection

for elemental characterization) . If we wish

to do microstructure phase identification or crystal quality, we go to transmission electron microscopy and use X-ray diffraction

for macrophase identification.

For other perspectives of the relative

features of these techniques and the instrumentation required the reader is referred to references [3-7].

LITERATURE CITED

1. Chu, W. K. , Nicolet, M-A. , Mayer, J. W.

,

and Evans, Jr., C. A., Analytical

Chemistry 46, 2136 (1974).

2. Cookson, J. A., and Pilling, F. D. , Thin

Solid Films 19, 381 (1973).

3. Coburn, J. W., and Kay, E. , CRC Critical

Reviews in Solid State Sciences 4_, 561

(1974)

.

4. Honig, R. E. , Surface, and Thin-Film

Analysis of Semiconductor Materials,

Thin Solid Films, 1975 (to be published).

5. Evans, Jr., C. A., J. Vac. Sci. Technol.

12, 144 (1975).

6. Evans, Jr., C. A., Surface and Thin Film

Compositional Analysis - A Description

and Comparison of Techniques, Anal.

Chem. Feature article (August, 1975).

7. Evans, Jr., C. A., Instrumentation for

Surface and Thin Film Analysis, Anal.

Chem., Instrumentation article, (August,

1975).

223
o

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CD CJ

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4->

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o -C

>> 03 r— >,+-> O 03 i

>,-!-> — c

c +-> T- O +->!- LO o o o C 4— •t- 3 -— C 4- i- 3

> H3 O > 03 O > ee CO 4->

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CD

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h- QJ O CO (_) ^ O O O CO O WZ CT) QJ LU >• QJ CO QJ i- CO c QJ CO O CO l—l 3 a.<c QJ CLCO — O LlJ o Q-i—

QJ 03 CO <: co -—

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224
TABLE 2

MATRIX OR CHEMICAL EFFECTS

Observed Extent Comment

1. AES Yes Few

El ements

Generally useable

2. BS No

3. ESCA Yes All Elements

Basis & justification

for technique

4. ISS Yes Depends on matrixInfluences extent of

ion neutralization

5. SIMS Yes Often Can be used but is generally a nuisance

TABLE 3

SAMPLE CONSUMPTION and/or ALTERATION

1. Non-destructive in all modes

a. BS (can get radiation damage of organics and

in channeling studies)

2. Non-destructive during analysis

Destructive for depth profiles

may get radiation (e~ or hv) damage

3. Sample consumed for all modes

a. ISS - sputtering rates may be slow with k He

+

b. SIMS - volume of material available controls

elemental sensitivities

225
TABLE 4

QUANTITATION (CONCENTRATION J« DEPTH SCALE)

Absolute Standards Mechanistic Procedures

1AES No Needed Just beginning

2. BS Yes

(± 5%)

Help in some cases (± 2%)Permits absolute

quantitation

3. ESCA No Needed

4. ISS No Needed Can use known scattering

cross-section for relative

concentration information

5. SIMS No Needed Most successful gives

± 200% in almost all cases

± 5-20% in many cases

TABLE 5

LATERAL ANALYSIS

1. Microanalysis [<_ 10 pm resolution)

a. AES via microfocused primary electron beam

b. SIMS via microfocused primary ion beam (ion microprobe)

or stigmatic secondary ion optics (ion microscope)

2. Millianalysis (^ 0.1 mm resolution)

a. ISS with millifocused primary ion beam

3. Macroscopic area (> 1 mm resolution)

a. BS generally uses 1 mm probing beams

b. ESCA illuminates large areas (> 1 mm2

) with x-rays

226
TABLE 6

IN-DEPTH PROFILING ANALYSES

1. Backscattering spectrometry samples all depths simultaneously

2. AES, ESCA, ISS & SIMS require ion sputtering for sequential

layer removal

CONCERNS WITH SPUTTERING

a. Analytical signal must come from flat-bottomed part of crater

b. Must be aware of sputtering artifacts

1. Preferential removal of one element over another

2. Crystal 1 inity effects

3. Sample may be multi-phase

227
S- £= o CL) 4- O O

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CU -C

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C

r- E 4-> E )

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LjJ c O 1 o< 0) o CL LO

U_

O cu 1 O o o o o c to o Q_ LO to CO

a> rc s_ r— o

CL. 4- c o c

UJQ QJ E E E +-> CC lo LO o r— E

u_ o O Q 4- UJ

cc

ZDoUJ

ME

i— CUE

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228

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CU E cn s_ o<: •i- sz .—

1— <_> • o •r- -r- C o 4- -C CM O 5 E 1 S_ \ o O. CUoeC LO co O -i- o _ ^ +-> O i

C 1 "O 1 i

cu o

-(-> +J E 'r— 1 3 i— * o to ro to ro C CU I +-> CD +-> LO O t- ro CM Z 00 1. _

00 on

+-> +-> 1 c: cu u CD E cu E

cu a. CU

S- 00 t. 3 cu 3

O"

cu (O cu

S- S- s_ -Q 3 E O =3 O 3 i— 3 ro i— 3

to

O

O O0 CJ

ro ro >> > o ^ >> >

+-> S- r— O cu i—

O

ro +-> "O -!-> ro +-> t- CU CU S_ cu cu CU E CU cu C 3 c o C 3

ai -o i- CU "O

CJ3 oS +-> o '

oa=t

oo

O CO

LU 00 oo

<c UJ I—

oo
TABLE 8

CONCLUSIONS

Havaniages u\ saQVanuayes

1 AFS G > oensiLivity to low L

aa . U 1 1 1 1 t_ U l L L.U l)Ua II C 1 La Lc

b. Minimal matrix effects b. 0.1% detection limits

r-C . M "i c y*c\7\ nal wci c ll l L i Ud j la l jo 1 o u

.

(Zr\r\r\ uUUU oil av^nuhn torhni niifl d 1 1 cuUUMU LcLiiniLjUc

2. BS a. Fast a. Low sensitivity especially

b. Ouant i ta ti vp in cone and for low Z depth b. Poor lateral resolution

3. ESCA a. Chemical information a

.

0.1% detection limits

b. Poor lateral resolution

c. Slow profiling

4. ISS a. Outer monolayer analysis a. Low sensitivity

b. Poor lateral resolution

c. Slow profiling

5. SIMS a. ppm detection limits for a

.

Quantitation and matrix

many elements effects

b. Isotopic resolution

c. Fast

d. Microanalysis

229
,H Ne Co Zn Zr Sn Nd Yb Hg Th

1

1

1 1 1 1 1 1 1 1 r~

Ion Scattering Spectrometry

2 10-^-

10"

10"

3 ioAuger*Bectron Spectroscopy;

. -^s. -

'

I

1i

i

I

i

s

I

II

I

I

I

I

I

I

I

MeV 4 He+ / _

Backscattering Spectrometry

/ / Secondary Ion

/ / Mass Spectrometry

/ i

20

1%

Figure 1. A comparison of bulk detection

limits vs. atomic number for five surface

sensitive analytical techniques.

lppm

40 60

Atomic Number

80 100lppb

THICKNESS (A)

Figure 2. Energy loss by 4 He in Ta 2 0 5 vs. known thickness.

Curves represent various calculated relationships.

230
Figure 3. Calculated detection limits vs. analytical area for AES and SIMS.

ELECTRONIC MATERIALS CHARACTERIZATION

AT CALTECH a UNIV. OF ILLINOIS

Mev 4 He BS

OPTICAL MICROSCOPY

1. MORPHOLOGY

2. ELEMENTAL

I DE NT SEM »/ X-RAY

1. CRYSTAL QUALITY

2. MICROSTRUCTURE

3. PHASE I DENT

MACRO- PHASE X-RAY

'DENT. DIFFRACTION

LATERAL

ANALYSIS

RAPID QUANT CONCENTRATION

S IN-DEPTH ANALYSIS

ION MICROPROBE

I ISS

I.LO-Z ELEMENT

2 MAJOR a MINOR ELEMENT OVERVIEW

I LATERAL RESOLUTION

2. LOW CONC. DOPANTS

8 IMPURITIES

OUTER MONOLAYER

ESCA

j

CHEMICAL OR

BONDING

Figure 4. Flow chart for electronic materials characterization at the California Institute of Technology

and the University of Illinois.

231
DISCUSSION OF THE PAPER

Participant : In one of the slides you

show sensitivity vs beam diameter for Auger

and SIMS, where does the hydrogen sit on that

diagram?

Evans : If you remember in the beginning we excluded hydrogen as being unobservable by

the Auger. As far as SIMS is concerned, we

are not really sure. I am only familiar with

the work in our own laboratory on hydrogen

and we have not really been looking at low

levels. We have been looking at the formation of niobium hydride in hydrogen charged

niobium. I would guess the ionization yield

for H probably sits between 10"^ and 10~6

ions/atom and as such H appears in the middle

of the diagram.

Newbury : Is it worth mentioning that with

the selection of the primary ion beam you

might, in fact, get enhanced sensitivity from

say, gold, carbon, or sulphur by possibly

using a cesium beam? It is possible by

selection of your beam to improve the

situation

Evans : Hughes and other laboratories are developing cesium ion sources for ion microprobes. If you do use cesium ion bombardment

and negative ion spectroscopy instead of

oxygen ion bombardment and positive ion

spectroscopy, those elements such as gold,

carbon, etc., which used to be very insensitive are now very sensitive. In fact, in

positive ion spectroscopy sodium is the most

sensitive element, in negative ion spectroscopy using cesium ion bombardment gold is

probably the most sensitive element. So as instrument developments come along we may

find improvements in the Auger curve but

there will definitely be a dramatic improvement in the SIMS curves as soon as a good

cesium ion source is developed.

Lucke : You mentioned one of the disadvantages

of SIMS was quantitation. I wonder what you

could say about the possibilities of eventually obtaining absolute concentration numbers.

Evans : There are laboratories working on the

fundamentals of the secondary ionization

process. The most successful of the resultant

techniques are those used by Anderson, and

Anderson and Hinthorn. Basically, they try

to calculate what the conditions are in the

ionization region. They assume that it is a plasma in local thermal equilibrium. Then

the actual concentrations are calculated from

the assumed ionization conditions and the

measured secondary ion intensities.

Lieberman : The alternative approach toward

quantitation is to develop working curves based on measurements of reference materials

having elemental concentrations similar to those of the unknown sample. The Electronic

Technology Division of NBS is endeavoring to develop and document techniques for the application of SIMS to silicon material and

device analysis. This work includes the

development of methods of sample preparation

and data analysis for reliable quantitative

SIMS analysis of impurities on silicon and

silicon dioxide surfaces.

232
WORKSHOP PARTICIPANTS

John Albers

National Bureau of Standards

Washington, DC 20234

Gordon Anderson

Naval Research Laboratory

Washington, DC 20375

James Anthony

National Security Agency

Fort Meade, MD 20755

Donna Bakale

Varian Vacuum

Springfield, NJ 07081

W. D. Baker

Naval Research Laboratory

Washington, DC 20375

William R. Bandy

National Security Agency

Fort Meade, MD 20755

Robert Barr

Technics Inc. Alexandria, VA 22310

Lt. Richard Bernard

Rome Air Development Center

Griffiss Air Force Base

Rome, NY 13440

J. B. Bindell

Bell Telephone Laboratories, Inc. Allentown, PA 18103

W. R. Bottoms

Electrical Engineering Department

Princeton University

Princeton, NJ 08540

M. G. Buehler

National Bureau of Standards

Washington, DC 20234

James W. Butler

Naval Research Laboratory

Washington, DC 20020

Sidney R. Butler

Lehigh University

Department of Metallurgy and Materials

Science

Bethlehem, PA 18015

Robert C. Cargill

Physical Electronics Industries, Inc. Eden Prairie, MN 55343

Forrest L. Carter

Naval Research Laboratory

Washington, DC 20034

Thomas R. Cass

Hewlett-Packard Laboratories

Palo Alto, CA 94304

Jar-Mo Chen

Martin Marietta Laboratory

Baltimore, MD 21227

A. Christou

Naval Research Laboratory

Washington, DC 20375

Thomas E. Cody

Applied Research Laboratories

Allentown, PA 18101

Donald K. Conley

Western Electric Company

Allentown, PA 18103

Dwight Cook

Bell Telephone Laboratories

Murray Hill, NJ 07974

Lawrence P. Cook

National Bureau of Standards

Washington, DC 20234

Becky P. Corbin

Commonwealth Scientific Corporation

Alexandria, VA 22314

Howard S . Danaceau

The Kearns Group

Rockville, MD 20850

David Darbo

NASA Goddard Space Flight Center

Greenbelt, MD 20771

John E . Davey

Naval Research Laboratory

Washington, DC 20375

L. E. Davis

Physical Electronics Industries, Inc. Eden Prairie, MN 55343

Howard M. Day

Naval Research Laboratory

Washington, DC 20390

Bruce E. Deal

Fairchild Semiconductor

Palo Alto, CA 94304

233
T. H. DiStefano

IBM Research

Yorktown Heights , NY 10598

Robert D. Dobrott

Texas Instruments, Inc. Dallas, TX 75222

L. G. Dowell

Union Carbide Corporation

Tarrytown, NY 10591

Kenneth L. Dunning

Naval Research Laboratory

Washington, DC 20375

John W. Dzimianski

Westinghouse Electric Corp., S.D.D

Baltimore, MD 21203

James Ehrstein

National Bureau of Standards

Washington, DC 20234

Nils E. Erickson

National Bureau of Standards

Washington, DC 20234

Charles A. Evans, Jr. Materials Research Laboratory

Urbana, IL 61801

R. J. Farrar

Harry Diamond Laboratories

Washington, DC 20438

Frank J. Feigl

Department of Physics

Lehigh University

Bethlehem, PA 18015

Charles Feldman

The Johns Hopkins University

Silver Spring, MD 20910

David File

Cameca, Inc. Elmsford, NY 10523

Judson C. French

National Bureau of Standards

Washington, DC 20234

Bill Gadzuk

National Bureau of Standards

Washington, DC 20234

Kenneth Galloway

National Bureau of Standards

Washington, DC 20234

Anthony J. Garratt-Reed

Massachusetts Institute of Technology

Cambridge, MA 02139

Richard Gerber

Aerospace Corporation

Los Angeles, CA 90009

Louis F. Giles, III. National Security Agency

Fort Meade, MD 20755

Michael N. Giuliano

Westinghouse/Advanced Technology Laboratory

Baltimore, MD 21203

Ralph Gorden , Jr. National Bureau of Standards

Washington, DC 20234

Leonard P. Grady

Deltaray Corporation

Burlington, MA 01803

Edward T. Graney

Physical Electronics Industries, Inc. Springfield, NJ 07090

Henry F. Gray

Naval Research Laboratory

Washington, DC 20375

Frank Grunthaner

Jet Propulsion Laboratory

Pasadena, CA 91103

George A. Haar

Naval Research Laboratory

Washington, DC 20375

Hiroshi Harada

University of West Lafayette

West Lafayette, Indiana 47907

W. L. Harrington

RCA Laboratories

David Sarnoff Research Center

Princeton, NJ 08540

James K. Hirvonen

Naval Research Laboratory

Washington, DC 2037 5

H. Ward Huang

IBM Corporation

Essex Junction, VT 05452

H. Hughes

Naval Research Laboratory

Washington, DC 20375

234
James M. Hurd

Tektronix Inc. Beaverton, Oregon 97077

Harry M. Hyman

RCA Corporation

Somerville, NJ 08876

N. L. Jarvis

Naval Research Laboratory

Washington, DC 20375

Marilyn A. Jasper, FIR

Night Vision Laboratory

Ft. Belvoir, VA 22060

Murzban Jhabvala

NASA Goddard Space Flight Center

Greenbelt, MD 20771

Francis Kub

Westinghouse ATL Baltimore, MD 21203

Chris Kuyatt

National Bureau of Standards

Washington, DC 20234

Alan J. Kushnir

Semi-Films

West Hurley, NY 12491

Max G. Lagally

University of Wisconsin

Madison, WI 53706

G. M. Lamb

NASA Goddard Space Flight Center

Greenbelt, MD 20771

Jan S. Johannessen

Stanford Electronics Laboratory

Stanford University

Stanford, CA 94305

Phil F. Kane

Texas Instruments Inc. Dallas, TX 75222

W. J. Keery

National Bureau of Standards

Washington, DC 20234

G. J. King

Army Night Vision Laboratory

Alexandria, VA 22307

T. D. Kirkendall

Comsat Laboratories

Clarksburg, MD 20734

W. J. Kitchen, Jr. Department of Defense

Ft. Meade, MD 20755

Bruce E. Knox

Penn State University

University Park, PA 16802

Robert J. Kopp

Fairchild Camera and Instrument Corporati

Mountain View, CA 94040

Ram Kossowsky

Westinghouse Research

Pittsburgh, PA 15235

R. Koyama

National Bureau of Standards

Washington, DC 20234

Donald E. Lander

AEI Scientific Apparatus, Inc. Falls Church, VA 22043

Graydon B. Larrabee

Texas Instruments Inc. Dallas, TX 75222

K. 0. Leedy

National Bureau of Standards

Washington, DC 20234

T. F. Leedy

National Bureau of Standards

Washington, DC 20234

Douglas H. Leong

Hughes Aircraft Company

Newport Beach, CA 92663

Leonard L. Levenson

University of Missouri-Rolla

Rolla, MO 65401

Miguel E. Levy

Hughes Aircraft Company

Culver City, CA 90230

Robert K. Lewis

Cameca Instruments, Inc. Elmsford, NY 10523

A. G. Lieberman

National Bureau of Standards

Washington, DC 20234

Wen N. Lin

Rockwell International Corporation

Anaheim, CA

235
Y. Liu

National Bureau of Standards

Washington, DC 20234

W. H. Lucke

Naval Research Laboratory

Washington, DC 20735

Fred Luehrs

Tousimis Research Corporation

Rockville, MD 20852

Ingemar Lundstrom

IBM Corporation

Yorktown Heights, NY 10589

Gary McGuire

Texas Instruments, Inc. Dallas, TX 75222

R. M. McLouski

Westinghouse

Baltimore, MD 21203

Joseph F. Martino

Night Vision Laboratory

Ft. Belvoir, VA 22060

Santos Mayo

National Bureau of Standards

Washington, DC 20234

David M. Metz

Penn State University

University Park, PA 16802

Reid A. Mickelsen

Boeing

Bellevue, WA 98006

T. A. Midford

Hughes Research Center

Torrance, CA 90509

Gaines W. Monk

Commonwealth Scientific Corporation

Alexandria, VA 22314

J. M. Morabito

Bell Telephone Laboratories

Allentown, PA 18103

Bernard Mrstik

Naval Research Laboratory

Washington, DC 20375

J. Munarin

Naval Ammunition Depot

Crane, IN 47522

J. S. Murday

Naval Research Laboratory

Washington, DC 20375

Brian A. Murphy

Semi-Films Technology

West Hurley, NY 12491

Madan Nanda

IBM/SDD

Manassas, VA 22110

Dale E. Newbury

National Bureau of Standards

Washington, DC 20234

Paul Norton

Bell Telephone Laboratories, Inc. Allentown, PA 18103

Mylous S. O'Dell

Night Vision Laboratory

Fort Belvoir, VA 22060

F. F. Oettinger

National Bureau of Standards

Washington, DC 20234

Victor J. Ohm

Rome Air Development Center

Griffiss Air Force Base, NY 13441

Masatoshi Ono

National Bureau of Standards

Washington, DC 20234

Philip F. Ordung

University of California

Santa Barbara, CA 93106

J. Oroshnik

U. S. Navy

Rockville, MD 20853

John M. Pankratz

Texas Instruments, Inc. Dallas, TX 75222

Donald L. Parker

Texas A & M University

College Station, TX 77843

Martin C. Peckerar

Naval Research Laboratory

Washington, DC 20375

Brad Phillips

Naval Ammunition Depot

Crane, IN 47522

236
Howard Phillips

Rockwell International Corporation

Anaheim, CA 92806

Dan Pierce

National Bureau of Standards

Washington, DC 20234

Larry E. Plew

Naval Ammunition Depot

Crane, IN 47522

Juergen L. W. Pohlmann

Army Night Vision Laboratory

Fort Belvoir, VA 22060

Cedric Powell

National Bureau of Standards

Washington, DC 20234

R. A. Reynolds

Advanced Research Projects Agency

Arlington, VA 22209

A. G. Revesz

Comsat Laboratories

Clarksburg, MD 20734

Giorgio Riga

Fairchild Camera

Mountain View, CA 94042

Frank G. Satkiewicz

Johns Hopkins University

Applied Physics Laboratory

Silver Spring, MD 20910

D. E. Sawyer

National Bureau of Standards

Washington, DC 20234

Harry Schafft

National Bureau of Standards

Washington, DC 20234

George H. Sigel

Naval Research Laboratory

Washington, DC 20375

Ronald L. Smith

National Security Agency

Fort Meade, MD 20755

Sherman D . Smith

Burroughs Corporation

Downingtown, PA 19335

Richard Stein

National Bureau of Standards

Washington, DC 20234

Yale Strausser

Varian Vacuum Division

Palo Alto, CA 94303

Barbara E. Sumner

Army Night Vision Laboratory

Fort Belvoir, VA 22060

Yen T. Tan

Kodak Research Laboratories

Rochester, NY 14650

Simon Thomas

Motorola SPD

Phoenix, AZ 85008

A. J. Tousimis

Tousimis Research Corporation, Inc. Rockville, MD 20852

0. D. Trapp

Technology Associates

Portola Valley, CA 94025

Richard E. Tressler

Pennsylvania State University

University Park, PA 16802

J. C. C. Tsai

Bell Telephone Laboratories, Inc. Reading, PA 19603

Charles Vath

Intel

Santa Clara, CA 95051

Theodore Vorburger

National Bureau of Standards

Washington, DC 20234

Bernie Waclawski

National Bureau of Standards

Washington, DC 20234

M. S. Wang

Monsanto Company

St. Louis, MO 63166

John C. Webber

IBM Corporation

Hopewell Junction, NY 12533

Clark W. White

Bell Telephone Laboratories, Inc. Murray Hill, NJ 07974

W. R. Whitford

ORTEC, Inc. Pearl River, NY 10965

237
Robert G. Wilson

Hughes Research Laboratories

Malibu, CA 90265

W. A. Wolstenholme

AEI Scientific Apparatus

Elmsford, NY 10523

K. H. Yang

IBM Corporation

Hopewell Junction, NY 12590

Carl Zanoni

Zygo Corporation

Middlefield, CT 06455

238
AUTHOR INDEX

AUTHOR

Austin, A. E. . .

Bottoms, W. R.

Christou, A. . .

Davis, L. E. . .

Day, H. M. . . .

Deal, B. E. . . .

DiStefano, T. H.

Dobrott, R. D.

Dunning, K. L.

Evans , Jr . , C. A.

Franz, J. M. . .

French, J. C. . .

Grunthaner, F. J.

Harrington, W. L.

Hughes, H. L. . .

Johannessen, J. S

Lewis, R. K. . .

Lieberman, A. G.

McGuire, G. E.

Morabito, J. M. .

Phillips, B. F. .

Phillips, D. H. .

Riach, G. E.

Spicer, W. E. . .

Strausser, Y. E.

Thomas , S . ...

Weisenberger , W.

White, C. W.

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. BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO. NBS SP 400-2?

2. Gov't Accession

No.

3. Recipient's Accession No.

4. TITLE AND SUBTITLE

Semiconductor Measurement Technology: ARPA/NBS Workshop IV,

Surface Analysis for Silicon Devices

5. Publication Date March 1976

6. Performing Organization Code

7. AUTHOR(S)

A. George Lieberman

8. Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/Work Unit No.

11. Cont rac c /(jr<in t No. ARPA Order No. 2397

Program Code 5D10

12. Sponsoring Organization Name and Complete Address (Street, City, State, ZIP)

Defense Advanced Research Projects Agency - 1400 Wilson Blvd.

Arlington, Virginia 22209

National Bureau of Standards, Washington, D. C. 20234

13. Type of Report & Period Covered Final

April 23-24, 1975

14. Sponsoring Agency Code

15. SUPPLEMENTARY NOTES

Library of Congress Catalog Card Number: 76-608043

16. ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a significant

bibliography or literature survey, mention it here.) This report contains the proceedings of the ARPA/NBS Workshop IV, Surface

Analysis for Silicon Devices, held at the National Bureau of Standards on April 23-24, 1975.

The Workshop, as part of an NBS program to develop measurement technology for

the field of semiconductor devices, was held to discuss the present capabilities and

future prospects of modern analytical beam techniques as applied to silicon, and

associated insulator films and device structures. Of particular interest were the

determination of impurity profiles, surface contamination, and interface characteristics. Techniques utilizing impinging electron, ion, neutral or photon beams were considered. The Workshop was directed at the analysts, the semiconductor

manufacturers who use the analysts' results, and the instrument people who design

and manufacture the analytical equipment. Transcripts of the discussions following

each paper are also included within these proceedings.

17. KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper name; separated by semicolons) Auger spectroscopy; depth profiles; electron beam induced

imaging; ESCA; insulator films; interface characteristics; internal photoemission ; ion

scattering spectroscopy; nuclear resonance profiling; photodepopulation; photovoltaic

imaging: Rutherford backscattering : SCANIIR; secondary ion mass spectroscopy; semiconductor devices; silicon; surface analysis; X-ray photoelectron spectroscopy.

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<HJ.S. GOVERNMENT PRINTING OFFICE: 1976 210-801/190 1-3

Announcement of New Publications on Semiconductor Measurement Technology

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