Sputtering of compound semiconductor surfaces. II. Compositional changes and radiation-induced topography and damage

This article was downloaded by: [Universitaets und Landesbibliothek]On: 10 December 2013, At: 00:27Publisher: Taylor & FrancisInforma Ltd Registered in England and Wales Registered Number: 1072954 Registered office: Mortimer House,37-41 Mortimer Street, London W1T 3JH, UK

Critical Reviews in Solid State and Materials SciencesPublication details, including instructions for authors and subscription information:http://www.tandfonline.com/loi/bsms20

Sputtering of compound semiconductor surfaces.II. Compositional changes and radiation-inducedtopography and damageJohan B. Malherbe aa Department of Physics , University of Pretoria , Pretoria, 0002, South AfricaPublished online: 24 Oct 2006.

To cite this article: Johan B. Malherbe (1994) Sputtering of compound semiconductor surfaces. II. Compositional changesand radiation-induced topography and damage, Critical Reviews in Solid State and Materials Sciences, 19:3, 129-195, DOI:10.1080/10408439408244589

To link to this article: http://dx.doi.org/10.1080/10408439408244589

PLEASE SCROLL DOWN FOR ARTICLE

Taylor & Francis makes every effort to ensure the accuracy of all the information (the “Content”) containedin the publications on our platform. However, Taylor & Francis, our agents, and our licensors make norepresentations or warranties whatsoever as to the accuracy, completeness, or suitability for any purpose of theContent. Any opinions and views expressed in this publication are the opinions and views of the authors, andare not the views of or endorsed by Taylor & Francis. The accuracy of the Content should not be relied upon andshould be independently verified with primary sources of information. Taylor and Francis shall not be liable forany losses, actions, claims, proceedings, demands, costs, expenses, damages, and other liabilities whatsoeveror howsoever caused arising directly or indirectly in connection with, in relation to or arising out of the use ofthe Content.

This article may be used for research, teaching, and private study purposes. Any substantial or systematicreproduction, redistribution, reselling, loan, sub-licensing, systematic supply, or distribution in anyform to anyone is expressly forbidden. Terms & Conditions of access and use can be found at http://www.tandfonline.com/page/terms-and-conditions

http://www.tandfonline.com/loi/bsms20

http://www.tandfonline.com/action/showCitFormats?doi=10.1080/10408439408244589

http://dx.doi.org/10.1080/10408439408244589

http://www.tandfonline.com/page/terms-and-conditions

http://www.tandfonline.com/page/terms-and-conditions

Critical Reviews in Solid State and Muterials Sciences . 19(3): 129-195 (1994)

Sputtering of Compound Semiconductor Surfaces . II . Compositional Changes and Radiation-Induced Topography and Damage

Johan B . Malherbe Department of Physics. University of Pretoria. Pretoria. 0002. South Africa

1 . Introduction ............................................................................................................................ 130 2 . Radiation-Induced Topography ............................................................................................. 131

2.1. Introduction .................................................................................................................. 131 2.2. Radiation-Induced Morphology Mechanisms .............................................................. 131

2.2.1. 2.2.2. 2.2.3. 2.2.4. 2.2.5. 2.2.6. 2.2.7. 2.2.8. 2.2.9.

Impurity Seeding Models ............................................................................... 132

Crystallographic Orientation-Dependent Models .......................................... 132 Surface Migration Model ............................................................................... 133 Ripple Models ................................................................................................ 133 Redeposition Model ....................................................................................... 133 Whisker Growth Model ................................................................................. 133 Microregion Model ........................................................................................ 133 InP Model ....................................................................................................... 134

Differential Sputter Erosional Models ........................................................... 132

2.3. Radiation-Induced Topography on InP ........................................................................ 134 2.4. Radiation-Induced Topography Development on Other Compound

Semiconductors ............................................................................................................ 138 Radiation-Induced Damage .................................................................................................... 140 3.1. Introductory Remarks ................................................................................................... 140 3.2.

3.2.1. GaAs ............................................................................................................... 141 3.2.2. InP and Other III-V Semiconductors ............................................................. 144 3.2.3. S i c .................................................................................................................. 146 3.2.4. II-VI Compounds ........................................................................................... 146 3.2.5. Hg,-, Cd, Te .................................................................................................... 148

3.2.5.1. Structural Damage .......................................................................... 148 3.2.5.2. Annealing of Hg,-, Cd, Te .............................................................. 149 3.2.5.3. Electrical Damage and Junction Formation .................................. 149

3.3. Concluding Remarks .................................................................................................... 151 Compositional Changes ......................................................................................................... 151 4.1. Introduction .................................................................................................................. 151 4.2. Preferential Sputtering .................................................................................................. 151 4.3. Kinetic Processes .......................................................................................................... 153

4.3.1. Gibbsian Segregation ..................................................................................... 154 4.3.2. Radiation-Enhanced Segregation ................................................................... 154

. 4.3.3. Segregation ..................................................................................................... 154 4.3.4. Radiation-Enhanced Diffusion ....................................................................... 154

3 .

Radiation-Induced Damage in Some Compound Semiconductors ............................. 141

4 .

~~

1040-8436/94/$.50 0 1994 by CRC Press. Inc .

129

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

4.3.5. Displacement Mixing ..................................................................................... 154 Implantant Surface Concentration ................................................................. 154

Sputtering of InP .......................................................................................................... 155 4.4.1. Experimental Measurements .......................................................................... 155 4.4.2. Factors Affecting the Accuracy of the Measurements .................................. 157 4.4.3. 4.4.4.

4.5. Sputtering of GaAs ....................................................................................................... 167 4.6. Sputter-Induced Compositional Changes in Other Compound

Semiconductors ............................................................................................................ 176 4.7. Summary ....................................................................................................................... 181

5. Conclusions ............................................................................................................................ 181 References .............................................................................................................................. 182 Corrections to Part I ............................................................................................................... 196

4.3.6. 4.4

Discussion of the Experimental Data ............................................................ 161 Predictions of Theoretical Models ................................................................. 165

ABSTRACT: Ion bombardment often leads to compositional changes in the burface layers of multicomponent targets. Such changes due to noble gas ion sputtering are discussed for InP and GaAs. The analyses show that the compositional change in InP (i.e., indium enrichment) is mainly due to preferential sputtering. In the case of GaAs. the changes are due to radiation-induced diffusion and segregation effects. Brief mention is made of compositional changes in a few other systems. The discussion on sputter-induced topography development deals mainly with InP because ion bombardment leads to dramatic topographical effects in this material. Ripple development on GaAs is also briefly discussed. Radiation damage has been well researched, and its mechanism and effects usually differ substantially when going from one semiconductor group to another. Bombardment-induced damage is briefly discussed for InP. GaAs. SIC, some 11-VI semiconductors. and HgCdTe.

KEY WORDS: compound semiconductors. bombardment-induced morphology, sputter-induced topography. sputter yield. ion-solid interactions, radiation-induced damage, structural damage. electrical damage, annealing, surface composition change, segregation. diffusion, preferential sputtering, Ill-VI semiconductors, 11-V semiconductors, compound semiconductors. GaAs. InP. Sic. CdTe, HgCdTe.

1. INTRODUCTION

In the first part of this review,' the importance of ion bombardment-induced effects in compound semiconductors was briefly outlined. Some of many ion-solid interactions due to bombardment were also mentioned. In many of these interactions, the resulting effects and processes cannot be discussed independently of each other. This was also evident in the discussion in Ref- erence 1 of sputter yields due to noble gas bombardment of compound semiconductors. The sputter yields depend inter alia on the three topics discussed in this part of the review, that is, bombardment-induced surface topography development, radiation damage, and compositional changes. For example, topography development increases the exposed surface area of the target, and this might change the sputter yields. The question of changes in the sputter yields from the different surfaces of a monocrystalline compound semiconductor target is

intimately connected to the radiation-induced damage in the target (see Section 3.4.1 in Refer- ence I). Similarly, bombardment-induced compositional changes in the surface layers of a target may cause a change in the sputter yields.

This review is published as two papers, with each divided into several sections. The main topic discussed in part 1 of this review was sputter yields of compound semiconductors when bombarded by noble gas ions. The aim was to compare the experimental sputter yield values and the theoretical predictions of the major sputtering theories. This is not a trivial exercise because these theories were developed for single elemental targets. There is a lack of experimental data on component semiconductors with the exception of GaAs. Even in the case of GaAs, there are few total sputtering yield measurements for ion energies above 10 keV. Because most experimental noble gas sputtering yield values exists for GaAs, only this system was compared with theoretical predictions.

130

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

This the second paper in this review considers three topics. The section on topography development deals mainly with sputter cones on InP because ion bombardment leads to dramatic topographical effects in this material. Ripple development on sputtered GaAs is also briefly discussed. Many publications have dealt with radiation-induced damage in compound semiconductors. Most of these publications deal with ions having energies on the order of 100keV. For these high-energy implantations, radiation-induced damage results cannot be extrapolated to the low- energy regime without some reserve. This is due to the proximity of the surface and the damaged/ crystalline interface, with thermodynamic properties often being different from those of the bulk and the damaged sections. Due to the lack of low- energy damage studies, one is forced to consider the high-energy, radiation-induced damage results in this section. Emphasis is given to two widely different systems with regard to radiation-induced damage: GaAs and HgCdTe. Several mechanisms may cause sputter-induced compositional changes. Evidence indicates that in the case of InP, the compositional changes are due to pure collisional effects (i.e., preferential sputtering). However, in the case of GaAs, several studies have shown conclusively that the sputter-induced compositional changes are due to bombardment-induced segregation of As to the surface of the GaAs, with an As-depleted region just below the surface. In general, few other compound semiconductor systems have been investigated thoroughly. This makes i t difficult to ascertain which particular compositional change mechanism(s) operate in most compound semiconductor systems.

2. RADIATION-INDUCED TOPOGRAPHY

2.1 Introduction

In the investigation of solid surfaces, the most commonly used method to obtain composition depth profiles is the use of a surface-sensitive technique, such as SIMS, or XPS and Auger electron spectroscopy (AES), combined with inert gas sputtering. The quality and accuracy of such depth profiles can be expressed in terms of the so- called depth resolution.2 Because of the complexity of the process of surface erosion by ion bom-

bardment, several factors can influence the accuracy of the r e s u k 2 Ion bombardment-induced surface texturing is the most important factor contributing to the deterioration of the depth resolution.” Kojima et a1.6 found a quantitative correlation between the increase in depth resolution and the increase in rms surface roughness,

Due to its importance, radiation-induced topography development has been studied exten- sively. Several reviews have appeared, including the reviews provided in the books edited by Behrish,7.* Kinakidis et al.? and Auciello and Kelly.Io The emphasis in these reviews and most publications on radiation-induced topography has centered on metals, but a large variety of experimental results obtained under various conditions (which are seldom comparable to one another) have also been reported on compound semiconductors. A number of different models have been suggested to account for the individual findings.

Radiation-induced topographical features on initially smooth substrates range from very small features that can only be seen in STM/atomic force microscopy (AFM) images1 to features that are easily discernible in an ordinary scanning electron microscope. Mainly the latter features are discussed in this section. Furthermore, the main emphasis is on radiation-induced topography development on InP, with just some brief and selective comments on other compound semiconductors. With the exception of a few remarks, the effect of reactive ion beam etching and ion beam- assisted etching on the development of surface topography is not considered.

It is appropriate to define some terminology. Following Ro~snage1.l~ the term cone will be used to denote protuberant features such as cones, pyra- mids, small islands, and whiskerlike structures observed after sputtering.

2.2. Radiation-Induced Morphology Mechanisms

In general, one can classify the different proposed mechanisms for bombardment-induced development of surface topography into two broad categories: sputter erosional theories and growth theories. In this section, the most important of these mechanisms are considered briefly to ascertain their applicability to morphology develop-

131

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

ment on the surface of compound semiconductors during ion bombardment.

2.2.1. Impurity-Seeding Models

Impurity seeding is the classic explanation of sputter cone formation and has been proven to induce surface morphology on seeded metal sur- f a c e ~ . ' ~ ' ~ Two kinds of mechanisms have been proposed to explain the phenomenon. According to the earlier explanation, an impurity having a lower sputter yield than its surroundings will act as the starting point for a cone due to the faster erosion of its surroundings. This surface discontinuity grows to a conical protrusion as a result of the variation of sputter yield with the angle of incidence of the ions. Characteristic for these seeded cones are their circular cross-sections and the groove or trench at the base of a cone.l4.I6 This groove is formed by locally enhanced sputtering due to ions reflected from the cone ~ a l 1 s . I ~ Radia- tion-induced topography development on semiconductor materials usually cannot be explained by this mechanism because the high purity and monocrystalline nature of the semiconductor materials and the cleanliness of the environment (UHV and differentially pumped ion guns) pro- hibit this.

WehneP observed that the seeding material need not have a lower sputter yield than its surroundings but must have a higher melting point. Rossnagel and RobinsonI8 have proposed that the seeded impurity atoms move about on the surface by thermally activated surface diffusion and gather into localized clusters. These clusters than initiate the formation of conical structures that develop with continued sputtering. Weher i s suggested that seed cones are the result of an interplay of whisker growth, surface movement of atoms, and the effects of sputtering. This suggestion is considered again in Section 2.2.5.

2.2.2. Differential Sputter Erosional Models

From Section 3.3.1 in Reference 1, it follows that the sputtering yield varies with the angle of incidence of the ions such that maximum sputter yield occurs in the 60 to 80" range, with the exact

angle being materiaNon d e ~ e n d e n t . ~ ~ . ~ ~ Further- more, the sputtering yield can also vary due to several other effects. These include the spatial distribution of the energy deposition function, the crystal orientation, grain boundaries, and (prebombardment and bombardment-induced) dislocations. Thus, surface topographical features can be enhanced and deformed by sputter etching.I4

In the Sigmund sputter-induced topography model,2' the rate at which material is sputtered from a point is proportional to the rate of energy deposited at that point. This energy deposition function is spatially dependent and thus can lead to the development of surface features. A combination of this effect and the angular dependence of the sputter yield can lead to a further enhancement of original features.

Carter and co-workers (e.g., References 14, 22 to 25) have developed models to describe the development of surface topography due to the erosional effect of ion bombardment. They have been able to predict some surface contour features using the variation of the sputtering yield with ion incidence angle and spatial variations of ion flux densities.

The above differential erosional theories have had reasonable success in their predictions of the radiation-induced topography. However, not all such surface contour features can be described by these theories. For example, any real growth development feature^?^.^^ are not explainable.

2.2.3. Crystallographic Orientation- Dependent Models

On pure copper single crystals, WhittonI6 could always induce sputter cones when sputtering the high index surface (1 1 3 I), but not on other index planes. Etch pits has been observed to be precursors of the pyramidal structures.

Pamler and c o - w o r k e ~ s ~ . ~ ~ , ~ ~ argued that channeling of the bombarding ions is a major cause of the development of sputter-induced topography in polycrystalline materials. As discussed in Sec- tion 3.8 of Reference 1, channeling of the ions will lead to a significant reduction of the sputtering yield compared to that of the nonchanneled ions. Clearly, this mechanism by Pamler and co- workers for topography development is not appli-

132

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

cable for the sputtering of single-crystal semiconductors.

2.2.4. Surface Migration Model

In the surface migration m ~ d e l , ' ~ J * impurity atoms are assumed to diffuse across the surface to eventually nucleate, forming islands or clusters of adatoms. This can result in the appearance of impurity seed cones on the sputtered surface, as discussed in Section 2.2.1.

2.2.5. Ripple Models

The mechanism of surface microroughening induced by sputtering as proposed by Sigmund2' (see Section 2.2.2) and surface diffusion have been used by Bradley and Harper30 to explain ripple formation in the sputtering of some compound semiconductors. This theory predicts that the ripple wave vector is parallel to the surface component of the ion beam direction when the angle of incidence of the bombarding ions is not too large (i.e., near-normal incidence). For near- grazing incidence, the ripple wave vector is perpendicular to the surface component of the ion beam direction. The wave length of the ripples depends on the temperature of the bombarded substrate. There is only some qualitative agreement between this theory and experimental data,26.31.32 especially with regard to the temperature dependence of the ripple wavelength.

Carter et al.25.33 identified two sources for the development of periodic structures on ion-bombarded surfaces: (1) the ordered dislocation array that often occurs in ion-irradiated crystalline substrates and that locally mediates projectile energy loss and surface-binding energy and (2) the over- lapping of high areal density, large individual etch pits, which leads to faceted terraced structures. Calculations by Carter et al.33 gave essentially the same result as those by Bradley and Harper,3o except for a factor of one half difference in the predicted wave numbers of the periodic structure.

2.2.6. Redeposition Model

According to the redeposition model, the presence of seeding impurities on the sample, in the

ion beam, or in the residual gas initiates cone formation. The cones develop further by the redeposition of sputter material and the effect of local differences in ion deflection. This effect becomes more pronounced with oblique ion bombardment. Experimental evidence exists for redeposition effects on cone shape de~e lopmen t .~~ This aspect has even been discussed in terms of a theoretical model.3s

However, there is a widespread ~ p i n i o n ' ~ J ~ - ~ ~ that redeposition plays only a minor role in the development of sputter cones. In fact, at normal incidence, redeposition would tend to broaden the cones and would make a minor contribution to the longitudinal growth of the cones. Also, redeposition only comes into effect when asperities already exist on the sputtered surface.

2.2.7. Whisker Growth Model

Ion irradiation can result in the growth of whiskers on certain target surface^.^^.^^^^ In fact, Wehner15 regards seed cones (see Section 2.2.1) to be the result of an interplay between whisker growth, surface diffusion, and the erosional effects of sputtering. This proposal is contradicted by the findings of Floro et al.3x that impurity seeding inhibits whisker formation. They also found no need for elevated target temperatures. It must, however, also be mentioned that WehnerI5 and Floro et al.38 used different substrate materials.

The main problem of the whisker growth model is the lack of a satisfactory explanation of the growth mechanism of the whiskers. In this regard, the two main suggestions are a kink site saturation mechanism4 and a dislocation loop mechanism. l 5

2.2.8. Microregion Model

In order to explain their extensive experimental findings on sputter cone formation on InP, Gries and MietheZ7 and Gries26 proposed the microregion model. The model assumes ion bombardment to produce an ensemble of micro- crystallites and noncrystalline aggregations of atoms in the InP. The majority of the crystalline microregions are so small that most interstitials, created in collision events between the bombard-

133

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

ing ions and substrate atoms, can reach an interfacial boundary rather than recombine with a bombardment-induced vacancy. These atoms are then transported by interfacial and surface diffusion to special growth points on the crystalline regions of the surface where crystal stubs proceed to grow. In order to account for the successful competition of the growth process with the damage production at the growth sites, these special growth points were postulated to be either screw dislocations or a step source of subatomic height (a Bauser-Strunk dislocation,41) either on a surface crystallite of the target material or on a faulted island. Ion bombardment-induced compressive stresses favor diffusion toward the surface. This crystal growth proceeds until the damage and sputter rate by ion bombardment overtakes the growth rate.

A crucial point in this model, which still needs direct experimental confirmation, is the growth mechanism.

2.2.9. InP Model

InP is a material notorious for extreme topography development during ion bombardment (see Section 2.3). Based on ideas in the microregion modePZ7 and on experimental scanning transmission electron microscopy (STEM)42 and infrared luminescence microscopy e~idence ,?~ . - '~ .~ GrieP traced the topography on InP back to localized epitaxial recrystallization in columnar forms from the radiation-induced amorphous/crystalline interface. These columns grow above the surface, forming a dense network of cones. This topography depends on the angle of ion incidence and on the interplay between reamorphization, recrystallization, above-surface growth, and sputtering.

2.3. Radiation-Induced Topography on InP

A relatively large number of papers have appeared on radiation-induced topography development on InP since the first paper by This is probably due to the pronounced effect in the case of InP and the importance of InP in the optoelectronic field.

Common to all reports is that surface cones were observed in the early stages of ion bombardment of InP wafers at room temperature. Differ- ent shapes of the cones have been reported. Al-

though the quality of the published scanning electron microscopy (SEM) pictures26.27.4742 differs significantly, it seems as if one can divide the bombardment-induced cones into two groups. Most of the reported sputter ~ones~~52.55-59.63 be- long to the first group.

The cones in the first group appear (Figure 1) at a dose of approximately 3 x 1OIs cm-2 as small protuberances (-10 nm) that grow (dose densities -10'5 to 10l6 cm-*) to sizes of 10 to 100 nm (Figure 2). In the low-energy range (E I 10 keV), the sputter-induced morphology has no49 or only a slight dependence on the energy of ions, with a slight increase in the surface roughness at very low (0.5 and 1 keV) energies.63 At a constant ion dose and ion energy, the sizes and density of these cones depend on the angle of incidence of the impinging ions.47.60.63-6S At an ion incidence angle of approximately 41" with respect to the sample normal, a maximum is obtained in topography de~elopment.6~ At glancing angles of incidence, the sputter-induced topography is at a minimum.47.60.64.65 Cones in the first group all developed on as-factory-received InP wafers even though they may have had different surface-cleaning treatments.

Cones in the second g r o ~ p ~ ~ . ~ ~ . ~ ~ . ~ ~ , ~ ~ initially look the same as those in the first group, but the protuberances develop into "spiky" or whiskerlike feature^^^.*^ (Figure 3). Wada47 as well as Malherbe and van der Berg63 showed clearly that these cones, on as-factory-received InP, origi- nated from seeding (redeposition) from the sample holder by impurity atoms (see Section 2.2.1). Sputtering at elevated temperatures also results in whiskerlike cones, which MacLaren et aL3' also interpreted as due to seed cone formation. Be- cause impurity-seeding cones have been studied in great detail on other materials, they are not to be considered further, and concentration is given to cones of the former kind. The InP wafers of Gries26 and Gries and Miethe27 developed nonseeded cones, but differed from the as-factory-received material in that they were preimplanted and annealed (1 x lOI5 Si+ cm-z, 350 keV). Therefore, at least the preimplanted surface would differ significantly from the as- factory-received material.

Most studies of radiation-induced topography on InP have been on as-factory-received samples, with different chemical cleaning steps.

134

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

FIGURE 1. SEM microphotograph depicting the initial stage of cone formation on factory-received InP wafers. The InP was bombarded with 10-keV Ar+ ions to a dose of 3 x 1015 cm-*.

On these samples, no electron and ion beam synergetic effects have been reported. In contrast, G r i e P and Gries and Miethez7 found synergetic effects on previously ion-implanted and implanted- and-annealed samples. Interestingly, there was also a difference between the synergetic effects of the latter two samples.

The development of a ripple-type topography has been reported for room-temperature argon sputtering of InP.26+'7.63 Malherbe and van der Bergb3 have shown that the ripples gradually developed from cones of the first group, with increasing dose density (Figure 4). The orientation of the ripples apparently depends on the angle of incidence of the impinging ions. Sputtering at a large angle of incidence (71" to the sample normal) resulted in the ripples being oriented perpendicular to the surface component of the ion beam

while for 45" sputtering, the ripples were in the direction of the ion beam.26727 Room- temperature sputtering by 0: ions at an incidence angle of 47" 66 and by Cs' as well as by 0; 31 ions at angles of incidence (with the surface normal)

of 25" (for Cs') and 42" (for 0;) also resulted in ripple formation. In neither of the latter two papers were the orientations of the ripples specified. MacLaren et al.31 reported that the spacing between the ripples depended on the sample temperature. No ripples were formed when the InP was sputtered at -50°C while sputtering at 80°C resulted in only cone formation." MacLaren et aL3' interpreted their results to indicate a surface diffusion mechanism for the ripple formation. They found only qualitative agreement between their results and the ripple formation theory of Bradley and Harper.30

The use of Cs' ions has been reported to lead to a suppression, but not elimination, of the cones.z6~z7~66~67 Some studies report that sputtering with the reactive oxygen ions,62.66 or with halogen gas ions such as iodine?* chl0rine,6~ and bro- mine,70 leads to suppression of sputter cones. In other s t ~ d i e s , ~ ~ . ~ ~ , ~ ~ . ~ ~ . ~ ~ - ~ ~ no such effects were observed. Criesz6 and Gries and MietheZ7 reported that flooding their samples with oxygen during sputtering leads to a suppression of cone forma-

135

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

FIGURE 2. SEM image of typical sputter cones of the first group, on InP bombarded with 8-keV Ar* to a dose of 1.5 x 1 0l6 cm-2.

FIGURE 3. SEM image depicting seeding cones (group 2 cones) at a point on the sample near the metal sample holder. Similar images can be obtained on implanted annealed lnP.2627

136

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

FIGURE 4. TEM micrographs of Cr-C replicas depicting the ripple development at glancing angle (71' to the sample normal) Ar* bombardment of InP. The energy of the ions was kept fixed at 5 keV, but the directions of the bombarding ions were different in the two cases shown. The following dose densities apply: (a) 5 x 10l6 A r cm-2 and (b) 2 x 1018 Ar+ cm-2. (From Malherbe, J. B. and van der Berg, N. G., Surf. lnterface Anal., in press. With permission.)

tion. Linders et al.J9 found that air exposure after small ion doses also prevents cone formation. A higher C1, gas pressure in the vacuum chamber caused a smoother surface.69

The above results suggest that the growth of InP cones is to some extent dependent on the chemical state of the sample surface (i.e., oxidized or not) and on the defect structure (i.e., unimplanted, implanted, or implanted and annealed). Cooling the samples to -183°C leads to a suppression of cone formati0n.~',50.5~ This clearly shows that cone formation on InP is not primarily an erosional process, but that the movement of atoms or InP molecules is a prerequisite for cone formation. This is in agreement with the results of

and Cries and Miethe27 found that the cones develop via a growth process. These facts exclude some of the cone formation theories (such as impurity-seeded cone formation, erosional theories, crystal orientation, or redeposition of sputtered species) as likely mechanisms.

There is some dispute in the literature about the composition of these ion bombardment-induced cones. Several publications claim the cones to be composed of pure indium.47s1-54.56,60.68.74-76

The main argument of some of these ~laimS'~J'J6.60.68 is based merely on the fact that P is preferentially sputtered from InP (see Section 4.4), leaving an In-enriched surface with local areas of metallic In present. These metallic In regions are believed to agglomerate to form the sputter cones. The lower areal density of the sputter cones when sputtering

with noble gas ions at glancing angles is regarded as confirmation of this modeV3 because preferential sputtering of P is reduced at glancing-angle ion bombardment (see also Section 4.4.3).

Direct experimental evidence of involvement of metallic In comes from Oliver and co-work- e r ~ , ~ ~ . ~ ~ who observed a high density of fine needles after 0.5-keV Ar+ sputtering followed by a 14-h vacuum thermal annealing at 500°C. The upper parts of the long needles were rich in metallic indium. However, this result is not applicable to room-temperature-sputtered InP. InP readily decomposes under thermal vacuum condition^.^^-^^ Noncongruent evaporation with preferential loss of phosphorus occurs from smooth, flat InP surfaces at temperatures above 366"C.77

Farrow and c o - w ~ r k e r s ~ ~ . ~ ~ , ~ ~ observed sputter cones on (100) InP samples subjected to several cycles of 0.5-keV argon ion bombardment and annealing (250 to 300°C for 30 min). Trans- mission electron microscopy (TEM) diffraction studies showed the cones to be composed of indium. The authors claimed the composition of the cones to originate from the bombardment step. However, in view of the above discussion, preferential evaporation of the more volatile P species from the InP cannot be ruled out. The high temperatures could also have caused the remaining In to agglomerate into cones.

In several papers, evidence is presented in favor of ion bombardment-induced cones on InP consisting of InP rather than In. Skinner et al.57 used TEM with EDX to determine the composition of the cones and found it to be predominantly InP. In a few studies,26*27.31.55 AES was used to determine the composition of the cones. Gries26 and Gries and Miethe2' determined the composition of a cone forest and found it to be InP. However, errors can be introduced into this kind of measurement due to the undulating nature of the analyzed surface. Using selected area Auger sampling, Kirk and Jonesss found the cones to consist of InP, while using scanning Auger microscopy (SAM), found no evidence of the presence of indium. For size ranges down to 50 nm, MacLaren et al." found no evidence of pure In islands. The cones had an InP composition, being perhaps slightly enriched in indium. Following Wilson et al.,*O they interpreted the cone formation as originating from preferential sputtering of P from the InP surface. They believe

137

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

the unbonded In atoms to agglomerate by surface diffusion into In islands, which act as seed points for cone formation (see Section 2.2.1).

A STEM studP2 investigated the composition and structure of low-energy (6 and 10 keV), radiation-induced cones on (100) InP. Conver- gent beam electron diffraction patterns proved the cones to consist of crystalline InP. The investigation showed that the irradiated surface region was not completely amorphous, but that crystalline regions also existed at all applied dose densities from 7 x 1015 to 2 x 10’7 cm-2. Weak-beam investigations revealed that most of the cones have the same orientation, viz., (loo), as the bulk InP. As proposed by the microregion m~del,’~.*~ the cones were found to grow only on these crystalline regions. The agreement of crystalline orientations of cones and bulk led G r i e ~ ~ ~ to formulate the “InP model” as a more likely mechanism.

A few other studies also presented evidence on the structure of the room-temperature, ion bombarded-induced cones on InP. GrieP showed a high-resolution micrograph depicting the crystal habit of the typical sputter cones during their early stages of development (Figure 5). This, as

well as the E M results of Skinner et al.,57 indicated crystalline InP cones. Radiation-induced recrystallization of the sputtered InP substrate was established by IR luminescence m i ~ r o s c o p y . 4 ~ ~ ~ ~ ~ ’

Crystalline InP cones are predicted by the microregion m ~ d e l * ~ , * ~ as well as by the InP

As mentioned in Sections 2.2.8 and 2.2.9, these models were originally developed to explain the mechanism and effects of radiation-induced morphology development on InP. These models can also be applied to other compound semiconductor systems (see Section 2.4).

To summarize, ion sputtering of InP leads to pronounced topography development. The cones are composed of crystalline InP and develop either by a recrystallization-based growth process or by an impurity-seeded mechanism.

2.4. Radiation-Induced Topography Development on Other Compound Semiconductors

Not many authoritative publications have appeared on the topic of topography development

FIGURE 5. High-resolution SEM image of the cone habit. The crystalline edges of the cones are discernible. (From Gries, W. H., Surf. interface Anal., 14, 61 1, 1989. With permission.)

138

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

for compound semiconductors other than GaAs and InP. Furthermore, the fact that accidental seeding of the target surface by impurity atoms occurs readily makes it difficult to analyze the publications that have appeared.

In general, one finds that noble gas ion sputtering of GaAs does not lead to sputter cone formation to the same extent as for InP. When easily discernible (with SEM) room-temperature sputter cones are formed on GaAs, they are attributed to impurities on the surface leading to seed cone formation.u2-x6 Sputtering with noble gas ions at elevated temperatures (225°C) produced an SEM- flat surface. On an STM level, maximum roughness developed at a dose of 6.9 x 1 Ot9 At+ cm-2 for 1-keV Ar+ ions.u7

Radiation-induced recrystallization is also believed to occur in G ~ A s . ~ ) As in the case of InP, evidence comes from infrared luminescence mi- c r o ~ c o p y ~ ~ and from STEM investigation^.^^ Analogous to the InP case,42 this recrystallization in GaAs is thought to be localized and to proceed from the amorphous-crystalline interface right to

the surface (but not, or not far, b e y ~ n d ) . ~ ) Poly- crystalline cones on an At+-bombarded AIGaAs/ GaAs superlattice also indicate a growth process.uu

Sputtering of GaAs with oxygen or cesium ions leads to the formation of a ripple structure perpendicular to the ion beam3'.32.66.*94 (Figure 6). The development of a ripple structure leads to an abrupt change in the sputter rate and a change in the secondary ion yield, which are important effects for SIMS investigations." In contrast to the latter result by Gericke et al.," an inspection of the SIMS depth profiles of AlGaAs/GaAs multilayers shows no such changes.32 The wavelength of the ripples has been reported to depend on the GaAs substrate temperature, angle of incidence of the bombarding ion, current density, ion dose, and ion en erg^.^^.^'.^.".^'.^^.^ When cooling the sample down to -30°C. no ripple formation was found after Cs' bombardment.)' MacLaren et aL3' interpreted the results to indicate that ripple formation is due to diffusion effects. They compared their results to a ripple formation theory by

. FIGURE 6. Ion bombardment of GaAs usually leads to the development of a "ripple-like" surface morphology, with the "ripples" being perpendicular to the angle of incidence of the ions. (From Gries, W. H., in Proc. SlMS V/ / / , Benninghoven, A., Janssen. K. T. F., Tiimpner, J., and Werner, W. H., Eds., John Wiley & Sons, Chichester, 1991, 323. With permission.)

139

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

Bradley and Harper30 and found only qualitative agreement. Using a slightly different system (i.e., GaAs/AIGaAs superlattices), Miethe et aL3* found that the theoretical predictions of the Bradley and Harper model are not substantiated by experimental results. They found that the angle of incidence is the most important parameter for ripple formation. The most pronounced effect was observed at 35” to 40” for 0; and 50” for Cs+ bombardment. Higher primary ion energy increased the ripple spacing, while the angle of incidence had no effect (+20%) on the ripple spacing. Heat- ing the sample to approximately 300°C resulted in smaller ripple spacing. O t h e r ~ ~ , ~ ~ reported results similar to those of Mie’the et al.32 regarding angle dependence.

G r i e ~ ? ~ argued that the ripple formation mechanism should be interpreted in terms of radiation- induced surface deformation combined with recrystallization of GaAs. The recrystallized subsurface “columns” serve as anchor points for individual ripples and thereby influence or determine the ripple spacing. Recently, AFM measurements gave support to this model.9’

As mentioned above, not enough papers by different groups have appeared on radiation-induced topography development in compound semiconductors. The situation varies considerably depending on the semiconductor in question. Thus, i t varies from several publications by a single group (e.g., the dichalcogenide WSe?s97) to publications by several groups on a particular semiconductor, such as GaAs and InP. There- fore, this section concludes with a few general remarks.

Noble gas sputtering of compound semiconductors containing In often leads to the appearance of sputter cones on the In many compound semiconductors, sputtering with 0; and Cs+ ions (which are important sputtering species in SIMS) leads to a more pronounced development of surface topography (in the form of cones or ripples) than for noble gas sputtering. In some materials (e.g., InP and InSbw), sputtering with these ions leads to a suppression of sputter-induced topography.

To summarize, one can only echo the conclusion of a recent summary of radiation-induced topography on semiconductor^:^^ this phenomenon, although long researched, is still relatively

poorly understood and hence difficult to predict. At present, this field is still in the fact-gathering stage.

3. RADIATION-INDUCED DAMAGE

3.1. Introductory Remarks

Radiation-induced damage in compound semiconductors has been one of the most intensively studied fields of ion-semiconductor interactions. Most of these studies were concerned with high- energy (20 to 200 keV) ions. Literature specifi- cally dealing with damage caused by low-energy ion bombardment is scarce. Low-energy bombardment gives rise to a thin damaged region at the surface. The proximity of the surface and the damage/crystal interface make a direct extrapola- tion of high-energy results to lower ones inaccu- rate in many cases.

The measurement of radiation-induced damage by low-energy ions is difficult. Many techniques have been used. An accurate and direct method is cross-section TEM. However, it suffers from the difficulties of sample preparation. A popular and straightforward technique is particle channeling combined with Rutherford backscattering (RBS). Channeling-RBS measures the difference between the backscattering yields of the investigating ion beam (incident along an “open channel” direction of a single crystal substrate) before and after damage occurred. By comparing the yield with that of a random (amorphous) substrate, the degree of damage to the region can be deduced. Using channeling-RBS, one can esti- mate the depth and thickness of the damaged region. With respect to low-energy radiation damage detection, channeling-RBS has one major drawback: due to energy straggling of the investigating ions (usually H+ or He+ ions), the depth resolution of this technique deteriorates quickly once the surface is penetrated.

Because the optical properties of semiconductors are also changed by radiation-induced damage, the damaged layers are often investigated by optical techniques such as Raman spectroscopy, photoluminescence, ultraviolet (UV) reflectivity, ellipsometry, etc. Likewise, a close correlation exists between structural damage and

140

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

changes in electrical characteristics of semiconductors. Therefore, methods such as DLTS, I-V, C-V-(w), the Hall effect, etc. have been employed to study radiation damage and profiles. Low-dose irradiation affects the electrical properties of semiconductor materials before any changes can be detected by methods such as TEM and RBS/channeling in the structural properties (such as amorphization, point and extended defects) of the semiconductor

Most of these characterization techniques are lacking in at least one of two aspects: (1) the depth resolution and quantification are poor, making it difficult to obtain accurate information on the interfacial region between the damaged layer and the undamaged single-crystal substrate, and (2) the sensitivity is often too poor to detect small amounts and regions of amorphization. It is therefore not surprising to find that there are many contradictory results regarding the extent of radiation-induced damage and/or amorphization in heavy ion-bombarded compound semiconductors.

The mechanism of ion-induced conversion of a crystalline state to an amorphous state is still being debated.Io3 Two models for the amorphization process have been proposed in the literature: ( 1 ) the heterogeneous model, which suggests that the individual damage clusters are amorphous and that complete amorphization occurs as a result of the accumulation and merging of individual damage clusters, and (2) the homogeneous model, which suggests that when the defect concentration reaches some critical value in a single crystal, that crystal becomes unstable and transforms to an amorphous state.Iw

3.2. Radiation-Induced Damage in Some Compound Semiconductors

3.2.1. GaAs

Some aspects of low-energy radiation-induced damage in GaAs were discussed in Section 3.4.1 of the first part1 of this review (with respect to ejection patterns in sputter yield measurements), in Section 2.4 (with respect to bombardment- induced recrystallization during bombardment- induced topography development), and also in Section 4.5 (ion bombardment-induced composi-

tional changes). These particular aspects are not discussed again. Here, we focus on structural damage. Very little attention is given to the influence of bombardment-induced changes to the electronic and optical properties of GaAs. In the case of GaAs, the threshold dose for changes of these two p r o p e r t i e ~ ~ ~ J ~ ~ - ~ ~ is much lower than for structural damage when measured by the conventional techniques, for example, TEM, channeling, etc.,Io7 in agreement with other semiconductors.10G111 The damaged regions measured by most electrical or optical methods usually extend much deeper than those measured by TEM or channeling, due to channeling of a small fraction of the bombarding ions.

Because radiation-induced damage in GaAs has been reviewed many times (e.g., References 112 to 121), only some general trends are briefly discussed. In this discussion, only a few references are given, concentrating on more recent ones. First, a warning: damage information gath- ered at high-energy bombardments does not necessarily pertain also to low energies. The structures are different for these two energy regions.'" An example of this is shown in Figure 7, depicting bright-field TEM images of GaAs, bombarded with 10- and 20-keV Si' ions to the same dose density.Iz3 The samples were subjected to rapid thermal annealing at 900°C for 12 s in arsine. There is clearly a tremendous reduction in the dislocation loop density in the lower energy-implanted sample.Iz2

There has been some dispute about the creation of amorphous layers in GaAs by high-dose, low-energy implantation using heavy-mass ions (e.g., References 114, 124, and 125). Most studies now indicate that GaAs is amorphized by low- energy (0.1 < Ep < 5 keV) ion bombardment (e.g., References 103 and 126 to 132) as well as by high-energy ion irradiation (e.g., References 133 to 138).

For low-energy noble gas sputtering, the damage depths are comparable to or slightly exceed the projected ranges of the ions in G ~ A S . ~ ~ ~ J * ~ J * * For the reduced energy range (for the definition of E, see Equation 59 in Reference 1) 0.3 c E < 2.8, a linear relationship has been found139-140 between the damage depth and the ion energy to the power 3 of the bombarding AP ions (Figure 8). For lower-energy (1 to 3 keV) argon bombard-

141

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

10 keV 20 keV

FIGURE 7. TEM images showing the reduction in dislocation loop density when using lower implantation energies in GaAs. Two-beam, plan-view TEM, bright field, grim, s > 0. (From Rubart, W. S., Jones, K. S., Seiberling, L., and Sandana, D. K., Mater. Res. SOC. Symp. Proc., 157, 677, 1990. With permission.)

I

U a , I U !

AI: A:, T d - GaAs

/ f '

/*

4/'

- = 0.60 Cu3

01 . I 0 1 2 3

Reduced Energy

FIGURE 8. Reduced damage depths of GaAs implanted with A r . Al+, and Te+ ions as a function of the bombardment energy in reduced units. (From Kido, Y. and Kawamoto, J.. J. Appl. Phys., 58, 3377, 1985. With permission.)

142

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

ment of GaAs, Konomi et a1.I2* found a linear relationship between the damage depth and the square root of the energy of the bombarding Ar' ions. This energy range corresponds to E = 0.088 to 0.27.

Both these relationships agree with experimental measurements of the dependence of the projected range on the primary energy of the ions. In terms of reduced units (i.e., E and pp) for the primary energy E and the projected range R,, Oetzmann et al.14'.i42 found that:

1 .0E'12 for ~ ~ 0 . 0 2 for 0.02 c E < 0.3

2.2& for 0 . 3 ~ ~ < 6 p, = 1 Jk213 ( 1)

using targets of amorphized Si, Ge, and A1 bombarded with As, Ge, Sb, Au, and Bi ions. This also agrees to some extent with the LSS theory, which predicts that for 0.002 < E I 0.1 143

An analogous result regarding the relationship between the range of the bombarding ions and the damage depth was found by Tognetti et a1.Iu They investigated the amorphization layer thicknesses of GaAs, InP, and Si caused by implantation of various ions with different energies, and found that this thickness X, can be well described by the equation

Xa = X, + (2.7 k O.5)AXd ( 3 )

where X, is the mean damage depth and AXd is the standard deviation of the damage profile as calculated by Winterbon.'45

Ion channeling naturally results in deeper damage.146J47 Electrical and optical defects usually greatly exceed the projected ranges of the ions.103.126 This has led to the suggestion that low- energy implantation creates a top amorphous layer followed by a transition region, where a fine- grain mixture of the microcrystals and amorphous GaAs coexist. The latter region gradually disappears in the bulk.Io3

In general, one finds that low-energy ion implantation of crystalline GaAs at room temperature leads to the creation of point defects, such as Frenkel pairs, consisting of a vacancy and

the displaced atom.II3 Although more complex defects are also present, the simple defects are more common. After annealing, the defects tend to conglomerate into more extended defects, such as dislocation loops. The annealing of continuous amorphous layers proceeds by solid-phase epitaxy. The growth front proceeds toward the surface. These layers generally show poor recrystallization characteristics and exhibit a high degree of residual disorder in the form of microtwins, stacking faults, and point defects.113.i I6 Thermal treatment in the range of 400 to 500°C results in defects annealing, leaving a high density of dislocation loops. These loops grow and annihilate at higher temperatures (-700°C). The remaining point-defect clusters anneal at higher tempera- t u r e ~ . " ~ Annealing by other methods, such as pulsed electron beam annealing (PEBA)133-'35 and rapid thermal annealing (RTA),1'7.148 also leads to the formation of dislocation loops and tangles (Figure 9). High-temperature capless annealing (furnace as well as faster methods, such as PEBA) can lead to local deviations from stoichiometry due to incongruent evaporation of As.lJ9 Disso- ciation of GaAs can be prevented by annealing in flowing ASH, gas.150.151

The presence of residual damage after annealing clearly shows the need to avoid amorphization of GaAs during implantation. Success has been achieved by implanting at elevated temperatures (2150"C).116.152 The increased mobility of point defects at higher temperatures can lead to dynamic annealing of the cascade regions. An elevated temperature anneal has the added advantage of a higher activation of the dopants. This is a significant advantage because it is more difficult to obtain high dopant activities in GaAs compared with

The question of the mechanism causing amorphization is still not completely solved. Con- flicting evidence is the prime reason for not being able to choose between the homogeneous or heterogeneous model (see Section 3.1). The mechanism also seems to depend on the mass and energy of the ions used. To illustrate this point, only a few examples are given. Channeling-RBS and optical transmission measurements on weakly damaged GaAs by 150-keV B+, 200-keV Ax+, and 300-keV Zn' ions support the homogeneous defect nucleation model.Is4 A similar conclusion can be drawn from high-dose, high-energy (200

143

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

FIGURE 9. 150-keV Cr+-implanted GaAs after different pulsed electron beam annealing conditions. (a)-(c) PEBA at 0.5 J cm-2. The amorphized GaAs started to recrystallize into grains with different orientations, (a) (1 12) and (b) (1 14) zone axes, because the melt front had not yet penetrated the damaged layer to allow epitaxial growth from the (001) substrate. (c) TEM micrograph shows that the structure consisted of large (>1 pm) grains containing dense dislocation tangles. Scale bar = 400 nm. (d) PEBA at 1.3 J cm-2. At energy densities >0.72 J cm-2, the melt front had penetrated the damaged layer. The TEM micrograph shows a cellular structure with Ga precipitates, segregation cells, and areas of high dislocation densities. Scale bar = 400 nm. (From Gaigher, H. L. and Alberts, H. W.. Radiat. Eff., 125, 373, 1993. With permission.)

to 330 keV), heavy-ion (Zn, Se, and Sn) implan- t a t i ~ n . ~ ~ ~ On the other hand, evidence from infrared (IR) luminescence microscopy on GaAs bombarded with argon and oxygen ions of low energy (several keV)81 supports the heterogeneous model in that the graphs of IR luminescence intensity vs. ion dose density are smooth over the entire range from 1O1O to lo'* cm-2, without any indication of a transformation discontinuity.Is6

As discussed in Section 2.4, some ion bombardment-induced topographical models need crystalline regions in the sputter-damaged section of the bombarded substrate. Apart from the evidence of a bombardment-induced recrystallization growth mechanism presented in Section 2.4, there is direct evidence of this phenomenon oc- curring in GaAs (e.g.. References 157 and 158). especially for light ion species.

3.2.2. InP and Other Ill-V Semiconductors

In general, most of what was said about damage in GaAs in Section 3.2.1 can be repeated muratis mutundis for InP. There are a few differences between GaAs and InP; for example, InP has a lower threshold damage density (dimen- sions of energy/volume) for amorphization; it is more easily amorphized, etc. 159 Several recent reviews have been published on radiation-induced damage in 1nP.I 13.1 14.152~164.161

Although radiation-induced damage in GaAs, and to a much lesser extent in InP, is well researched and reasonably well understood, the same is not true of other 111-V semiconductors. More systematic studies are needed to give an overall picture.Is'

In a systematic study, Jones and Santana,lsy using TEM, determined the threshold damage density for amorphization for a number of compound semiconductors. The results are shown in Figure 10. The numerical value for GaSb might be slightly erroneous due to the neglect of changes in the density of the substrates. (GaSb swells due to radiation-induced damage.162-'6s) The studyI5' clearly shows that the 111-V semiconductors are amorphized by ion bombardment, but that there are significant differences between different semiconductors regarding the energy deposition required for amorphization.

Figure 11 confirms the different behavior of the different 111-V semiconductor^.^^^ The arsenides clearly require higher deposited energies in the crystal lattice to obtain a damage concentration equal to that in the phosphides. The slope of the arsenide curve in Figure 11 is also steeper than that of the phosphides. It is interesting to note that the behavior of the elemental semiconductors Si and Ge is close to that of the phosphides. This indicates that the partially ionic bonding in the phosphides has no influence on damage production. The high mobility of point defects

1 44

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

1019l I I I I I I I I

1.0

0.8

0.6

0.4

0.2

0

Si AlAs InP lnAs GaP GaAs GaSb ZnSe

I I I , , I , I , I

J = 8 0 K 200 k e V Ar+

0 -A - -

/ + T 7 7 lnAs

I x GaAs I InP

A GaP

+ Ge a

-

7 . -

A

dB X I

0 I Q S i /Y9+ I

A ? " , , , X7./

- - I

- -

I , , I I t , I

Semiconductor

FIGURE 10. Threshold damage densities for amorphization for several (mostly Ill-V) semiconductors. Ion bombardment was performed at 77 K using 20-keV Si+ to a dose of 1 x 10'5 Si* cm-2. (From Jones, K. S. and Santana, C. J., J. Mater. Res., 6, 1048, 1991. With permission.)

FIGURE 11. Normalize_d integral defect density NdN,"ax as a function of the relative average vacancy concentration N,JN,, obtained from TRIM simulations, for 200-keV Ar+ implantation at 80 K into various Ill-V and elemental semiconductors. (From Wesch, W., Nucl. Instrum. Methods, 868, 342, 1992. With permission.)

145

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

seems to be responsible for the substantial influence of defect annealing and transformation in the arsenides, whereas in the other materials a higher stability of the primarily produced defect clusters can be assumed.152 The crystalline-to-amorphous transition apparently proceeds via an accumulation process, and in the arsenides via a collapse- like p r o ~ e s s . ~ ~ - ' ~ ~ For GaAs, this last generalized statement is not necessarily correct because some results contradict this (see Section 3.2.1 for a brief discussion).

To summarize, although the 111-V semiconductors exhibit many common features with respect to radiation-induced damage, there are significant differences. More systematic studies are needed in which the behavior of the different semiconductors is compared.

3.2.3. Sic

Silicon carbide is a very popular material and has been the subject of many ion bombardment studies (e.g., see Reference 166 for a recent review). This is due to the fact that it has applications as an indirect wide-band-gap semiconductor as well as a nonoxide ceramic. It exists in over 100 polymorphs. The two forms of main interest for electronic application are a- and p-Sic. a-Sic consists mainly of the 6H polytype, while the p- or cubic phase has the zincblende structure.IM

Silicon carbide is easily amorphized by ion implantation at room temperature (e.g.. Refer- ences 166 to 182). After annealing, polytype transitions can take place.183,1x4 Implantation at elevated temperatures (-500°C for p-SiClXs and -750°C for a-SiCI7' leaves the S i c monocrystalline during implantation, albeit with residual damage. At room temperature, amorphization occurs at energy densities of about 16 to 20 eV per atom for a-Sic and about 25 eV per atom for p-Sic, which corresponds to 0.2 to 0.3 dpa (displace- ments per atom). RBS/channeling studies indicate that amorphization starts at a depth corresponding to the maximum energy deposition position and spreads out with increasing ion dose. The accumulation of damage is approximately linear with fluence until amorphization occurs. 175~179 This amorphization is accompanied by swelling of up to 30%.17'

Furnace annealing below 1400°C results in a very slow epitaxial regrowth of amorphous Sic. Above 1 500°C, damage annealing becomes very rapid in a narrow temperature interval180 (Fig- ure 12).

Interesting mechanical and chemical changes occur during ion bombardment. The micro- indentation hardness at first increases during the early stages of damage accumulation but then decreases after amorphization to about 60% of that of the crystalline counterpart.I7l Implantation also causes reduction in the friction coefficient as well as in the wear rate (for details and explana- tions, see Reference 186). The oxidation and chemical etching rates of the amorphous state are higher than for the crystalline state.i74

According to McHargue and Williams,'M the amorphization kinetics, annealing kinetics, and property changes during ion bombardment are broadly compatible with the homogeneous model for amorphization (i.e.. the critical accumulation model).

3.2.4. II- VI Compounds

The defects created by ion implantation in 11-VI compounds are different from those created in 111-V and IV s e m i c o n d ~ c t o r s . ~ ~ ~ As discussed in Sections 3.2.1 and 3.2.2, the following generalized properties apply to radiation damage in 111- V (and, in principle, also for the IV) semiconductors. In covalent semiconductors with high binding energy, radiation-induced defects are usually localized within the region where ions deposit their energy by nuclear collision. These defects are primarily point defects. In most cases, it is also possible to amorphize the substrate by high-dose implantation.

11-VI semiconductor compounds are charac- terized by low binding energies and high ionicity. Extended defects (dislocation loops) are often observed. These defects can propagate well beyond the depth at which ions are stopped.188 Due to ionic forces, recombination of defects is more likely to happen, and amorphization of the crystalline material is not readily obtained. Conse- quently, even for large implanted doses of heavy ions, the density of defects, observed by channeling techniques or TEM, does not increase indefi-

146

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

ORNL-DWG 86-1 1308

0.35

h

E 0.30 =t

> 0.25

v

U w

J

u)

0 I a.

2i

a

3 0.20

0.i5 a

8 0.10 I F 0 - 3 0.05

0

t

52Cr (260 keV, ixiOi6/crn2) IN a-SiC

1 I

0 200 400 600 800 i000 4200 4400 4600 ANNEAL TEMPERATURE ("C)

FIGURE 12. Width of an amorphized Sic layer as a function of isochronically annealing temperatures (10 min at each temperature ). a-Sic was amorphized by 260-keV implantation of Cr+ to a dose of 1 x 10l6 cm-2. (From Bohn, H. G., Williams, J. M.. McHargue, C. J., and Begun, G. M., J. Mafer. Res., 2, 107, 1987. With permission.)

nitely and saturation is reached. Similar results are also found for the 11-VI insulating. ionic compound MgO. IX9

Because CdTe is a typical 11-VI semiconductor with regard to its radiation damage behavior, a very brief discussion of this behavior is given below. It has been reported that neither high-energy implantation (e.g., References 190 to 195) nor low-energy implantation (e.g., Ref- erence 196) leads to amorphization of monocrystalline CdTe, although considerable damage in the form of extended defects is caused to the crystal structure. The damage layer extends well beyond the calculated projected range of the bombarding ions, indicating a high mobility of the defects. The deep radiation damage in CdTe is explained by enhanced diffusion of the defects

due to dislocation loop-induced s t r e s ~ . ' ~ ~ . ' ~ ~ For a more extended discussion of the mechanisms involved in deep radiation damage in a related material (i.e., Hg,-,Cd,Te), the reader is referred to Section 3.2.5.

Even for high-energy bombardment, a low- damage surface layer is observed in CdTe. Its thickness is independent of the ion projected

This indicates a high efficiency of the surface to act as a defect sink.193 The results by Lu et who reported an absence of a LEED pattern after 1-keV Ar+ bombardment, can probably be explained by this defect annihilation behavior of the CdTe surface.

Implantation into CdTe can cause electronic- type conversion due to radiation-induced defects (e.g., References 198 and 199).

147

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

3.2.5. Hg,,Cd,Te

Due to its importance as a narrow-gap semiconductor material, the effect of ion implantation into Hg,-,Cd,Te (MCT) has been studied exten- sively. Details of this effect can be found in the review papers by Bubulac,200 DestCfanis,20'.'o' and Bahir and F i n k n ~ a n . ~ ~ ~

3.2.5.1. Structural Damage

The structural damage caused by the bombarding ions in crystalline Hg,-,Cd,Te is in many respects similar to that caused by implantation of single-crystal metals. The main types of structural radiation-induced damage that are observed are extended defects, viz., dislocations and dislocation loops. The first direct observation of structural defects was made by Bubulac et al.,?ol using TEM. They showed that for boron implants, vacancy-type dislocation loops could be observed. With mercury ion implantation, stacking faults were f o ~ n d . ? ~ ~ . ? ~ ~ These were observed even at a depth equal to ten times the mean range of the mercury ion. Subsequent TEM s t u d i e P showed interstitial-type dislocation loops induced by boron implantation. These defects appear in the implanted region only for large implanted doses

cm-?). For smaller doses ( loi4 cm-'), no identifiable defects were observed by TEM. This suggests that dislocation loops are created by the coalescence of simpler defects. The channeling- RBS spectra for implanted MCT207-2w are also characteristic of extended radiation defects. This is in contrast to the spectra obtained for 111-V compound semiconductors as well as for Si or Ge, where primarily point radiation defects dominate.

Another very important property of the radiation-induced damage in MCT is the fact that the damage reaches a saturation value before amorphization. This indicates that there is a strong recombination of the defects during the bombardment process. The defect saturation effect for irradiated MCT was already discovered in the first damage studies in 1980 by Bahir et al.,207.208 using the channeling technique combined with RBS. This saturation of defects was subsequently con- firmed by several other studies (see Reference 201 for details).

The above similarities with radiation-induced damage in metals make it easy to explain the development and saturation of radiation damage in MCT using theories developed for radiation damage in metals. The primary effect of ion bombardment is the production of point defects in a crystalline substrate. As the bombardment dose steadily builds up, there is an increase in the number of vacancies and interstitials, and they coalesce to form clusters. These interstitial or vacancy clusters gradually form dislocation loops of the interstitial and vacancy types. With increasing dose, the loops increase until they inter- act and eventually form complicated dislocation tangles. Further irradiation will simply result in the rearrangement of these tangles by the process of slip and climb but without an increase in defects.

Another similarity with radiation-induced damage in metals is the creation of deep radiation damage (i.e., damage far beyond the range of the bombarding ions) in MCT. For light implanted ions such as P,?07 B,?") and Ar,?07 the damaged region is near the penetration depth of the ions, whereas for heavy ions such as Hg207.211 and In,2m.212.?13 the damaged region may spread much deeper than the projected range of the ions. For mercury implantation, radiation-induced damage (stacking faults) is observed at a depth equal to ten times the mean range of the Hg ion.'"? This difference between heavy and light ion irradiation suggests that the deep damage might be spike dependent. The extension of the damaged zone to depths much deeper than the ranges of the implanted ions is also found in irradiated single- crystal fcc For metals, the proposed mechanism for deep radiation damage is based on the assumption that the bombarding heavy ions create spike regions in the metal ~ u b s t r a t e . ~ ' ~ . ~ ' ~ . ? ' ~ As discussed in Section 2 of the first part' of this review, the very high temperatures in a spike region cause an expansion in the spike region, resulting in a strong stress gradient. Such stress gradients push the dislocations deeper into the single crystalline substrate. The efficiency of such a process would depend sensitively on the Peierls- Nabarro force, which has to be surmounted in order to move a dislocation. This Peierls-Nabarro force is material, crystal-structure, and crystallographic-orientation dependent. The Peierls-

148

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

Nabarro force is very small in fcc metals, explaining the above deep damage effect in such met- ais.217

3.2.5.2. Annealing of Hg,+ Cd,Te

A suitable annealing process for MCT is of paramount importance to repair implantation damage and to activate implanted impurities. However, annealing of MCT is not as easily achieved as with most other semiconductor materials. The primary problem is that Hg is lost during the annealing step. Although low-temperature (5200°C) postimplant anneals can remove some of these defects without too severe a loss of Hg, the effect is limited.*19 In general, one finds that most of the damage can only be annealed out at temperatures from 250 to 35OoC, in which case Hg loss becomes a major problem. Conventional furnace annealing can only be done under a satu- rated vapor pressure of mercury, usually obtained by sealing the MCT in a quartz tube.*20.2?1 Alter- natively, the Hg loss can be reduced by employ- ing methods with short annealing times such as RTA2?I-?l3 and laser annealing.207.”2“ However, a capping layer is also often necessary with these methods.

Another problem is the rapid diffusion of some implanted ion species, such as indium, during annealing.*” By comparing the range profiles of implanted indium in MCT at room temperature and at 77”K, Magel and S i g m ~ n ” ~ concluded that the longer tails of the implanted profiles at room temperature indicate that radiation-enhanced diffusion caused the long tails. Implanted b~ron,’~’ . ’~~ and aluminumZm showed very little diffusion.

Implantation causes the chemical bonds of some of the substrate molecules to break. Thus the main element diffusing in the postimplant anneal from the implanted source is displaced Hg from the lattice.228

Other phenomena may occur in or near the implanted region during the annealing step, that is, redistribution, segregation, or gettering of some of the impurities, such as Na, through the entire implanted layer. For MCT, this effect was first observed by Bubulac et al.229 for Li implantation into MCT when Li was present at levels higher

than - lOI3 cm-3. DestCfaniszol observed similar behavior for sodium impurities in boron-implanted samples. Usually there is a decrease in alkaline impurity concentration beyond the ion-implanted profile. This behavior has been e ~ p l a i n e d ? ~ ~ . ? ~ ~ by a decrease in vacancy concentration beyond the path of the ions. Fast-diffusing alkaline impurities should decorate mercury vacancies, and their profiles should be related to a decrease in metal vacancy concentration associated with mercury in-diffusion. The source of the mercury atoms needed to fill vacancies could originate from the damaged region, where many free interstitial mercury atoms may be formed during the ion collision cascade. This effect is assumed to be important in junction formation mechanisms.

3.2.5.3. Electrical Damage and Junction Forma tion

Attempts to form n-p junctions on MCT in a controlled manner using ion implantation of a chemical dopant have resulted in only limited success. Problems are encountered that usually do not occur in silicon, germanium, and III-V semiconductors. All the studies dealing with ion implantation of MCT have shown that radiation damage is of major importance in junction formation. In particular, the following two effects dominate:200-203

1. Unannealed implanted samples always exhibit n-type conductivity. This is attributed to radiation damage, such as radiation-induced point defects, free point defects such as Hg atoms, residual impurity or implanted species, deep-level donor complexes,231 and resonant levels.??* The n-type region can extend up to a few microns further than the stopping range of the implanted specie^.^^^^^.^^^.^^^

2.

The electrical damage induced by ion bombardment is introduced at very low implantation dose densities.232 For boron implantation, electrical measurements show a strong n-type doping of lo1* ~ m - ~ for such small-dose densities as lOI3 cm-2.201 At such low doses, no structural damage can be identified by TEM or RBS. (Using

1 49

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

heavier ion species, such as In, the structural damage can be detected even at dose densities on the order of 10l2 cm-2.2w) The dislocation loops detected by these methods at higher dose densities are electrically inefficient and are the cause of low carrier mobility before annealing.

Because the implantation process leads to an n-doped layer that depends very little on the implanted species, implantation conditions, and initial doping of the crystal used, in some cases, satisfactory n-p junctions have been achieved using implantation with low-energy (- 1 keV) ions (see References 203 and 233 for references). The main element diffusing in the postimplant anneal from the implanted source is displaced Hg from the

Thus, the junctions form away from the damaged region. In some undesirable cases, diffusing species such as defects, point defects, and implanted atoms diffuse through the extended defects of the material, causing the formation of a graded n-type electrical profile. The predomi- nant diffusing species will depend on the processing conditions and on the defect structure of the layer.

The use of a true chemical dopant is prefer- able because it allows precise control of doping concentration profiles, type of junction, accurate prediction of junction position, and reduced number of noise sources. Notwithstanding the abovementioned problems, n-p or p-n junctions have been obtained after annealing of implantation defects and activation of implanted impurities

arsenic.226 Results concerning boron212~230~234~231238 are somewhat controversial, although most studies claim electronic activation after annealing.

Radiation damage can also be used in making ohmic contacts on MCT. As mentioned above, due to the Hg loss problem, low annealing temperatures are used. This causes the damage in the near-surface region not to be totally annealed out. The remaining near-surface n+-type region serve as a good ohmic contact.

The complex phenomena associated with electrical damage were schematically illustrated by Bubulac,2m reproduced here as Figure 13. Curve 1 represents a typical implanted species atom concentration profile, while the corresponding

such indium,?l2.? 13.221 phosphorus,2?3.~2J.234,235 and

;MMO@ILE EL. ACTIVE OEFECTS . IMC SPECIES RELATE0 - 3 - :MPUNl'€O SPECIES ATOM

CONCENTRA~ION 8 - CARRIER CONCENTnAnON

4S-:MCUNTEO

3 - CARRIER CONCENfl lAnON ANNLILEO

I -HOEILL E L ACTIVE 9EFECTS

8 M P U W T U Y E R OLPECT REUTEO

c OIFCUSION SOURCE FOR

1 I 3

Wg VACANCY

IAOIAl lON OAYAGL . ENHANCED DIFFUSION - N f 3 ACTIVATED 3 OE- (urn1

ACTIVE DEFECT3 E L E C T n I U U Y IEXKNDEO OEFCCTSI 'MWRIlV I

. IMPUUATW REOISTRIWTIOW. SEGREGATION. GETERING - ~ ~ - J A C * N C Y EvoLunom

FIGURE 13. Phenomena associated with ion implantation of MCT. (From Bubulac, L. O., J. Cryst. Growth, 86, 723, 1988. With permission.)

150

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

junction electrical profile before annealing is given by curve 2 and that after postimplant annealing by curve 3. The as-implanted electron concentration profile, as shown in curve 2, exhibits a double peak feature. The shallower peak represents the immobile, electrically active defects related to the implanted species, and the deeper peak represents the mobile, electrically active defects that pen- etrate much deeper into the MCT substrate. This profile marks the region of a diffusion source that operates in the postimplant anneal.

The insight gained from studies on the above mechanisms of n-p junction formation led to the development of a qualitative model by Bubulac.200 This model considers ion implantation as a limited diffusion source of either Hg or the implant species, and not as a direct source of n-p junction formation. Consequently, two substantially different ion implantation techniques were developed: a traditional technique in which the displaced Hg diffusion source operates to form n-on-p junctions and a classic technique based on activation of the implanted species. The implanted species, as diffusion source, can act to form p-on-n or n-on-p junctions. This view of the above mechanisms of n-p junction formation was used to iden- tify the conditions required to place the junctions several microns deep into the epitaxial layer, where the material is of high quality and essentially damage free. The electrical profile of the junction could be designed for specific applications.2m Thus, junctions formed by electrical activation of implanted impurities in an impurity-doped background are now possible in both configurations, n-on-p and p-on-n.

3.3. Concluding Remarks

Knowledge of radiation damage and annealing properties in semiconductors is of utmost importance to the electronic devices industry. This explains the extensive coverage in the literature. Nevertheless, there are still many problems to be solved.

One such topic, which was only briefly considered in this section, is that of heavy ion bombardment-induced recrystallization at high doses. This field is becoming increasingly important from both a technological point of view and a funda- mental scientific point of view.

The radiation-induced damage behavior of each semiconductor family differs dramatically from that of the others. In the predominantly covalent 111-V semiconductors with high binding energy, radiation-induced defects, as detected by TEM or channeling, are usually localized within the region where ions deposit their energy by nuclear collision. These defects are primarily point defects. In most cases, it is also possible to amorphize the substrate by high-dose implantation. In contrast, the 11-VI compounds are not amorphized by room-temperature implantation. Instead, extended defects (dislocations loops) are observed. This can be attributed to low binding energies and high ionicity in 11-VI compounds. In the case of large implanted doses of heavy ions, the density of defects reaches a saturation level before amorphization sets in. Radiation damage extending well beyond the ion range is observed. Electronic-type reversal can occur after implantation of the crystalline substrate.

4. COMPOSITIONAL CHANGES

4.1. Introduction

Sputtering of multicomponent materials often leads to compositional changes in the selvedge. There are many obvious reasons why accurate knowledge of the extent and mechanisms of sputter-induced compositional changes in compound semiconductors is of great importance to technology. Compositional changes become even more important with the trend toward shallow doped layers, as discussed in Section 1 of the first part' of this review.

Many factors can cause such compositional changes. A brief review of these factors is given in this section as well as an in-depth discussion of noble gas ion bombardment-induced compositional changes in GaAs and InP, followed by a brief discussion of such changes in some other compound semiconductor systems. Once again, compositional changes in a target material due to chemical sputtering effects are excluded.

4.2. Preferential Sputtering

Early on, all sputter-induced compositional changes in multicomponent targets were explained

151

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

in terms of preferential sputtering (see Reference 239) for a short historical review). Several review papers'0*239-24s have been published on preferential sputtering. Most reviews have concentrated on the preferential sputtering of binary alloys, oxides, and insulators. Little mention is made of preferential sputtering effects in compound semiconductors.

Preferential sputtering occurs when the composition of the flux of sputtered particles is different from their concentrations on the surface of multicomponent targets. It is due to the primary collision effects of ion bombardment.

The preferential removal of a component from the surface region of a target leads to the formation of an altered layer with a composition different from the bulk stoichiometry. The thickness of the altered layer is comparable to the range of the ions in the target. At temperatures where thermal diffusion is not important, the thickness of the altered layers is constant, and steady-state conditions are obtained. As already discussed in Sec- tion 3.1 of the first part' of this review, matter conservation leads to an equilibrium situation, where the composition of the flux of sputtered particles is equal to the bulk stoichiometry.

A useful parameter in the discussion of preferential sputtering effects is the component sputtering yield Y: of an element i in a multicomponent substrate. This is defined by the equation?29

where Yy is the partial sputtering yield and C: is the equilibrium surface concentration of element i. Using Equations 9 and 10 of the first part' of this review, and Equation 4, it follows that at steady-state sputtering of a binary compound AB, the ratio of the component sputtering yields is given by

ing energy differences of the constituent atomic species in the target. In a linear collision cascade, the momentum and energy transfers will be different for the different constituents, resulting in different ejection probabilities for the different atoms. This will, in general, cause preferential sputtering of the lighter component and a subsequent enrichment of the heavier element in the surface region of the target. However, this mass difference effect may be opposed by the difference in the surface-binding energies of the constituents. A component with a lower surface-binding energy will, in general, have a greater ejection probability than those with higher binding energies, resulting in an enrichment of the component with a higher surface-binding energy in the selvedge.

Several theoretical models have been proposed to quantitatively explain preferential sputtering (e.g., see References 239, 245, 246). A commonly used approximation is one based on Equation 5, where the ratio of the component sputter yields Y; is assumed to be equal to the ratio of the elemental sputter yields. The equilibrium sputtered surface composition ratio can then be expressed as

C' c, Y, (6)

A=-.- '; ',

where Yi is the sputtering yield of the pure element i.

In his preferential sputtering niodel. HaffZ4' assumed equipartition of beam energy and radiation-induced diffusion. He derived the following formula

( 7 )

A more advanced and popular theory is the Sigmund preferential sputter theory in the linear cascade regime.'" From this theory, it follows that the ratio of the component sputtering yields can be approximated by

y' ('A/'B) A = ( 5 ) Y; (Ck/C\B)

where C, and C, are the bulk concentrations.

preferential sputtering effects are the composition of the target, mass differences, and surface-bind-

In general, the main factors that contribute to I-?m

(8)

152

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

where Mi and Ui denote the atomic mass and surface-binding energy, respectively. Substitution of Equation 8 into Equation 5 yields

The exact value of the parameter m is uncer- tain because m depends on the energy. The parameter m in Equations 8 and 9 originates from the power cross-section that Sigmund used in his theory. For this cross-section, m = 1 at high energies (Rutherford scattering), m = 0.5 for medium energies (10 to 100 keV), and 0 I m I 0.5 for lower energies. For most applications in sputtering, the relevant values for m are in the range 0 I m I 0.2248.Z4y In the case of preferential sputtering of oxides, Malherbe et al.246 found that for 2m = 0.33, good agreement exists between a large amount of experimental data and theory. There- fore, to minimize the number of parameters in this section, we use this value for 2 m in the following applications of Equations 8 and 9.

Another difficulty in the application of Equa- tions 8 and 9 is the value of the surface-binding energy term Ui. In the case of metallic alloys, the heats of formation of gaseous atoms from elements in their standard states are usually taken to represent this energy.

Several other choices for the values of the surface-binding energies exist. This problem has been discussed in theoretical treatments of oxide ~ p u t t e r i n g . ~ ~ . ’ ~ ~ . ’ ~ ~ ~ ~ ~ ~ Malherbe et al. calculated the binding energies of the oxygen and metallic atoms by using a model based on a modification of the Pauling formalism25’ for the bond energies in a covalent bond. If their formalism is adapted for compound semiconductor AB, one obtains for component B

2 U, =+D(A-B)++D(B-B)-+(E, - E ~ ) (10)

and for the metallic component A

and H, the sublimation energy of the pure metallic element A.

The pair-binding mode1252+253 has also been used to calculate the surface-binding energies in a theoretical treatment of alloy preferential sput- t e ~ i n g . ~ ~ ~ Using the above notation, this model states that the ratio of the surface-binding energies is given by

u B= U A

C,D(B - B) + C,D(A - B) (12) C,D( A - B) + C A D( A - A)

Ke11y244.249.250 proposed a model in which linear cascade sputtering is assumed up to, but excluding, the top monolayer. For the surface-binding energy terms, Kelly used a model based on nearest neighbor bond strengths and arrived at the following expression

Ca = I AH: -C,h, -[C,(AH”,’ +C,(AH:)I

-2C,C,(AH: +AHL)h, +C,C,h~]US)//

+(‘B (13)

and mutatis mittandis for CBSr where DHia is the heat of atomization. According to Kelly,?* the heat of mixing term h, can be set equal to zero without really affecting the final answer.

For the sputtering of compounds under spike conditions, Sigmund24* derived the following equation

where T denotes the effective spike temperature. The problems associated with this temperature were discussed in Sections 2 and 3.7 of the first p a d of this review. The surface-binding energy term is given by243

=*h(aD, (15)

where a,, denotes the Debye frequency.

4.3. Kinetic Processes

where D (A - B) denotes the bond energy of the diatomic molecule AB, and &+the electronegativity

Several secondary ion collision processes can also contribute to surface compositional changes

153

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

of multicomponent targets. The most important of these processes are briefly discussed in Sec- tions 4.3.1 to 4.3.7. Nuclear transmutation of the substrate atoms is not discussed.

In practice, two or more of these processes take place simultaneously, but usually it is impossible to determine the relative contribution of each individual process to the resultant surface composition.

The combined effect of the different processes on compositional changes can be simulated in kinetic model^.?^^-^^^ Although these models help one understand the mechanisms causing such compositional changes during sputtering, the many parameters required in some kinetic theories and the scarcity of supportive experimental data make the validity of some conclusions somewhat ques- tionable.

4.3.1. Gibbsian Segregation

In many multicomponent materials, the surface composition differs from the bulk composition at elevated temperatures. This compositional change, known as Gibbsian (or surface, or equilibrium) segregation, is due to a thermodynamic force that causes the system to minimize its surface free energy (for a review, see Refe ren~es”~ , ’~ Gibbsian segregation and preferential sputtering are intimately connected because both are strongly dependent on surface-binding energies. For metallic alloys, the two processes often oppose each other, that is, surface enrichment of species A due to Gibbsian segregation in an alloy AB is opposed by a tendency for depletion of the same component A under ion sputtering.239

4.3.2. Radiation-Enhanced Segregation

Bombardment-induced defects can cause radiation-enhanced segregation in the target. For instance, diffusion of defects toward the surface (which acts as a sink) can result in an increase in concentration of a component in the surface. In contrast to Gibbsian segregation, this process is a nonequilibrium one. The enriched layer may dis- solve into the bulk after ion bombardment has ceased, provided the diffusion coefficient is high enough.257

4.3.3. Segregation

If the exchange of atoms between the surface and inner layers is sufficiently fast, the surface composition of a binary compound AB under thermal equilibrium is given by2-

where AG, is the segregation free energy. This is the energy change associated with exchange of an atom B in the surface layer with an adjacent atom A in the bulk.

4.3.4. Radiation-Enhanced Diffusion

Ion bombardment causes the concentration of point defects in the target to far exceed the concentration at thermodynamic equilibrium. Because the diffusion coefficient is proportional to the point defect concentration, these radiation-induced defects can lead to an enhancement of the diffusion coefficient by several orders of magnitude compared to the equilibrium one. The ion bombardment thus effectively lowers the temperature at which diffusion processes dominate compositional changes. This radiation-enhanced diffusion can then lead to surface compositional changes.

4.3.5. Displacement Mixing

The bombardment-induced cascade event leads to a spatial relocation of atoms and pro- duces point defects in the substrate. This atom relocation process is called displacement mixing (or atomic mixing) and comprises recoil implantation and cascade mixing. Recoil implantation takes place when one type of atom is preferentially transported in the beam direction due to preferential momentum transfer. Cascade mixing is a random-walk process resulting from the movement of secondary recoil collisions. Thus, the displacement process can also, to a first approximation, be described by a diffusion-like model.

4.3.6. lmplantant Surface Concentration

During ion bombardment, the ions become trapped in the substrate and subsequently change

154

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

the surface composition. The surface concentration of the implantant depends on several parameters, such as the sputter depth, ion range distribution, chemical reactivity of the ion/substrate combination, etc. Assuming a Gaussian range distribution for the bombarding ions, Schulz and W i t t m a a ~ k ~ ~ ~ derived the following equation for the equilibrium implantant concentration distribution c(x) in the presence of sputtering:

2Y

where x is the depth inside the sputtered surface, Y is the sputter yield, and R, and AR, are the ion projected range and range straggling, respectively. The error function is given by

Equation 17 has been employed to calculate ion range distribution parameters.'*26R The saturation surface concentration that occurs at equilibrium sputtering easily follows from Equation 17:

2Y

Equation 19 has been used in comparisons with measured surface concentrations of noble gas implantants (e.g., References 269-278). For high-dose density implantations, the surface concentrations of the implantants were much smaller than the predicted. In the case of low dose density implantation of noble gas ions, the retention of the implantant closely agreed with model prediction^.^'^ The retention characteristics of implantants at high-dose densities also strongly depends on the chemical reactivity between the implantation species and the target material.273 In an AES investigation of low-energy (0.5 to 5 keV) nitrogen bombardment of silicon, a Si,N, layer was formed,267 while similar implantations of CuZa resulted in much lower surface concentrations of nitrogen.

Singer et a1.280 used AES to determine the dependence of the surface concentration of implanted argon in GaAs on the argon energy for

0.5- to 5-keV A P sputtering. The results by Singer et al. exhibit some qualitative agreement with Equation 19 because the surface concentration of argon decreased with increasing ion energy. The concentrations varied from approximately 1.8 to 0.4 at%. According to Equation 19, these surface concentrations give unrealistically high sputter yield values (see Section 3 of the first part' of this review for typical sputter yield values in this energy range). For 10- to 40-keV krypton sputtering of GaAs, Carter et al. also found that application of the Schulz-Wittmaack results in too high sputtering yield values.

From the above discussion, it follows that the surface concentrations of noble gas implantation species during sputtering of compound semiconductors are small but not insignificant. Further- more, there are often large mass and size differences between the implanted atomic species and the substrate atoms. This can cause the energy transfer process to be significantly different in an implanted substrate, with consequences for the surface composition.

4.4. Sputtering of InP

4.4.1. Experimental Measurements

When InP is sputtered by noble gas ions, the composition of the surface layer is dramatically altered. Typical AES spectra (in differentiated form) obtained from three InP surfaces - as- received, vacuum cleaved, and Ar' sputtered - are shown in Figure 14. The Auger spectrum for the vacuum-cleaved InP was obtained from a (1 10) cleaved surface, while the other two were obtained from (100) faces. The spectrum from the as-received sample (Figure 14a) shows C and a large 0 Auger peaks due to hydrocarbon contamination and oxide formation on the InP. Dis- tinct differences are seen between the P peak, and to a lesser extent the In peaks, for the as-received and vacuum-cleaved surfaces. The drastic change in the P (120 eV) peak is due to the LVV transition being more sensitive to the chemical environment than the MNN transition of In.281 The peak changes have been discussed in more detail e l s e ~ h e r e . ~ ~ ~ . * ~ ~ The composition of the sputtered surface was calculated to be 40 _+ 3 at% P and 60 f 3 at% In. Thus, Ar+ bombardment leads to an

155

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

I I I I I I

0 100 200 300 400 500 600 KINETIC ENERGY (eV)

FIGURE 14. Differentiated AES spectra of (a) as-received InP (100) surface, (b) vacuum- cleaved (1 10) surface, and (c) a (100) surface cleaned by 0.5-keV Ar+ sputtering at a 40' angle of incidence. Note the changes in the In/P peak-to-peak ratio between the spectrum in b and c, indicating that compositional changes were introduced by the argon ion bornbard- ment. (From Barnard, W. O., Malherbe. J. B., and Myburg, G., S. Afr. J. Phys., 14,22, 1991. With permission.)

156

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

indium-enriched surface due to preferential sputtering of P.

In enrichment on the InP surface has also been found for sputtering by other noble gas ions and by reactive ions such as Cl,2”.2x5 H,285-287 0 , 3 1 s 6 2 etc. Laser irradiation also leads to a preferential loss of P from InP.2x8

Published findings on surface compositional change measurements due to sputtering of InP by low-energy (EP I 10 keV) noble gas ions are summarized in near chronological order in Table 1. In this compilation, the composition of the sputtered surface is given by the following notation: InP, with x = C;/C;, where Cf is the surface concentration of species i after steady-state sputtering.

4.4.2. Factors Affecting the Accuracy of the Measurements

A striking feature of the data presented in Table 1 is the large scatter in the data. This indicates a lack of consensus as to the correct value for the sputtered surface composition. These discrepancies may be due to a number of reasons, which are discussed below.

An important factor that may influence the accuracy of the data in Table 1 is the appearance of cone-like structures on the sputtered surfaces. Ion bombardment-induced cone development was discussed in detail in Section 2. In summary, while the cones originate via a growth process and consist in the main of monocrystalline InP, there is also an excess of metallic In on the surface. However, the mere existence of cone-like structures may distort AES and XPS measurement signals from such surfaces. This is due to the effect of surface roughness on such analyses. The excitation/analyzing beams in most studies had diameters that were much larger than the individual cones (30 to 90 nm) as well as the average distance between them (50 to 200 nm). Several of these cones are therefore analyzed in a single investigation.

For obvious reasons, the presence of reactive gases in the vacuum system may also detrimen- tally influence the accuracy of the measurements.

A related factor affecting the accuracy of surface compositional concentrations is the incorporation of atoms of the bombarding species. As

discussed in Section 4.3.7, the argon surface concentration in the sputtered substrate is low but significant. For the low sputter energies considered in Table 1, the argon concentration should decrease with increasing energy.

The different analyzing depths and principles of the various techniques used to obtain the data in Table 1 are certain to contribute to the variations in the data. For instance, most AES and XPS studies indicate preferential sputtering of P from the InP. Using energy loss spectroscopy (ELS), Tu et al.*% found the first two to three monolayers to consist only of In after sputtering, while their AES spectra also showed the presence of phosphorus. In fact, by using the AES spectrum for a vacuum-cleaved (1 10) InP surface as a reference sample, the sputtered surface can be calculated to have a composition of approximately 1:l for In and P (see Reference 312 for details of the calculation method). This suggests that no preferential sputtering has occurred. A partial explanation for these contradictory results (also with most other AES measurements; see for example, References 3 1,47,57,59,6 1,64,98,28 1,282,29 1,297-302, 307, 308, and 3 10-3 14) may be the presence of oxygen (not taken into account in the above-mentioned calculation) in the AES spectrum of the sputtered InP surface.?96 The ISS measurements by Barcz et a1.293.294 also showed no preferential sputtering effect for InP, in contrast to most AES and XPS results. ISS is more surface specific than either AES or XPS in that it analyzes only the first monolayer, while AES and XPS have analyzing depths extending over many monolayers. This kind of discrepancy has long been known to exist in preferential sputtering studies on other binary systems. Nelson3I5 found no preferential sputtering on AgAu using ISS, while Holloway and Bhattacharya3I6 did, using AES. The ISS/(AES, XPS) discrepancies are usually explained in terms of radiation-enhanced diffusion and surface segregation and their effects on different sampling depths, ion bombarding energies, and ion current densities.

Diffusion and segregation effects become more pronounced with increasing substrate temperature. The sputter energies and ion dose densities used in most sputtering studies are too low to cause a significant increase in the sample temperature. However, simultaneous bombardment

157

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

TABLE 1 Summary of Published Results of Surface Compositional Changes in InP Due to Low Energy (EP I 10 keV) Noble Gas Ion Sputtering

Ion energy (keW

0.48

0.5 - 0.5 0.75 0.5 0.5 4 0.5, 10 25, 40 1 .2 .3 2

0.5

1 1.5 3 0.5 0.5- 5

0.5

0.25- 1.5

8

c0.5

co.5

2

1-3

Angle Of Sputtered surface

Incidence (7 cornposition Analyzing

technique(@

XPS. AES. LEED

SlMS AES

AES AES , AES

XPS ISS.

FIBS ELS,

SlMS

AES

AES. RHEED

Ion sF=ie(s)

Ar

Ar Ar

Ar Ar Ar H Ar

Nem

Ar Ar

Ar

Ar

Ar

Ar

Ar

Ar

Ar

Ar

Ar

Remarks Ref.

InP 289 -

45 - - - - - - -

45 -

0

0 0 0 45 -

-

-

35.7 41 50 61 68 77 81

89 as

-

- -

45

Textured surface Preferential sputtering effect observed -

51 290

In enriched - InPo64

InPo, In enriched

InP05 In enriched InP

281 29 1 56 In islands

292 293, 294

In enriched In

295 296

- ELS: 2-3 ml In AES: still some P present;

Prolonged 3-keV in microclusters on surface

Ar- sputtering induced SEM contrast (r - 0.5 pm), but SAM gave no evidence of In microdroplets

59

AES 57 Textured surface Sample cooled to -183%; no surface topography

From their Figure 3. using elemental sensitivity method, (similar to our method)

sputtering Cones after prolonged

In islands formed

297 AES

AES 47, 64

60 ELS XPS

In ctystallites with a height of four monolayers and coverage e = 0.25 -

298-301 AES. EPES. ELS

AES 302

303

304, 305

In-rich ML on P ML

In enriched ELS: formation of In islands

No difference between (1 00) and (1 10) surfaces and ion energy. Top two monolayen: InPo,. half of which is In metal, 20% bulk-like InP. and remainder In bonded to P with average composition InP,. Deeper layers: =InP but In metal and InP, still present

metallic type; textured surface

created

P decreases with increasing angle of incidence in the region 35" 5 e 5 90"; In islands

26% of total In atoms are of a

ELS: In islands

preferential sputtering of

ELS, XPS

XPS

61

98

306

ELS, AES

ELS. AES

XPS. ELS

Ar 0.5 0

Ar 1, 3, 6 - - - Ar

158

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

TABLE 1 (continued) Summary of Published Results of Surface Compositional Changes in InP Due to Low Energy (EP I 10 keV) Noble Gas Ion Sputtering

Analyzing ion technique(.¶) specie(s)

AES. Ar

SEM Ar ELS

XPS Ar

ARAES Ar ee = 120

ARAES Ar

XPS. Ar

AES Ar

e. = 72"

ELS

AES. Ar SAM. XPS, SEM

AES

AES

AES

Ion energy (keV)

0.3

0.2. 0.5,

0.5

0.5 1 2 4 5 4

4

0.5-5

0.8

3

Angle of

incidence (9

-

0-75

-

55

55

55

18 40 42 70 60

Ar 0.5-5 0 1 1 30 41 52 60 71

Ar 0.25 80 0.5 80 0.75 80 1 .o 80 2.0 80

11 21 41 51 61 71 75 77

Kr 0.5-5 1

with an intense electron beam (as is often used in AES) can cause a significant increase in the sample ternperat~re.~~' Also, InP decomposes at elevated temperatures, leaving the substrate In

Due to heating (causing decomposition), bond breaking, charging, and contamination ef- f e c t ~ , ~ ' ~ ~ ~ ~ ~ ~ ~ ~ an intense electron beam (used as an excitation source in some techniques such as

enriched.77-7?,318

Sputtered surface composition

In enriched

Remarks Ref.

Some metallic In: sub- 307

0.8 keV: In droplets on surface 308 0.2. 0.5 keV: no In droplets Broken In-P bonds 309

microscopic islands

Surface modifications due to P-preferential sputtering: XPS: InP, In-In bonds, and InP2 bonds; ELS: In metal clusters - size, -tens of A

310, 311

Textured surface after 282. sputtering: 312 cornposition not dependent on prior treatment

substrate temperature

enrichment in In

50 nm)

decreases if substrate temperature decreases

ing

ion energy and not dependent on prior treatment

Cone formation, depending on 31

AES: cones: InP with slight

SAM: no In island (resolution

XPS: In metal peaks: intensity

Sputter cones after sputter- 313

Composition independent of

314

Sputter cones after sputter- 283 ing; composition independent of ion energy

AES) can cause some changes in the surface composition of compound semiconductor^.^^^ These changes are usually small at the normal electron beam densities used in AES measurements and should not affect measurements of radiation-induced compositional changes in InP.

A key problem associated with the measurement of surface compositional changes is that of obtaining a standard surface composition that can

159

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

be used as a reference. For most monocrystalline compound semiconductors, such a reference can be obtained by cleaving a sample in ultrahigh vacuum. For most surface-sensitive analytic methods (such as AES, XPS, ELS, etc.) employed in surface compositional change measurements, the analysis depths extend over a few monolayers. The assumption during measurements performed on a cleaved sample is that the signal is from the bulk composition. When using a technique with an electron beam as excitation source, one must beware of channeling effects. As shown in Figure 15, these can be very Despite the strong intensity variations of the individual Auger signals for InP in Figure 15. the ratio of the In and

P signals remains nearly constant (probably within experimental error) over a wide range of angles of incidence of the primary electron beam. This is also true for GaAs, although the signal ratio of Ga to As is not quite as constant as the corresponding ratio of InP. More details regarding this cleaving method can be found in References 282.3 12,3 13, 324, and 327.

Incorrect specification of the ion beam parameters, such as the ion bombardment energy, might also have an effect on the sputtered surface composition. This effect is discussed in more detail in Section 4.4.3.

A beam parameter that definitely has a significant effect on the sputter surface composition

n

5 4

r

U

(1 Q

E W c3 I) U

0 a I- (r

Q. Q I

1 2 \ 0 0

0 I I l l I l l I l l I l l I

-40 -20 0 20 40

InP ( I 10) b)

n

5 4 Y

Q Q I LT W c3 I) 4

0 a I-

LT

Q Q. I a I

> -

1

INCIDENCE ANGLE [4]

FIGURE 15. Polar scans in the plane that contains the [i 101 + [OI 01 -+ [I lo]+ [IOO] + [I i 01 directions done with a CMA of (a) cleaved GaAs (1 10) and (b) InP (1 10). The intensity variations in the Auger peaks are mainly due to the channeling of the incident electron beam, producing variations of the ionization cross-section when the incident direction is close to a Bragg direction. Another factor that contributes to the intensity variations is the angular anisotropy. Diffraction effects can be ignored due to the poor angular resolution of a CMA. In the lower part of the figures are the peak-to-peak height ratios of the above Auger lines, showing very little variation with respect to the angle of incidence. (From di Bona, A., Facchini, A., Valerie, S., Ottaviani, G., and Piccirillo, A., Mater. Res. SOC. Symp. Proc.. 223, 197, 1991. With permission.)

160

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

is the angle of incidence of the ion beam. Unfor- tunately, this angle is seldom specified in the papers summarized in Table 1, making comparison between the measurements difficult. This angle dependence of the sputtered surface composition is discussed in more detail in Section 4.4.3.

4.4.3. Discussion of the Experimental Data

This section presents some general conclusions regarding the different measurements of noble gas bombardment-induced compositional changes in InP as summarized in Table 1.

From the compilation in Table 1, it can be seen that prior treatment of the surfaces had no influence on the equilibrium sputtered surface

0.80

0.75 0

0 [1L

.- -+

2 0.70

0.65 0

composition of I n P . 2 8 2 3 3 ' 2 . 3 ' 3 Furthermore, the equilibrium sputtered surface composition of InP was found to be independent of the crystal orientation.282.283.304.3'"313 The (loo), (1 lo), vicinal InP, and differently pretreated ( 100) surfaces were investigated. From Figure 16, it can be concluded that no significant (more than the mean deviation from analysis to analysis) difference exists between the I n P peak ratio for two different crystal orientations. InP is a semiconductor that is readily amorphized by Ar' bombardment29s (see also the discussion in Section 3.2.2). The low-energy ions used in most sputtering studies have very small projected ranges, thus aiding in the destruction of the single-crystal structure in the near-surface region. Any ion-channeling effects to influence the surface composition are accordingly mini- mized.

I I I 1 'I

A r ' - InP

T

I I I I I I

1 2 3 4 5

Ion Energy E (keV) 6

FIGURE 16. Steady-state sputtered surface composition ratio (i.e., InP,) as a function of ion energy for two different crystal orientations. The error bars represent only the statistical variation between ten measurements. (From Malherbe, J. B., Appl. Surf. Sci., 70171, 322, 1993. With permission.)

161

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

There is disagreement among the few results given in Table 1 regarding the dependence of the final surface composition on the ion energy. Peacock59 and Valeri and Lolli3'O reported an energy dependence, while other au-

dence (see also Figure 16). Models of preferential s p ~ t t e r i n g 2 ~ ~ ~ * also make no provision for energy dependence either. Only near the threshold energy for sputtering do simple energy transfer considerations show that preferential sputtering effects should be ion-energy dependent.32E In the sputtering experiments summarized in Table 1, the ion energies were well above the threshold sputtering energy of -25 eV for 1nP.I

In Figure 17, a depth profile is shown of the In and P atomic fraction ratio after first sputtering with 4-keV Ar' and then switching to 0.5-keV Ar+ sputtering.310 The I n P ratio as a function of sputter depth is completely different from the corresponding Ga/As ratio for GaAs. There is no indication of a segregated layer at the sputtered surface or of a depleted layer just below the surface. Using angle-resolved AES and X-ray photoelectron spectroscopy measurements, Valeri and Lolli310 found that the layer in which the composition was different from that of the bulk was much smaller than the projected range of the argon ions in InP and interpreted these results to indicate that preferential sputtering is the mechanism that causes the bombardment-induced compositional changes in InP.

Several s t ~ d i e s ~ ~ . ~ . ~ ~ ' . ~ ~ . ~ ~ ~ . ~ ~ ~ have shown that the equilibrium surface composition depends on the angle of incidence of the bombarding ions. However, the angle of incidence is seldom specified in the papers summarized in Table 1, making comparisons between the reported data difficult. In Figure 18, all those measurements are shown in which the angle of incidence was specified. A wide scatter exists between the different results. Reasons for this scatter have been discussed in Section 4.4.2. Note that the development of surface topography on InP during low-energy ion bombardment is also ion-angle dependent (Sec- tion 2.3).

In general, it was found that with grazing incidence, the preferential sputtering of P from InP becomes less. This observation is in line with results on some other system^,^^^-^^* such as Ni-Fe,333 C~Inse , , "~ Ta20,,32E and CUT^.^^'

~ors98.282.283.29~295.3O4.3 12.3 I 3 found no such depen-

In line with the scarcity of experimental data on the angle dependence of equilibrium-sputtered surface compositions of multicomponent materials, there are also very few models to explain these results.

Using other systems, Taglauer and co-work- e r ~ ~ * ~ - ~ ~ ~ found that an oblique ion beam incidence leads to an enhanced forward emission of the heavier components of these systems. This is explained by the angle dependence of the sputtering due to an incoming projectile generating a primary recoil atom that leaves the surface.330 A Monte Carlo computer simulation showed that there is a relatively low probability of recoils of the heavier atomic species under- going large angle scattering events with low- energy transfer, in contrast to the primary recoils of the lighter atomic species. This leads to a more stoichiometric surface composition at glancing angle sputtering for the following reason: the two main factors that influence pure preferential sputtering are the differences between the masses and between the surface-binding energies of the constituent surface atoms in a multicomponent material. In systems where the mass difference is large (as is the case with InP), preferential effects are usually pronounced, leading to an accumulation of the heavier species in the substrate. Therefore, ion bombardment at a glancing angle of incidence will di- minis h t hi s accumulation.

Another explanation for this angle dependence may be the creation of spikes. The amount of energy Q per unit length needed to vaporize and decompose the InP into individual gaseous atoms is given by Q(eV/A) = HJa,, where H, is the vaporization and decomposition enthalpy, and a, the lattice parameter. Using the values given in Moses,335 one obtains Q = 1.18 eV/A for InP. According to the PRAL computer code,336 which is also contained in the TRIM code,337 this is much less than the average energy loss due to nuclear collisions of an Ar' ion along its track in the InP, viz., 34 eV/A for 3-keV Ar' ions, indicating the existence of spike regions in the InP. There is a greater probability at glancing angle incidence for such a spike region to end on the surface of the target. In such a region, all the atoms have such large energies that mass and surface energy difference effects are negligible. All these atoms with an outward velocity compo-

162

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

- 7

rd - a a

a2

a Irl u 3 a

H H H I 0 & I > 0

E

cr: z 0

H u a [r,

U

E 0 H

H

a

H

a

H

a

10

5

1 .o

0.8

0.6

0.4

I I I I I I 1) I I GaAs (1 10)

A s ( 3 1 eV) 1.. . .* \ ...... -- .-._._. -.-..

Ga(55 e V ) \

I*+--" -0-%-+

AA 1 k e V --+0.5 k e V

4 keV --+ 0.5 k e V GaAs { InP 0 4 kev --+0.5 keV

FIGURE 17. (a) Peak-to-peak heights of the main Ga and As Auger peaks in GaAs as a function of the sputter time for 4-keV and 0.5-keV argon ion sputtering. The GaAs sample was first sputtered by 4-keV Ar+ ions and then by 0.5-keV Ar+ ions to show the transition between the high- and low-ion energy regimes. (b) Sur- face atomic fraction ratios In/P and GdAs as a function of sputtering time. These ratios were calculated from the Auger peak-to-peak heights to show the change in surface concentration when switching from a higher argon ion energy (4 keV as well as 1 keV) to a lower energy (0.5 keV). In the case of GaAs, the transition between the steady-state regime of 4-keV sputtering and the transient regime of 0.5-keV sputtering is shown for both the high-energy (Ga LMM 1070 eV and As LMM 1228 eV) Auger peaks, indicated by the filled circles and triangles, and the low-energy (Ga MMM 55 eV and As MNN 31 eV) Auger peaks, indicated by the open triangles. (From Valerie, S. and Lolli, M., Surf. interface Anal., 16, 59, 1990. With permission.)

163

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

2

X

.- O 1 -w U [li 0.9

\ 0.8 -

0.7

0.6

0.5

0.4

I I

+ Ar - InP

0 This study

Bornard et a l A Malherbe e t a l 0 Peacock

0 Valer i e t a l A

v Jardin e t a l

Hou e t a l

0 Lau et a l 0

0

0 A

V 0

I I

0.01 0.1 1

c o s oi FIGURE 18. Equilibrium-sputtered surface composition ratio x (i.e., InP,) as a function of the angle of incidence Bi of the bombarding Ar* ions. The other published results were all collected from References 59 to 61,306, 312, and 313. (From Malherbe, J. B., Appl. Surf. Sci., 70/71, 322, 1993. With permission.)

nent will be emitted from the target. This “hot spot” contribution will reduce the nonstoichio- metric part of the sputter yields from the normal sputtering process.

From Figure 18, it can be seen that the data of Hou et aLg8 and Malherbe and co-workers283.312.313 can be fitted to straight lines. From Equation 5, it follows that the ratio of the equilibrium sputtered surface fractional compositions q of InP, divided by the ratio of the bulk fractional compositions Cp, can be related to the component sputter yields

Y$239 The approximation given in Equation 6 re- lates this ratio to the ratio of the elemental sputtering yields Yi. Using these relationships and the data in Figure 18, it follows that

where f is the slope of the above-mentioned straight line fitting, 8 is the angle of incidence of the ions, and Y represents either the elemental or compo-

164

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

nent sputter yields. The values for f are 0.11 for the data of Hou et al.98 and 0.17 for the data of Malherbe313 (note that incorrect values were specified in Reference 3 13). This cosine dependence is similar to the predictions of the Sigmund sputter model for angle dependence of the total sputtering yield, as discussed in Section 3.3.1 of the first part! of this review, viz., Y = (cos Cl)-f with f = 1.7 k 0.7. If one assumes that each individual sputtering yield Y in Equation 20 also has the same dependence on the ion angle of incidence, as predicted by the Sigmund equation for the total sputtering yield, then Equation 20 becomes

In terms of this model, the above experimental ion angle dependence results for InP are explained by fin # fp.

Most results of the many experiments on compositional changes in InP due to noble gas sputtering have indicated that preferential sputtering is the mechanism that causes the changes observed. Segregation- and diffusion-related experiments, however, are difficult to perform reliably on InP due to the fact that InP decomposes at elevated temperature^.^^-'^ Only a few experi- m e n t ~ ~ ' . ' ~ . ~ ~ have been performed at temperatures below room temperature. The results from the two groups seem to be in some contrast to each other. Skinner et a1.56.57 found an enhanced preferential sputtering of phosphorus, while MacLaren et al.31 reported that the In Auger peak shape suggested less preferential sputtering of phosphorus. The one angle-resolved AES study310 did not produce any real evidence of strong segregation effects. Valeri and Lolli3l0 found that the In/P ratio increased as the sampling depth was decreased. This is consistent with a preferential sputtering mechanism. The only contrary evidence is the ISS experiment by Barcz et al.,293.294 which was discussed above.

From a preferential sputtering mechanism, one would expect the enriched atomic species (in this case In) to be evenly dispersed throughout each atomic layer in the altered region. Two types of experiments, however, suggest that some kind of diffusion mechanism may lead to agglomeration of In clusters in the surface of the

InP. The first of these is based on chemical infor-

ELS~~61~300.303,306.307.310 measurements. These measurements show evidence of pure In metal peaks together with other peaks. The size of these pure In clusters should be on the order of tens of Bingstroms or larger.310 The other evidence comes from the XPS measurements by MacLaren et al.31 who found a decrease in the In(meta1)-XPS signal when the substrate was cooled to below room temperature during sputtering.

To summarize the ion bombardment-induced compositional changes in InP, experimental studies show that there is an enrichment of In in the surface due to preferential sputtering of P. The equilibrium surface composition is independent of the ion energy in the low-energy region (0.5 to 10 keV) and of the sample orientation. The angle of incidence of the bombarding ions, however, has an effect on the final surface composition.

mation obtained by X P S ~ I - ~ . ~ ~ ~ . ~ O ~ X W ~ O J I I and

4.4.4. Predictions of Theoretical Models

In this section, we fit the few theoretical preferential sputtering models, mentioned in Section 4.1, to the experimental data for the case of low- energy argon ion bombardment of InP. If all the quantitative argon sputtered surface compositions given in Table 1 (in total 50) are arithmetically averaged33x (thus excluding energy or angle- dependence effects), one obtains a value of InP,,, * o,19. The majority of the references in Table 1 that do not provide quantitative information of the sputtered surface composition also indicate the preferential sputtering of phosphorus from InP.

The elementary preferential sputtering theory of Haff247 predicts an equilibrium sputtered surface composition of InPo.72, which is near the mean experimental value of InP. This reasonably good agreement between theory and experimental values is somewhat surprising when taking the simplicity of the Haff theory into account.

As mentioned previously (Section 4. l), a problem in the application of the Sigmund preferential sputter theory is the value of the surface-binding energy term Ui. The heats of formation of gaseous atoms from elements in their standard states AH; were tabulated for In and P by AH: (In) =

165

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

2.52 eV and AH;(P) = 3.44 eV. Using these values in Equation 8, one obtains CYCb = 0.80, which predicts an indium-enriched sputtered surface in agreement with most experimental results, albeit this value is higher than most experimental measurements.

When using the Sigmund preferential sputtering theory (Equation 9) and Malherbe et al. formalism (Equations 10 and 11) for the surface- binding energies of In and P, exactly the same value for the equilibrium sputtered surface composition of InP is obtained as that predicted by the preferential sputtering theory of Haff.247 Substi- tuting the values D(In - P) = 2.05 eV, D(P - P) = 5.07 eV (see Reference 339), E,, = 1.78, and E, = 2.19 (see Reference 340) into Equations 9 to 1 I , one obtains CgC;,,= 0.72. This value is in good agreement with most experimental data.

The Sigmund equation with the pair-binding model (see Equations 9 and 12) also gives a value reasonably near the mean experimental value for the ratio of the surface concentrations after prolonged sputtering. Again, using the bond ener- g i e ~ ~ ' ~ of the diatomic molecules in Equation 12 and substituting in Equation 9, the value CYC:, = 0.85 is obtained.

By setting the heat of mixing term h, in Equa- tion 13 equal to zero and substituting the heats of formation of gaseous atoms from elements in the standard states for AH:, one obtains CYC;,, = 1.17 for the Kelly model. This value predicts phosphorus enrichment during sputtering, in complete disagreement with all experimental studies and the other theoretical predictions considered in this article.

The main difference between this theory of Kelly (i.e., Equations 12 and 13) and the others considered in this article (Sigmund248 and HafPJ7) is the fact that the atomic mass difference effect of the substrate is neglected in his theory. In the case of InP, this mass difference is large, while for GaAs it is negligible. Simple ballistic energy transfer considerations show that this difference is of major importance in the former case.328

The above calculations confirm that the effect of surface-binding energy is of secondary importance to the mass ratio factor. This conclusion is supported by the results of Malherbe and Ba~nard ,~ '~ using the Trim-89 computer code.337 This code uses the same averaged surface-bind-

ing energy for both In and P. The results, therefore, reflect only the mass difference effect on the sputter yields of the two atomic species. For 0.5-keV Ar+ bombardment of stoichiometric InP, the relative sputter yield of Y,,,/Y, = 0.85 indicates a preferential sputtering of the P species.

To compare the experimentally determined surface composition with the Trim-89 calculations, Malherbe and Barnard312 used a conservation principle: at steady state, the ratio of the sputter yields of the two atomic species must be equal to the ratio of the atomic bulk concentra-

Furthermore, they assumed that the sputtered particles originate only from the top layers and thus neglected the dependence of the concentration profile on depth. Based on these premises, the composition of the target was varied for each bombarding energy and angle of incidence until a sputter ratio of Y,,,/Y, = 1 was obtained.

This method yielded a sputtered surface composition of InP,,,, for normal-incident 0.5-keVargon ions on InP. In contrast to most experimental results and several other measurements,98,29*?95.3~.31~.3'~ the Trim-89 calculations indicated a small energy dependence effect, with a decrease in the preferential sputtering of P with increasing bombardment energy. For normal-incident Ar+ bombardment, the calculated sputtered surface composition changed from InP,,,, for 0.5-keV Ar+ to InP,,, for 5-keV Ar+ ions. The difference (4%) between these two calculated compositions is smaller than the experimental error (210%) in the measurements given in Table 1.

The calculations also show a small angle of incidence dependence, from InPo,,6 for 0.5-keV Ar+ at normal incidence to InP,., when the angle of incidence was increased to 70". Although this angle dependence is smaller than most experimental values, the tendency is the same: less preferential sputtering of P with increasing angle of incidence in the range of 0 to 70".

Because of the incorrect surface-binding energy values used, one cannot really expect quantitative agreement between the Trim-89 calculations and the experimental values. For a review of the problems of computer simulations of the sputtering process, the reader is referred to Reference 34 1.

Thus, there is a large variation in the predictions of the present linear cascade preferential sputtering theories for the sputtered surface com-

166

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

position of InP. This is not surprising because these theories are certainly too simplified to agree with the real complexity of the sputtering process. To improve the agreement with experimental data, highly nonequilibrium collision cascade dependence, angle dependence, energy dependence, and matrix dependence terms (such as surface-binding energy, which varies with the surface composition and structure) are needed in these theories.

To summarize, the predictions of the sputtering theories considered and the Trim-89 Monte Carlo code show that the mass difference between the In and P atoms is the major factor contributing to the preferential sputtering of InP. The effect of the surface-binding energies of the two atomic species is of secondary importance.

4.5. Sputtering of GaAs

The (100) and ( 1 10) GaAs surfaces are the two most widely used and studied compound semiconductor surfaces. Consequently, several pa-

preferential sputtering of these surfaces. Table 2 contains a synopsis of published results on the equilibrium sputtered surface composition of single-crystal GaAs subjected to noble gas ion bombardment. To facilitate comparisons, the following notation was used in Table 2 for the composition of the sputtered surface: GaAs, with x = Ci,/CE;,, where C; is the surface concentration of species i, after steady-state sputtering, as determined by the various techniques.

There is overwhelming evidence that sputtering leads to Ga enrichment in the surface region of GaAs. Typical AES spectra before and after argon ion bombardment of a vacuum-cleaved (1 10) GaAs surface are given in Figure 19. This figure clearly shows a reduction in the main high- energy peak of As compared with the high-energy Auger peak of Ga after sputtering. This indicates an overall Ga enrichment in several monolayers in the surface region of GaAs due to the relatively large sampling depth of these Auger electrons. The low-energy Auger peaks (with much smaller sampling depth), however, exhibited peak shape changes before and after sputtering. It is therefore inadvisable to draw any conclusion from these low-energy differentiated peaks about possible Ga enrichment in the top surface layers.

pers85.1 32,293.294.299.3 10.342-368 have appeared on the

It is interesting to compare the predictions of the Sigmund preferential Sputtering theory to the experimental values given in Table 2. In contrast to InP, the mass difference between Ga and As is negligible. Thus, according to Equations 8 and 9, the difference in the surface-binding energies of these two types of atoms will play the dominant role in the preferential sputtering of GaAs by Ar‘. Indeed, a comparison of the sublimation ener- gies335 shows that the sublimation energy of elemental As (i.e., H, (As) = 29 kcal mol-l) is smaller than that of elemental Ga (Le., H,(Ga) = 64.9 kcal mol-I), explaining, in agreement with most measurements, the preferential sputtering of As from GaAs. However, substituting these Val- ues into Equation 9 and taking 2m = 0.33 gives a sputtered surface composition of GaAs,,. This As concentration value is considerably smaller than the experimental measurements summarized in Table 2. It is therefore tempting to conclude that this Ga enrichment does, like InP, originate solely from preferential sputtering effects.

Only a few of the measurements in Table 2 do not support the conclusion of Ga enrichment. Two of these are the results by M ~ G u i r e ~ ~ ’ and Will- i a m ~ . ~ ~ ~ In the latter case, the surface composition was determined using RBS, which, according to the author, was not really sensitive enough for accurate values. The results obtained by Bhattacharya et a1.82 are not really applicable to the present discussion. These results were obtained on samples sputtered at elevated temperatures. From the discussion in Section 3.4.1 of the first part1 of this review and the studies by Singer et a1.280.369.370 it follows that enhanced diffusion and recrystallization take place during sputtering.

In Figure 20, deconvoluted angle-resolved XPS results by Bussing et al.348 are given for argon-bombarded GaAs. These results indicate that the top surface layer consists of stoichiometric GaAs, but that there is a region with an As deficiency below the surface. The thickness of this subsurface As-depleted region increases with increasing Ar‘ bombardment energy. A depth profile as in Figure 20 is typical of a surface segregation mechanism induced by the ion bombardment of an ion-bombarded compound or al-

Several other studies also support the segregation model of Bussing et a1.M8 Studies by Singer

10y.245

167

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

TABLE 2 Summary of Published Results of Surface Compositional Changes in GaAs Due to Noble Gas Sputtering

G a h 5urface

(loo). (111) n-type 1 x 10” cm’

Analyslr tKh- nique

AES

AES. UPS

AES

AES

GDOS

AES

AES

XPS

ISS

AES

Ion 5pOClm

Ar

Ar

Ar

Ar

Ar

Ar

Kr Xe Ar

Ar

Ne

Ar

Ion e * w IE(kW1

1.5

0.9

0.6

1

2

0.5 0.5 2 2 5

2

10 15

40

0.05

0 . 5 5

lon current denrlty

U W C 4

-

-

20 x 1oJ

-

-

-

- -

- 2.5

10,000

0.3-100

Ar 2C-100 -

RBS cham Ar 1-3 - neling

AES Ar 3 3

lon angle Sputtered of Incidence wrfnce

composition

Ga

Slight Ga

G a o n

G a s , 05

enrichment

enrichment

Ga enrichment

G a s , 3o G a s , i d

GaAs GaAs, o(i Ga

Ga enrichment

enrichment GaASo m GAS0 m

GaAs, m

-

Ga enrichment to stoichiometry

-

GaAs

Ga enrichment

Ga enrichment increased with increasing A r energy between 0-2 keV

Steady-state sputtering after removal of less than 100 A

As deficiency decreased monotonically from the surface to the bulk

Much higher As-deficiency with low energy (50 eV) but high current density (1 mA/cmz)

Ga enrichment due to theno- dynamic processes rather than preferential sputtering

Above a critical temperature T,. no surface compositional changes occurred; below T,, there was Ga enrichment

T, depended on ion energy and ion flux

At a fixed energy, T, increased with increasing ion flux; at a

Ref.

351

350

343

354

85

347

352

355

293, 294

280. 369. 370

fixed flux, with increasing energy Above temperature results are explained in terms of annealing of radiation-induced damage

Ga enrichment was due to preferential sputtering of As from a disordered surface

Below T,, the sputtered surface composition depended on the ion energy and on the ion flux: Ga enrichment increased with increasing ion energy and ion current density

for accurate surface composition determination 3 x 1016 cm2 was minimum dose for steady-state sputtering

Compositional changes due to heating or sputtering occurTed only in surface layers 5 20 A

Ion sputtering restored stoichiometry of heated GaAs due to Sputtering of excess Ga

Technique not sensitive enough 295

353

168

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

TABLE 2 (continued) Summary of Published Results of Surface Compositional Changes in GaAs Due to Noble Gas Sputtering

GaAs surface

Analysis tech- Ion nlque specles

ISS, RBS,

PL AES

Ar

Ar

Ion energy [E(keV)I

110

1.5 3 5

Ion current Ion angle Sputtered densitv of incidence surface

100 75 100 75 100 75

XPS Ar 1.5 300 (normal 3 300 to sample) 5 300

XPS (1 0" from surface plane)

AES. c h a n n e I - ing

RBS. TED

AES

Ar 1.5 300 3 300 5 300

Ar Xe

Ar

Ar

0.05- 1.5

30

2

75 75 75

75 75 75

20-1 000 -

170-368 0

to -

Remarks Ref.

Radiation-induced vacancies were 349 much deeper than the range of implanted ions

AES: 345, Ga enrichment at all ion 346, energies 356

Sensitivity factors in Handbook of Auger Electron .Spectroscopy incorrect

Ga enrichment actually more pronounced

Postimplantation annealing reduced As depletion but did not restore stoichiometry

energy sputtering tended to anneal slightly more rapidly than 5 keV damage

Damage created by lower

XPS: With a deeper analyzing depth. the XPS and AES results were in qualitative agreement

With a shallow analyzing depth, the XPS data showed a Ga enrichment at the surface for 1.5 keV. while at higher energies (3 and 5 keV) the surface was stoichiometric or As enriched, but As depleted deeper in the substrate

The depths of depletion appeared to correlate with the ranges of the ions

Formation of elemental Ga and As Electrical measurements (I-V. C-V, DLTS) on in situ manufactured Au Schottky contacts after sputtering:

increased ideality factor n and reverse current I,

"Soft" diode characteristics -

Barrier height decreased Damaged layer was donor-like Comparison with XPS data indicated that the sputter-induced As vacancies are donor-like

Electrical changes were explained using a combined effective work function model and the creation of a donor-like surface damage layer

At 50 eV. very little Ga- 357 enrichment

Ga enrichment increased with increasing ion energy up to 1 keV for A r bombardment

incorporation Below 200 eV. no ion

RBS: GaAs sputtered at 225°C TED: GaAs sputtered at 443"C,

82

still crystalline 364

169

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13


Analysis Ion current Ion angle Sputtered denslty of incidence surface

[ i W c m l (7 composition Remarks R d .

- - As enrichment; Ga enrichment (E > 0.5 keV) 358 increased with increasing A r energy

Ga enrichment XPS peak shape analysis showed

Ion specks

Ar

G a b surface

t e c i niqw

(1W XPS. n-type u-v

reflectivity

that the degree of amorphization increased with increasing Ar. energy; at 4 keV. there was still some crystallinity left

U-V reflectivity measurements showed that at 4 keV, 45% of the near-surface region had been transformed to amorphous GaAs;

I-V Schottky contact parameters became increasingly "soft" with increasing Ar* energy, for diodes fabricated on Ar'- bombarded GaAs

Surface composition = GaAs. Subsurface depletion of As; width increased with increasing energy; at deeper depths, a return to GaAs

Proposed mechanism: preferential sputteringcaused by surface segregation enhanced by sputter- assisted diffusion in the near- surface region

Ga enrichment decreased exponentially

Compositional changes controlled by kinetic collision process

RHEED: sputtering caused amorphization; recrystallization after 430°C anneal

13.7% oxygen on 0.5-kV sputtered surface; less than 2% oxygen on all other surfaces

increased chemical reactivity compared to chemically cleaned surface: this reactivity increased with ion energy up to 2 kV

Fermi-level pinning position: before Ar', E, - 0.68 eV: after Ar'. E, - 0.51 eV

Ar: real barrier height reduction

Sputtered GaAs showed an

348

132

365

342

359

366

310 31 1

360 360

ARXPS Ar 1.5. - 3.5

20 GaAs

AES. Ar 2 RHEED

7.5 80 G a s o

GaAso 6SD

XPS Ar 0.5 t 2 3

(100) ICTS, AES n*, n = loi7 cm3

Ar 0.3

0.6 0.9 1.2 1.5 1.8 0.1-0.2 iss Ar At 550 K: As enrichment due to

segregation with Ar' bombardment

defect-related process

superlattice

depletion and deeper altered layer

subsurface region

As diffused to surface via a

GaAs substrate of AIGWGaAs

Higher energies - more As

Outer layer richer in As than

(100) undoped

Ne

Ar

Ar

5

AES 3

AES with mechanical shields; e. = 120

0.5 1

2 4

170

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13


Analysis GaAs surface

(110) undoped

(100) and ( 1 1 1 ) s.i.

(1 10) ( 100)

8 x 1017-

Si doped,

2 x lo', cm-3

technique

AES

XPS. e. = 720

Auger eleclron diffrac- tlon. XPS dlf- fraction

XPS

XPS

AES

Ion Ion current Ion anale Sputtered Ion energy

species [E(keV)1

5

0.5

1 2 4

5 Ar 4

Ar 1

3 Ar 1. 3

Ar 0.65 1.5 3 0.65 1.5 3 0.65 1.5 3 0.65 1.5 3

- density of incidence surface

[i(CIA/cm7 (7 composition

GaAs, ,,' Gdso 9,'

GaAs, MQ

GaAs, wQ

GaAsoa,'

GaAs, wb

55 GaA%,* GaAs, p7Q

lc-1000

1.6 0.2-40

10 20 40 10 20 40 10 20 40 10 20 40

a Low-energy Auger peaks were used in the calculation. High.energy Auger peaks were used in the calculation

et a1.280.369.370 showed that below the recrystallization temperature T,, Ga enrichment increased with increasing Ar+ bombarding energy and/or ion current density. This indicates that a segregation/ diffusion mechanism must operate during the sputtering of GaAs by argon ions. The series of studies by Valeri and co-workers310~311~34-1~3~~361 (see also Figure 17) confirm the ARXPS results and segregation model by Bussing et aLU8 The GaAs profiles shown in Figure 17 also indicate that the argon bombardment gave rise to an As-depleted region below the surface. The As depletion and the thickness of the altered layer increased with increasing ion energy.

45

45 45

18 18 18 42 42 42 51 51 51 70 70 70

Remarks Ref.

Composition profile of GaAs (1 00)

Altered layer depth + R,; 50 A at

XPS: pure Ga. As, and GaAs Auger electron diffraction: 0.5 keV,

= GaAs( 1 10)

1 keV. 110 A at 4 keV

structural order persisted; 4.5 keV. amorphized

profile dominated by As radiation- enhanced Gibbsian segregation at room temperature due to radiation-induced point defects

Sputter-induced composition

GaAs067 Steady state at 12 x lots Are cm-* 362, Ga-Ga bonds 367,

GaAs, 7, Annealing at 575 K: GaAs,,, 360 Ga 363

GaAs,, Results independent of crystal 327 GaAso7, orientation, (100). (110). and GaAs, ,,, vicinal planes

GaASO69 to02

GaAs,,, t OM

GaAs07, :ow

enrichment

GaA% 77 I 0 05

GaAsO03 I O M

GaAsO 73 I O M

GaA%Bl ,001

Gab, 7(1

GaAsO 73 z 0 M

The ISS studies by Oman-Rossiter et a1.359 and Barcz et a1.293~294 do not seem to support the above segregation model. Because ISS analyzes only the top surface layer, such studies should, according to this model, find only stoichiometric GaAs. Oman-Rossiter et a1.359 found that both Ar+ and Ne+ bombardment resulted in As-stabi- lized surfaces. In contrast, Barcz et a1.293.294 found a Ga-enriched surface after Ne+ bombardment. More ISS studies are needed to clarify these discrepancies.

The effect of crystal orientation on the final surface composition after sputtering was investigated by four groups280.327.~~~7.361 using (100)

171

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

Ga - 5

I I i

f 0 130 260 390 520 650 780 910 1040 1170 1300 1430 l!

KINETIC ENERGY (eV)

FIGURE 19. Typical differentiated Auger spectra obtained from (a) vacuum-cleaved (1 10) GaAs, (b) 1 -keV and (c) 3-keV Ar-sputtered (1 00) surfaces. (From Malherbe, J. B., Barnard, W. O., Strydom, I. Le R., and Louw, C. W., Surf. Interface Anal., 18, 491 , 1992. With permission.)

0

and (1 10) GaAs. Within experimental error, no crystal orientation effects could be detected in the results by Malherbe et al.327 Also, Singer et a1.280 did not report any effect, leaving the impression of orientation independence. McGuireM7 stated that the differences obtained for GaAs (1 11) and GaAs (100) crystals are probably within experi-

mental error (no error bars are given in his publication). Valeri and di Bonaw investigated Au- ger peak-to-peak height ratios of the high-energy Ga and As peaks of vacuum-cleaved GaAs (1 10) surfaces, as a function of the polar angle along the (110) plane (see Figure 15). On the cleaved surface, the peak-to-peak heights exhibited several

172

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

R(e.;r-7--$ 1.5 KeVAr+

.85 0

0

90 30 60 0

CAfl '*O I I

0.7 1 QaAs 0

R(8) I 90

0.6 30 60 0

cA!J l a 0 l 0.5

0 ~

0 0.5 1.0 1.5 2.0 R($or\ 0.6 0 0

0 30 60 90

0.5

0 0 0.5 1.0 1.5 2.0

"4

FIGURE 20. ARXPS results for Ar+-sputtered GaAs, using the Ga 3d and As 3d XPS peaks. Relative As intensity [IAE/(IAs + IGa)] is plotted as data points R(B), with curve fit [-] using the Laplace transform of the step-function CDP depicted (at.% As vs. relative depth, x/h): for Ar+ ion energies of (a) 1.5, (b) 3.0, and (c) 5.0 keV. Note the region of subsurface As depletion, and the progression of the depleted region to greater depth with increasing ion energy. (From Bussing, T. D., Holloway, P. H., Wang, Y. X., Moulder, J. F., and Hammond, J. S., J. Vac. Sci. Techno/., 86, 1514, 1988. With permission.)

distinct peaks and valleys. A 0.5-keV Ar+ sputter- Most studies280.3 10.31 1.327.342.343.345.346.348.356-

ing caused these diffraction effects to become less distinct but still to be present. This indicated that a structural' order was still present. However, a 4.5-keV Ar+ sputtering caused these peaks and valleys to disappear, indicating an amorphization of the GaAs substrate.

358.360,361,365,369.370 found that higher-energy Ar+ ions cause a greater depletion of As than lower-energy Ar'. The effect is especially noticeable at the lower energy. In Figure 21, all AES and X P S measurements on low-energy argon ion-sputtered GaAs given in Table 2 are plotted as a function of

173

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

1.2 x 0 a

.m CI,

1.0 a

Lc

m a 0.6

- on argon sputtered GaAs

L V

I S \

0.4 4 0 1 2 3 4 5 6

Argon Ion Energy E (keV)

FIGURE 21. Dependence of the surface composition of low-energy argon ion sputtering of GaAs as a function of the bombarding argon energy. Only AES and XPS measurements are shown. The solid line represent the best fit between the data and an exponential function energy term. The data were represented by and taken from: D, DeLo~ise3~2.36~.368; K, Kang et M, Malherbe et al.327; Mc. M c G ~ i r e ~ ~ ~ ; S, Sakalas, and Zhukauska~3~~; V, Valeri et aL310 311344360; Va, Van 00s t rom~~~; W, Wang and H o l I o ~ a y ~ ~ ~ ~ ~ ~ ~ 3 ~ .

the ion energy. Both these techniques (AES and XPS) have analyzing volumes ranging over several monolayers. Thus, they measure an average composition of these volumes and are unable to give direct proof of the segregation model by Bussing et al. However, in agreement with this model, the results exhibit the correct tendency, as a function of the ion energy. The curve in Fig- ure 21 represents the best fit between the data when fitted to an exponential function, i.e.,

f - 5

(22) -- LA’ = x = exp(4.088E) CLa

In Figure 21, the surface composition data were compared with each other, regardless of the angle of incidence of the bombarding argon ions.

Only two s t ~ d i e s ’ ~ ~ . ~ ~ have reported on the dependence of the sputtered surface composition on the angle of incidence of the ions. Malherbe et a1.327 used three different ion energies (0.65,

174

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

1.5, and 3 ke\-' ;It four different angles of incidence (1 8,42.5 i . and 70"). The results of the two extreme energizs (0.65 and 3 keV) are shown in Figure 22. The results of 1.5-keV Ar+ bombardment showed ~ t r : same tendency as that of the other two ion znergies. Within experimental error, the angle <.t' incidence of the ions has no effect on the s:z;Idy-state sputtered surface composition. M c G U ~ R ~ ~ ~ took measurements at two angles of incidence, 8 = 0 and 75", for 2-keV A P on both GaAs ( IlW) and GaAs (1 1 1). He reported a difference of S to 9% in the Auger peak height ratio of As/Ga \vith the ratio being larger for the glancing incidence ions on both crystal orientations. A comparison to the results of Malherbe et a1.327 in Figure 21 shows that the latter measurements are also \vithin this variation range.

The absence of angle dependence is in contrast to findings on InP312,313 and some other sys-

1 .o

0.9 X

0

0

Q5 0.8 0 c3 \ 4

.- +

cn

0.7

0.6

It probably also explains the broad agreement between all the published AES- and XPS-measured surface compositions of Ar+-sputtered GaAs determined on different instruments with different angles of incidence of the bombarding ions.

To summarize, low-energy noble gas sputtering of GaAs at room temperature results in a surface composition depleted in As. The main mechanism that causes this compositional change is radiation-enhanced segregation, probably in conjunction with preferential sputtering of the As species. In line with this segregation mechanism, Ga enrichment increases with increasing ion energy E, with the biggest increase between 0.65 and 1.5 keV. No sample orientation effects have been observed or any dependence on the angle of incidence of the bombarding Ar+ ions.

I I I I I I I

+ Ar - G ~ A s v .65 keV

3 keV

0 10 20 30 40 50 60 70 80

Angle of Inc idence of Ions ( ")

FIGURE 22. Effect of the angle of incidence of the argon ions on the As/Ga ratio for ion energies of 0.65 and 3 keV, respectively. (From Malherbe, J. B., Barnard, W. O., Strydom, 1. Le R., and Louw, C. W., Surf. Interface Anal., 18, 491, 1992. With permission.)

175

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

4.6. Sputter-Induced Compositional Changes in Other Compound Semiconductors

Sputter-induced compositional changes may be a disadvantage in one case and an advantage in another. For the electronic device industry, in most cases, it is beneficial to have a clean stoichiometric monocrystalline compound semiconductor substrate with known electronic properties. However, in some cases, a different surface layer has advantages. For example, good ohmic contacts on CdTe require a Te-enriched surface l a ~ e r . 3 ~ 1.372

To rectify the compositional changes and structural damage of sputter ion bombardment. a frequently used technique employed in LEED studies can be applied to compound semiconductors. This entails the use of several cycles of alternative sputtering of the surface with low- energy noble gas ions followed by an annealing step.373.374 For materials that can be cleaved in vacuum, it has been demonstrated that the LEED patterns are indistinguishable from cleaved surfaces or from surfaces treated as just explained. This method has consequently been used for several compound semiconductor surfaces, for example,GaSb,37s GaAs,376 PbTe,337 1r1Sb.'~~ CdS,379 CdTe,3R" etc. By means of a suitable choice of the various parameters, any compositional Changes induced by the sputtering process can be rectified by the annealing step. The crucial step in this method is the choice of the annealing parameters, as it can introduce unwanted artifacts. For example, too moderate a heat treatment may not anneal the radiation damage. An excessive heat treatment may cause segregation of impurities to the surface (e.g.. 450°C annealing causes segregation of carbon in GaSb38'), high temperatures may cause decomposition of the compound (e.g., above -460°C, InP starts to decompose, with loss of the volatile P specie^^'-^^), etc.

A large number of publications have dealt with the question of ion bombardment-induced compositional changes in compound semiconductor materials. In Table 3, the equilibrium sputtered surface compositions of binary compound semiconductors, other than InP and GaAs, are summarized for noble gas ion sputtering. This table does not cover all such published data for all

the compound semiconductors (it aims to be rep- resentative rather than comprehensive/exhaustive), not does it include noble gas bombardment-induced compositional changes on ternary and higher compound semiconductors.52~324~381~391~3w~41~37

Ion bombardment-induced compositional changes of oxide compound semiconductors (e.g., In,O, BaO, CdO, HgO, and ZnO) have been discussed p r e v i o ~ s l y , ' ~ ' , ' ~ . ~ ~ ~ . ~ ~ ~ and hence are not included in this table and discussion.

One noticeable feature of Tables 1 to 3 is that often there are large discrepancies among the results of different studies on the same material with the same technique. Hence, one must beware of making any predictions and establishing rules about sputter-induced compositional changes on the strength of a single study.

Noble gas sputtering of the 111-V semiconductors usually leads to either enrichment of the group I11 atomic species in the selvedge or to a stoichiometric surface layer. Little is known about the main mechanisms that cause the sputter-induced compositional changes in these compounds other than InP and GaAs. For example, based on purely preferential sputtering considerations, one would expect phosphorus to be preferentially sputtered in Gap. Both main factors determining preferential sputtering, viz., mass difference and surface-binding energy differences, favor the preferential sputtering of the lighter and weaker- bonded phosphorus atoms. This prediction also corresponds to most measurements presented in Table 3. However, ISS studies indicate the presence of stoichiometric GaP on the top surface layer, in agreement with a bombardment-induced segregation model.

The IV-IV compound semiconductor family is best represented in Table 3 by Sic. Few publications have appeared on the effect of sputtering on the composition of Sic. It is a material difficult to study. It is largely covalently bonded with a tetrahedral coordination. S i c occurs in over 100 polymorphs, which have similar atomic arrange- ments in the plane perpendicular to the symmetry axis (C-axis) but with different stacking sequences. The nearest-neighbor bonding is tetrahedral, but the second-nearest neighbors determine whether the structure is cubic, hexagonal, or rhombic.lM Some sputtering studies have been performed on amorphous S i c films or polycrystalline substrates.

176

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

TABLE 3 Summary of Published Results of Surface Compositional Changes in Compound Semiconductors Due to Noble Gas Sputtering

Analysis Ion Ion current Ion angle Semiconductor tech- Ion energy density of Incidence surface nlque species [Ep(keV)] [j(@cm) (9

CdS( CdS

- 120) AES Ar 1 54 XPS Ar 4 - 30

- 120) AES Ar 1

CdSe (111)

CdTe (100). I O O O l t

(111), (110), and others

CdTe CdTe single

CdTe(ll1) CdTe(ll1) and

crystal

CdTe( 1 10)

CdTe single crystal + polycrystalline

CdTe( 1 1 1 )

Te surface

polycrystalline

both Te and Cd surfaces

CdTe

CdTe(ll1)

CdTe( 1 10)

- AES Ar - XPS. Ar 1, 2, 6, - LEED 10. 15

- AES Ar - AES. Ar -

AES Ar 05-4 - AES Ne 1 12

- XPS

AES.

XPS XPS, AES. LEED

XPS

AES, LEED

AES. LEED

Cu,Te XPS

GaN XPS polycrystalline

GaP(m) P face AES

- Ar -

- Ar - Ar 1 -

Ar 4 -

Ar 0.6 -

Xe 0.6 4.5 1 2 3

1 2

Ar 0.6 4.5

J

Ne 0.6 4.5

1 2 3

Ar 4

Ar 2

-

15

Ar 0.9 -

Remarks

Sputter-induced compositional changes were independent of: prior surface treatment crystal orientation:

(0001) or A (cadmium) (Oooi) or 8 (sulfur)

Explanation for nonstoichiometry: forward .%altering of S below the surface and subsequent exposure of these atoms through continued sputtering

LEED patterns regardless of Ar ion energy and without annealing

Ref.

382 383 384

385

381

52 386

387 The Auger Cd(MNN) and Te(MNN) 388, peaks indicated no preferential sputtering, but the Te(NO0) and Te(MNN) peak ratios did indicate a Cd preferential spul- tering from a 5-A surface layer

LEED: amorphous surface after sputtering

LEED: diffuse patterns displaying enhanced threefold symmetry indicated significant sputter- induced damage

AES: very slight preferential sputtering of Cd, independent of ion energy and mass

LEED: good patterns after sputterring

No segregation or diffusion effects were detected

Preferential sputtering of nitrogen and production of metallic Ga; a surface layer >2 nm deep consisted of almost equal amounts of metallic Ga and GaN

The sputtered surface composition was independent of crystal orientation

389

390

196

383

192

391

383

392

350

177

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

TABLE 3 (continued) Summary of Published Results of Surface Compositional Changes in Compound Semiconductors Due to Noble Gas Sputtering

Analysis Semlconductor tech- surface nique

Gap( 1 1 1) Ga face UPS

GaP RBS GaP ISS

GaP(11l)B AES, disappearance polential, ionization spectroscopy

GaSb

GeSi HgTe (111). (110) and others

HgTe( 1 10)

ISS

RBS XPS, LEED

AES

lon Ion current ion angle Sputtered lon energy density of incidencs surface

species [Ep(keV)] [I(irAlcml] (9 composition

Gap, .3 Ar 0.9 - -

Ar 40 50 - GaP Ne 25 2.5 - GaP

Ar 0.1-2 t 0 Ga enrichment

Ne 25 2.5 - GaSb

Ar 40 -50 - GeSi Ar 1. 2, Te - -

5. 6. enrichment 9. 10. 15

Ne 0.6 1 2 3

Ar 0.6

4.5

4.5

Remarks

A similar study on GaAs showed only slight Ga-enrichment

Disappearance spectroscopy: Disordering of surface by Ar. due to a cascade process

At low doses, P vacancies made big contribution to disordering

High doses: decreased rate of defect formation due to recombination of vacancies with interstitial atoms andlor disappearance at the surface or dislocations. etc.

Amorphiration dose = 2 x 10’5 Ar- cm-’. but still some short- range order

Defect formation cross-section was proportional to ion energy

radiation caused an appearance of free states in the band gap and the decay of surface free states in the conduction band

Ionization spectroscopy:

XPS: Te enrichment increased with Ar- energy

Concentrations of Hg and Te did not add up to loo%, with a maximum deficit of 15?4 at 15 keV

LEED patterns became increasingly blurred and disappeared at higher Ar ion energies

Sputtering caused a high concentration of surface and subsurface Hg vacancies; these acted as sinks for deeper Hg atoms, which diffused due to radiation-induced defects

Low-energy Ar ions produced thinner layers with larger gradients than Ar ions of higher energy, due to the smaller range of lower ions

which increased with ion energy and ion mass

preferential sputtering effects

Preferential sputtering of Hg.

Depth profiles indicated only

Ref.

393 293. 294

394

293, 294

393 38 1

39 1

178

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13


Analysis Ion Semiconductor tech- Ion energy surface nlque species (Ep(keV)]

lnAs

InAs(ll1)

lnSb

lnSb(ll1) (iii)

ISS

XPS. ELS ISS

AES

InSb( 1 1 l )A AES lnSb ( i i i ) B A and B surfaces AES lnSb AES

InSb(ll1) XPS. ELS

InSb( 100) AES

lnSb(100)

PbTe(ll1) n. p = lo1’ -1 0” cm-3

PbTe Sic

p-SiC( 100) n-type

n.p = l O I 7 -10” cm-)

SiC(OOO1)

Xe

Ne

Af

Ne

Ne

Ne

Af -

Af

Xe Ne Af Kr

AES, Kr LEED

XPS. Ar LEED

AES Ar Ne

AES, Ar ELS, Kr LEED

LEED

1 2 3 0.6 1 2 3 25

E S 6?

25

15

15

3 -

E 5 6?

2 0.5 2 2 0.3- 0.5

3

2 0.25-8

E < 0.7

AES He 3

Ion current density

[ i ( W c W

4.5

2.5

- 2.5

1000

1000

0.45 -

-

- - - - -

-

- 1000

-

Sic( 100) AES Ar 0.5, -

3C-Sic . AES, Ar 0.5, - 5

(100) LEED 3

Ion angle of Incidence

(3

49

-

-

-

30

30

-

-

75 0 75 75 -

-

60 -

-

-

-

60

Sputtered surface

composition

HgTe1,,0,

HSTei M t om w e 1 I 6 z 0 cd

HgTelMzOC6 HgTel I 7 t Om

HgTe127?0m HgTeO 16 f 0 M lnAs

InAs, 5

lnSb

In ennchment

InSb,,

InSb, , In

Ins45

enrichment

InSb,,

InSb,,

InSb, InSb,, In

InSb, 13

enrichment

PbTe

PbTe, SI enrichment

C enrichment

C enrichment

-

C ennchment

Remarks

No segregation or diffusion- effects were detected

ELS: In islands after annealing at 300°C for 1 h

Sputtering at elevated temperatures (T 2 350%); results showed a spatial anisotropy in the yields of In and Sb due to single crystal effects

Excess of 10-20 at% In; void formation and anomalous swelling of the irradiated layer

ELS: In islands after annealing at 300% for 1 h

Auger p-t-p ratio of SbAn decreases with increasing B

(1 x 1) LEED pattern with high background; an improved LEED pattern with annealing up to 250°C

LEED: Ar. bombardment did not destroy crystal structure on surface

Comparison of sputtering yield measurements and theory suggests enrichment of heavier species

LEED: Simultaneous 1 -keV bombardment and 650°C annealing gave best LEED patterns of the bombarded samples; room-temperature ion bombardment followed by 700°C anneal yielded no LEED patterns

in room-temperature. ion- bombarded samples

C enrichment in near-surface region was dose dependent; Si enrichment in deeper layers Conclusion: radiation-induced segregation

Si-C bonds were broken by argon bombardment

After ion bombardment, there was a loss of short-range order by disruption of Si-C bonds

LEED: patterns only obtained on samples that were simultaneously ion bombarded by low-energy ions and heated

ELS: SIC bond peak disappears

Ref.

293, 294

98

293, 294

395

396

397

98

347

398

399

400 401

402

403

404

405

1 79

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13


Analysis Ion Ion current Semiconductor tech- Ion energy density surface nique species [Ep(keV)] [i(llrvcmq

a-Sic, 22 AES Ar

a-Sic AES Ar

SiGe XPS Ar

SnTe(ll1) XPS. Ar (100)

p - 2 Y 1020 cm3 LEED

3 7

2. 4. 30 keV: 30 3.2

- 3

3 -

The data given in Table 3 are often contradictory. Some studies show Si enhancement on the surface, whereas others show C enhancement. Most studies indicate that noble gas sputtering causes amorphization of the substrate, with Si-C bonds being broken. Confirmation of the bond- breaking mechanism comes from studies already discussed in Section 3.2.3, viz., change in microhardnes~,~’~ reduction in the friction coefficient and wear rate due to possible formation of a lubricating amorphous carbon film,186 and an increase in the oxidation and chemical etching rate”? of noble gas ion-bombarded Sic. However, in a recent experimental study, Miotello et a1.108 found that Si-C bonds at the surface undergo a change of bond type. They concluded that sputter-induced changes in S i c are due to the chemically guided motion of defects rather than ballistic effects.

The 11-VI semiconductor family is represented by three widely different semiconductors - CdS, CdTe, and HgTe - with respect to their sputter behavior. Based on preferential sputtering considerations, one would expect to find Cd enrichment on the surface of noble gas ion-sputtered CdS. However, all experimental studies indicate the opposite. This probably means that sputter-

Ion angle Sputtered of incidence surface

(9 composition Remarks Ref.

- SiGe

- SnTe

Sputter-deposited layer 406, 407

C accumulated at the surface 408 without depletion in the subsurface region

Si also accumulated at the surface with Si depletion near the Ar. range

Electron bombardment caused C movement due to a positive surface charge

Si-C bonds were formed at the surface, resulting in no change of bond type

Composition changes were due to the chemically guided motion of defects rather than ballistic effects

409

399 LEED: Ar. sputtering did not destroy the crystal structure on the surface

induced segregation or diffusion effects are the dominant mechanisms involved in the compositional changes due to sputtering. However, the studies cited in Table 3 do not provide any proof of this conjecture.

With respect to sputter-induced changes, the semiconductor CdTe has a behavior more in common with metal alloys than with the covalent bonded 111-V compound semiconductors. It is not easily amorphized by sputtering. Most of the studies also indicate that there are no sputter-induced compositional changes in CdTe. As indicated by LEED studies, i t remains monocrystalline even to the top surface layers. AES and XPS indicate that the surface remains stoichiometric. The commonly used high-energy peaks of Cd and Te usually measured in both AES and XPS imply relatively large electron attenuation lengths. Thus, these techniques analyze several monolayers of the selvedge, and not just the top surface layer. In an AES study by Lu et aI.,I9’ no compositional changes were observed after neon ion sputtering when the high-energy Cd (MNN) Auger peak at 376 eV and the Te (MNN) Auger peak at 483/491 eV were used. However, when the low- energy Te (NOO) Auger peak at 3 1 eV was used

180

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

together with the Te (MNN) peak, a Te-enriched top surface layer was obtained.19’ Clearly, more studies using other top surface-sensitive techniques, such as ISS, are needed to clarify this matter.

Noble gas ion bombardment of HgTe leads to a loss of the Hg species in the substrate. The HgTe monocrystal is also structurally damaged by the bombardment, as manifested by LEED patterns.381 The degradation of the LEED patterns increases with increasing ion energies. The Te enrichment on the target surface increases with increasing ion energy and ion mass.381.391 Stahle et al.391 investigated the possibilities of ion bombardment-induced Gibbsian segregation and bombardment-induced diffusion. They concluded that preferential sputtering was the main mechanism for compositional changes in HgTe.

4.7. Summary

Many more studies are needed to clarify the mechanisms by which ion bombardment-induced compositional changes occur in the different compound semiconductors. It is impossible to make generalized predictions on ion bombardment-induced compositional changes because different semiconductors behave differently, even within the same family. This would seem to be due to the fact that many different mechanisms can, and probably do, operate simultaneously in most compound semiconductors.

5. CONCLUSIONS

Ion bombardment of InP leads to severe topography development. On clean surfaces, the bombardment-induced topography is in the form of small cone-like protuberances. The sizes and areal density distribution of these cones depend on the ion species, ion dose density, ion energy, and angle of incidence of the bombarding ions. Ion bombardment of GaAs leads to ripple development, with roughly the same dependence on the ion beam parameters as InP. There are few systematic studies on bombardment-induced topography development on other compound semiconductors.

The mechanisms and effects of radiation- induced damage in compound semiconductors

depend strongly on the semiconductor group. The 111-V semiconductors are usually easily amorphized at low-dose densities (on the order of lOI4 cm-2), but there is also evidence of some recrystallization taking place at higher dose densities. The radiation damage behavior of the 11-VI and HgCdTe semiconductors is very similar to that of metals.

Several mechanisms can cause bombardment- induced compositional changes in multicomponent materials. In the case of argon bombardment of InP, the experimental results indicate that the cause of surface compositional changes is mainly preferential sputtering, while for argon bombardment of GaAs, bombardment-induced diffusion and segregation effects are the dominant mechanisms. In general, few other compound semiconductor systems have been investigated thoroughly. This makes it difficult to ascertain the particular compositional change mechanisms that operate in most compound semiconductor systems.

This review ends with a few general remarks. Ion-solid interactions encompass several different and complex mechanisms. A few of these interactions are reasonably well understood, and one often finds good agreement between experimental results and theoretical predictions. The following topics fall into this category: physical sputtering in the linear cascade regime, the ranges of ions in matter, sputter-induced compositional changes in GaAs and InP, and radiation damage in several compound semiconductor systems.

There is, however, disagreement between experiment and theory in several of the topics covered in this article. The degree of discrepancy varies from slight for threshold sputtering to severe for radiation-induced topography. Other topics included in this broad category are spike sputtering (spike temperature) and heavy ion bombardment-induced recrystallization of radiation damage in compound semiconductor systems. Further study of recrystallization, topography development, and compositional changes is of particular topical interest.

ACKNOWLEDGMENTS

The author expresses his sincere appreciation and gratitude toward Dr. Werner Gries (Darmstadt)

181

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

for reading the document and making excellent 15. Wehner, C. K., Cone formation as a result of whisker w Y

growth on ion bombarded metal surfaces, J. Vac. Sci. Technol., A3, 1821, 1985.

16. Whitton, J. L., Experimental studies of morphology development, in Erosion and Growth of Solids Stimu-

comments and suggestions for improvements. The extensive input of many colleagues, Mane Man, my father, and family is also gratefully acknowl- edged.

REFERENCES

1.

2.

3.

4.

5 .

6.

7.

8.

9.

10.

I I .

12.

13.

14.

1 82

Malherbe, J. B., Sputtering of compound semiconductor surfaces. I. Ion-solid interactions and sputter yields, CRC Crit. Rev. Solid State Phys. Muter. Sci.. 19. 55, 1994. Hofmann, S., Practical Surface Analysis by Auger and X-Ray Photoelectron Spectroscopy. Briggs. D. and Seah. M. P., Eds., John Wiley & Sons, Chichester, 1983, chap. 4. Zalar, A., Improved depth resolution by sample rota- tion during Auger electron spectroscopy depth profiling, Thin Solid Films, 124. 223, 1985. Pamler, W. and Wangemann, K., Depth profile analysis of aluminum metallization in microelectron- ics: optimization of depth resolution. Sirrf. Interface Anal., 18, 52. 1992. Sobue, S. and Okuyama, F., Topographical features of rotated Auger samples sputtered with inert-gas ions, J. Vac. Sci. Technol., A8, 785, 1990. Kojima, I., Fukumoto, N.,and Kurahashi, M., STM and AES study of the relation between ion sputter induced roughness and depth resolution, J . Electron. Specfrosc., 50, C9, 1990. Behrisch, R., Ed., Sputtering by Particle Bombard- ment 11. Springer-Verlag, Berlin, 1983. Behrish, R. and Wittmaack, K., Eds., Sputtering by Particle Bombardment Ill, Springer-Verlag. Berlin, 1991. Kiriakidis, C., Carter, G., and Whitton, J. L., Eds., Erosion and Growth of Solids Stimulated by Atom and Ion Beams. Martinus Nijhoff, Dordrecht, 1986. Auciello, 0. and Kelly, R., Eds., /on Bombardment Modification of Surfaces, Elsevier, Amsterdam, 1984. Michely, T. and Comsa, C., The scanning tunneling microscope as a means for the investigation of ion bombardment effects on metal surfaces. Nucl. Instrum. & Methods, 882, 207. 1993. Wilson, I. H., Scanning tunnelling microscope studies of sputtering, Surf. Interface. Anal.. 20,637, 1993. Rossnagel, S. M., Effects of surface impurities and diffusion on ion bombardment-induced topography formation, in Erosion and Growth of Solids Stimu- lated by Atom and /on Beams, Kiriakidis, G., Carter, G.. and Whitton. J. L., Eds., Martinus Nijhoff, Dordrecht, 1986. 181. Carter, G., Navinsek, B.,and Whitton, J. L., Heavy ion sputtering induced surface topography development, in Sputtering by Particle Bombardment II, Behrish, R., Ed., Springer-Verlag. Berlin, 1983, 640.

lated by Atom and Ion Beams, Kiriakidis, G., Carter, G.. and Whitton. J. L.. Eds., Martinus Nijhoff, Dordrecht, 1986. 15 1.

17. Nindi, M. M., Stulik, D. C., Millham, C. B., and Larson, T. J., Three-dimensional computer simulation of shape and groove formation of a developing cone under energetic Xe atom particle bombardment, Nucl. Instrum. Methods, B44. 79. 1989.

18. Rossnagel, S. M. and Robinson, R. S., Surface diffusion activation energy determination using ion beams, J . Vac. Sci. Technol., 20, 195. 1982.

19. Oechsner, H., Sputtering - a review of some recent experimental and theoretical aspects, Appl. Phys.. 8, 185. 1975.

20. Oechsner, H., Untersuchungen zur Festkorperzer- staubung bei schiefwinkligem Ionenbeschub poly- kristalliner Metalloberfllchen im Energiebereich um 1 keV, Z. Phys.. 261. 37, 1973.

2 I . Sigmund, P., A mechanism of surface micro-roughening by ion bombardment, J. Muter. Sci.. 8, 1545, 1973.

22. Carter, G., Theory of surface erosion and growth, in Erosion and Growth of Solids Stimitlared by Atom arid Ion Beams. Kiriakidis. G. . Carter, G., and Whitton. J. L., Eds., Martinus Nijhoff, Dordrecht, 1986, 70.

23. Nobes, M. J., Geometric methods of analysis, in Ero- sion and Growth of Solids Stimlrlated by Atom and loti Beams, Kiriakidis. G.. Carter, G., and Whitton, J . L.. Eds., Martinus Nijhoff, Dordrecht, 1986. 103.

24. Carter. C. and Nobes, M. J., The theory of development of surface morphology by sputter erosion processes, in Ion Bonibarilment Modification of Sirrfaces. Auciello, 0. and Kelly, R., Eds.. Elsevier, Amster- dam, 1984. 163.

25. Carter, C., Nobes, M. J., Katardjiev, 1. V., and Whitton, J. L.. The role of bombardment induced defects in the initiation of topography on ion bombarded f.c.c. metals, Defect Diffits. Forirm. 57-58.97. 1988.

26. Cries, W. H., Radiation-induced sample modification in surface analysis: InP as an extreme example, Surf. Interface Anal.. 14, 61 I , 1989.

27. Cries, W. H. and Miethe, K., Synergetic effects during sputter-assisted depth profiling: growth-dominated topography development on InP and a model of the atomic mechanism, Microchim. Acra. I , 169, 1987.

28. Hiisler, W. and Pamler, W., Effects of crystallinity on depth resolution in sputter depth profiles, Surf. Interface Anal.. 20. 609, 1993.

29. Pamler, W., Wangemann, K., Kampermann, S., and Hosler, W., Depth resolution in Auger depth profile analysis of aluminum metallization in micro- electronics: the effect of crystalline texture, Nucl. Instrum. Methods, B5 I , 34, 1990.

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

30. Bradley, R. M. and Harper, J. M. E., Theory of ripple topography induced by ion bombardment, J. Vuc. Sci. Technol., A6, 2390. 1988.

31. MacLaren, S. W., Baker, J. E., Finnigan, N. L., and Loxton, C. M., Surface roughness development during sputtering of GaAs and InP evidence for the role of surface diffusion in ripple formation and sputter cone deve1opment.J. Vuc. Sci. Technol.. A 10.469, 1992.

32. Miethe, K., Schwarzbach, D., Betz, W., Nickel, H., Losch, R., Schlapp, W., Stingeder, G., and Grasserbauer, M., SIMS depth profiling of an AI,,,,G~,,,As/GaAs superlattice with 0; and Cs+ primary ions using different energies and angles of incidence, in Proc. SIMS VIII, Benninghoven, A., Janssens. K. T. F., Tumpner, J.. and Werner, H. W., Eds.. John Wiley & Sons, Chichester, 1991, 395.

33. Carter, G., Katardjiev, I. V., and Nobes, M. J., The production of periodic surface features by sputtering and related processes, Defect Diffus. Forum. 57-58. 49, 1988.

34. Auciello, 0. and Kelly, R., On the role of the primary beam and of scattered or sputtered particles in the faceting of cones on bombarded surfaces, Nucl. Instrum. Methods. 182/183, 267. 198 I.

35. Belson, J. and Wilson, I. H., Theory of redeposition of sputtered flux on to surface asperities, Nucl. Instrum. Methods. 182/183, 275, 1981.

36. Wilson, I. H., Belson, J., and Auciello, O., Second- ary effect in ion bombardment-induced erosion. in Ion Bombardment Modification of Sirifaces. Auciello, 0. and Kelly, R.. Eds.. Elsevier, Amsterdam. 1984.225.

37. McDonell, W. R., Whisker growth - a new mechanism for helium blistering of surfaces in complex radiation environments, J . Nirrl. Mater.. 76/77, 258, 1978.

38. Floro, J. A., Rossnagel, S. M., and Robinson, R. S., Ion-bombardment-induced whisker formation on graphite, J. Var. Sci. Technol.. A l , 1398, 1983.

39. Poker, D. B., Schubert, J., Alexandrou, A.. Froehlingsdorf, J., and Stritzker, B., Growth of Sn whiskers after low temperature implantation of 20 keV He or H, Nucl. Instrum. Methods. B19/20, 185, 1987.

40. van Vechten, J. A., Solberg, W., Batson, P. E., Cuomo, J. J., and Rossnagel, S. M., Kink site saturation mechanism for whisker growth under sputtering conditions, J. Cryst. Growth, 82, 289, 1987.

41. Bauser, E. and Strunk, H., Dislocations as growth step sources in solution growth and their influence on interface structures, Thin Solid Films, 93, 185, 1982.

42. Malherbe, J. B., Lakner, H., and Cries, W. H., Composition and structure of ion-bombardment-induced growth cones on InP, Surf. Interface Anal., 17, 719, 1991.

43. Cries, W. H., Recrystallization-affected topography on semiconductors at SIMS conditions, Proc. SIMS VIII, Benninghoven, A., Janssen, K. T. F., Tiimpner, J., and Werner, H. W.. Eds.. John Wiley & Sons, Chichester, 1991, 323.

44. Cries, W. H., On the Origin of Topography Develop- ing on keV Ion-Irradiated InP and GaAs. Poster EMJ- ThP19, 38th Natl. Symp. American Vacuum Society. Seattle, 1991. unpublished.

45. Cries, W. H., On the surface-analytical implications of commonly used simplifying assumptions about (wafer) surfaces, S. Afr. J . Phys., 16. 95. 1993.

46. Wilson, I. H., The topography of sputtered semiconductors, Rudiut. E’., 18, 95. 1973.

47. Wada, O., Ion-beam etching of InP and its application to the fabrication of high radiance InGaAsPAnP light emitting diodes,J. Electrochem. Soc., 131,2373, 1984.

48. DeMeo, N. L., Donnelly, J. P., O’Donnell, F. J., Geis, M. W., and O’Connor, K. J., Low power ion- beam-assisted etching of indium phosphide. Nucl. Instrum. Methods. B7/8, 814, 1985.

49. Linders, J., Niedrig, H., Sebald, T., and Sternberg, M., The sputtering of InP with 3 keV Ar- and 0,-ions under various angles, Nucl. Instrum. Methods, B 13, 374. 1986.

50. Bouadma, N., Devoldere, P., Jusserand, B., and Ossart, P., Ion beam etching and surface characterization of indium phosphide, Appl. Phys. Lett.. 48, 1285. 1986.

5 1. Dowsett, M. G., King, R. M., and Parker, E. H. C., SIMS evaluation of contamination on ion-cleaned (100) InP substrates, Appl. Phys. Lett.. 31, 529, 1977.

52. Farrow, R. F. C., The use of ion beams in molecular beam epitaxy, Thin Solid Films. 80, 197. 1981.

53. Oliver, J., Faulconnier, P., and Pokier, R., Elec- tron-stimulated oxidation on InP (100) surfaces studied by Auger electron spectroscopy and by low- energy-loss spectroscopy, J. Appl. Phys.. 51, 4990, 1980.

54. Oliver, J., Low-energy electron losses by plasmon excitations as control means of the growth of epitaxial layers, J. Appl. Phys., 5 1, 5406, 1980.

55. Kirk, D. L. and Jones, C., The influence of chemical treatment upon the composition and nature of (100) surfaces of indium phosphide, J . Phys. D . 12, 651, 1979.

56. Skinner, D. K., New method of preparing (100) InP surfaces for Schottky barrier and ohmic contact formation, J. Electron. Muter.. 9. 67. 1980.

57. Skinner, D. K., Swanson, J. G., and Haynes, C. V., Ion beam induced composition changes during Auger sputter profiling of thin A! films on InP, Surf. Inter- face Anal., 5 , 38, 1983.

58. Liau, Z. L. and Zeiger, H. J., Surface-energy-induced mass-transport phenomenon in annealing of etched compound semiconductor structures: theoretical modeling and experimental confirmation, J. Appl. Phys., 67, 2434, 1990.

59. Peacock, D. C., An AES. SAM and RHEED study of InP (1 10) subjected to ion bombardment and annealing treatments, Vucuum, 33. 601. 1983.

60. Hou, X.-Y., Yu, M.-R., and Wang, X., The variation of In islands on InP surface versus ion sputtering

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

angle observed by electron energy loss spectroscopy, Chin. Phys. Letr.. 2. 31, 1985.

61. Jardin, C., Robert, D., Achard, B., Cruzza, B., and Pariset, C., An AES and ELS study of InP (100) surface subjected to argon ion bombardment, Su$ Interface Anal., 10, 30 I , 1987.

62. Treichler, R., Cerva, H., Hosler, W., and Criegern. R. V., Primary beam damage at the bottom of SIMS craters in GaAs, InP and Si: a TEM and AES study, in Proceedings SIMS 7, Benninghoven, A.. Evans, C. A., McKeegan, K. D., Storms, H. A., and Werner, H. W., Eds.. John Wiley & Sons, 1990,259.

63. Malherbe, J. B. and van der Berg, N. C., Argon bombardment-induced topography development on InP, Surf. Intedace Anal., in press.

64. Wada, O., Ar ion-beam etching characteristics and damage production in InP, J . Phys. D , 17,2429, 1984.

65. Nagata, F., Shimotsu, T., and Kakibayashi, H., Cross-sectional observation with transmission electron microscope, JEOL News, 22, 14, 1985.

66. Duncan, S., Smith, R., Sykes, D. E., and Walls, J. M., Surface morphology of Si (100). GaAs (100) and InP (100) following 0; and CS' ion bombardment, Vacuum. 34, 145. 1984.

67. Miethe, K. and Kuphal, E., Sudden surface roughening followed by levelling observed in SIMS analysis of InP-containing Ill-V multilayers, Surf. Interfuce Anal., 20, 742, 1993.

68. Chew, N. C. and Cullis, A. C., Iodine ion milling of indium-containing compound semiconductors, Appl. Phys. Lett., 44, 142, 1984.

69. Mutoh, K., Nakajima, M., and Mihara, M., Reac- tive ion-beam etching of InP with CP?, Jpn. J. Appl. Phys.. 29. 1022, 1990.

70. Takimoto, K., Ohnaka, K., and Shibata, J., Reac- tive ion etching of InP with Br,-containing gases to produce smooth, vertical walls: fabrication of etched- faceted lasers, Appl. Phys. Lett., 54, 1947. 1989.

71. Tadokoro, T., Koyama, F., and Iga, K., A study on etching parameters of a reactive ion beam etch for GaAs and InP, Jpn. J . Appl. Phys.. 27. 389. 1988.

72. Donnelly, V. M., Flamm, D. L., Tu, C. W.. and Ibbotson, D. E., Temperature dependence of InP and GaAs etching in a chlorine plasma, J. Electrochem. Soc.. 129, 2533. 1982.

73. Homma, Y., Okamoto, H., and Ishii, Y., Surface microtopography and compositional change of cesium-ion-bombarded semiconductor surfaces, Jpn. J . Appl. Phys., 24, 934, 1985.

74. Farrow, R. F. C., Cullis, A. C., Grant, A. J., and Pattison, J. F., Structural and electrical properties of epitaxial metal films grown on argon ion bombarded and annealed (001) InP, J. Cryst. Growth. 45.292, 1978.

75. Cullis, A. C. and Farrow, R. F. C., A study of the structure and properties of epitaxial silver deposited by atomic beam techniques on (001) InP, Thin Solid Films. 58, 197, 1979.

76. Barna, A. and Pkz , B., Simple method for the preparation of InP based samples for TEM investigation, J. Electron Microsc. Technol., 18, 325. 1991.

77. Farrow, R. F. C., The evaporation of InP under Knudsen (equilibrium) and Langmuir (free) evaporation conditions, J . Phys. D . 7, 2436, 1974.

78. Massies, J. and Lemaire-Dezaly, F., Correlation of the tine structure of In,, M4,5 N4,5 Auger spectrum and thermal degradation at InP surfaces, J . Appl. Phys.. 55.3136, 1984.

79. Massies, J. and Lemaire-Dezaly, F., Auger electron spectroscopy and electron loss spectroscopy compara- tive study of vacuum annealing effects on InP surface, J . Appl. Phys.. 57, 237, 1985.

80. Wilson, I. H., Belson, J., and Auciello, O., Second- ary effects in ion bombardment induced surface erosion, in Ion Bonrbardment Modification of Surfaces. Auciello. 0. and Kelly, R., Eds., Elsevier. Amster- dam, 1984.225.

8 I . Hoinkis, O., Miethe, K., Betz, W., and Cries, W. H., Application of infrared luminescence microscopy for ion beam diagnostics and for measurement of ion dose densities from 10" to loN6 cm-', Fieseniirs J. A n d . Chem.. 341, 101. 1991.

82. Bhattacharya, S. R., Chose, D., Basu, D., and Karmohapatro, S. B., Surface topography of Ar' bombarded GaAs (100) at various temperatures, J. Vac. Sci. Technol.. A5. 179. 1987.

83. Bhattacharya, S. R., Chose, D., and Basu, D., En- ergy dependence of the sputtering yield of gallium arsenide by argon ion bombardment. lndiun J. P ure Appl. Phys.. 25. 328. 1987.

84. Bhattacharya, S. R., Chose, D.,and Basu, D., Mass and energy dependence of the sputtering yield of gallium arsenide, Nircl. Itrstrimr. Methods. B47.253. 1990.

85. Creene, J. E.. Natarajan, B. R., and Sequeda- Osorio, F., Sputtering of metal alloys containing second-phase precipitates. J. Appl. Ph.vs.. 49. 41 7, 1978.

86. Walkow, B., Loeb, H. W., Freisinger, J., and MoBner, C., Studies of GaAs surface roughness and organic masks resistance depending on the ion-beam energy, Fieseriitrs J . A n d . Cheni.. 34 1. 248. I99 I .

87. Zekorn, S., Fries, Th., Becker, C., and Wandelt, K., Topography of sputtered semiconductor surfaces by STM in Air, PAys. Status Solidi A. 131. 99, 1992.

88. Okuyama, F. and Fujimoto, Y., Evidence for atomic recondensation on ion-bombarded solid surfaces. Surf. Sci.. 210, L223, 1989.

89. Cries, W. H. and Lakner, H., private cornmunica- tion. 1993.

90. Cericke, M., Lill, T., Trapp, M., Richter, C.-E., and Hupfer, A., Surface roughening during depth profiling by secondary ion mass spectrometry (SIMS) in GaAlAs and GaAs. Fresenius J . Anal. Chem.. 341, 31. 1991.

91. Stevie, F. A., Kahora, P. M., Simons, D. S., and Chi, P., Secondary ion yield changes in Si and GaAs due to topography changes during 0; or Cs' ion bombardment, J . Vac. Sci. Technol.. A6. 76, 1988.

92. Friedbacher, C., Schwarzbach, D., Hansma, P. K., Nickel, H., Crasserbauer, M., and Stingeder, C., A combination of atomic force microscopy and second-

184

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

ary ion mass spectrometry for investigation of AI,Ga,.,As/GaAs super lattices, Fresenius J. Anal. Chem., 345,615, 1993.

93. Karen, A., Okuno, K., Soeda, F., and Ishitani, A., A study of the secondary-ion yield change on the GaAs surface caused by the 0; ion-beam-induced rippling, J. Vac. Sci. Technol., A9, 2247. 1991.

94. Duncan, S., Smith, R., Sykes, D. E., and Walls, J. M., Summary abstract: surface topography of electronic materials following oxygen and cesium ion bombardment, J. Vac. Sci. Technol., Al, 621, 1983.

95. Schimmel, Th., Fuchs, H., Graf, K., Sander, R., and Lux-steiner, M., Annealing behaviour of ion- bombarded WSe,, P h y . Status Solidi A, 131, 89. 1992.

96. Schimmel, Th., Fuchs, H., Sander, R., and Lux- Steiner, M., Atomically resolved STM imaging of ion-bombarded WSe,. Ultramicroscopy. 4 2 4 , 6 8 3 , 1992.

97. Schimmel, Th., Fuchs, H., and Lux-Steiner, M., Generation and manipulation of atomic-scale structures with the STM, Phys. Srarus Solidi A, 131, 47, 1992.

98. Hou, X., Yu, M., and Wang, X., Experimental evidence for indium island formation on clean InAs ( I l l ) and InSb ( 1 1 1 ) surfaces, Chin. Phys., 7, 854, 1987.

99. Carr, B. A., Friedland, E., and Malherbe, J. B., Effect of Ar+ ion bombardment on the electrical characteristics of Al/n-Si contacts, J. Appl. Phys.. 6 I , 2566, 1987.

100. Carr, B. A., Friedland, E., and Malherbe, J. B., Effect of annealing on the Schottky barrier height of Al/n-Si Schottky contacts after Ar+ ion bombardment, J. Appl. Phys.. 64, 4775. 1988.

101. Malherbe, J. B., and Friedland, E., Myburg, G., Carr, B. A., Bredell, L. J., and Pavlovska, A., Modification of metal Schottky contacts on silicon by ion implantation, Nucl. Instrum. Methods. B35, 247, 1988.

102. Malherbe, J. B., Weimer, K. P., Friedland, E., and Bredell, L. J., The effect of silicon ion beam mixing on the barrier height of Sb/n-Si Schottky contacts, Appl. SUI$ Sci.. 70/71. 442, 1993.

103. Vaseashta, A., Elshabini-Riad, A., and Burton, L. C., Effects of low energy AP ion bombardment on GaAs, Muter. Sci. Ens.. B9, 489, 1991.

104. Wendler, E., Wesch, W., and Gotz, G., Defect production in ion implanted GaAs, GaP and InP, Nucl. Instrum. Methods, B55, 789, 1991.

105. Wang, Y. G. and Ashok, S., Gallium arsenide surface modification by low-energy argon and nitrogen ion implantation, Nucl. Instrum. Methods. B39, 461, 1989.

106. Shwe, C. and Gal, M., Measurement of damage profile in semiconductors: a sensitive optical technique, Appl. Phys. Lett.. 62, 5 16. 1993.

107. Pribat, D., Dieumegard, D., Croset, M., Cohen, C., Nipoti, R.,Siejka, J., Bentini, G. G.,Correra, L., and Servidori, M., Ion implantation of silicon in gallium arsenide: damage and annealing charac-

terizations. Nucl. Instrum. Methods. 209/210. 737, 1983.

108. Malherbe, J. B., Weirner, K. P., Bredell, L. J., Friedland, E., and Myburg, C., Ion beam mixing of Sb Schottky contacts on n-Si, Muter. Res. SOC. Symp. Proc.. 157. 221, 1990.

109. Myburg, G., Malherbe, J. B., and Friedland, E., Ion bombardment of nickel contacts on silicon, Nucl. Instrum. Methods, B35, 457, 1988.

110. Auret, F. D., Malherbe, J. B., Net, M., and Myburg, G., A DLTS study of defects formed in silicon during ion beam mixing, Nucl. Instrum. Methods, B35, 234, 1988.

1 1 1. Malherbe, J. B., de Witt, B., and Berning, G. L. P., The effect of Ar' ion implantation on the electrical characteristics of Cr/p-Si Schottky diodes, J. Appl. Phys.. 71, 2757, 1992.

112. Pearton, S. J., Ion implantation in GaAs. Solid State Phenom., 182, 247, 1988.

1 13. Pearton, S. J., Ion implantation for isolation of 111-V semiconductors, Muter. Sci. Rep., 4, 313. 1990.

114. Pearton, S. J., Chakrabarti, U. K., Perley, A. P., and Jones, K. S., Ion milling damage in InP and GaAs. J. Appl. Phys., 68, 2760, 1990.

115. Wendler, E., Wesch, W., and Gotz, G., Investiga- tion of the amorphization process in ion implanted A,,,Bv compounds, Nucl. Instrum. Methods, B63.47, 1992.

1 16. Sadana, D. K., Mechanisms of amorphization and recrystallization in ion implanted 111-V compound semiconductors, Nucl. Instrum. Methods, B7/8, 375, 1985.

17. Gill, S. S., Rapid isothermal annealing of ion implanted gallium-arsenide, Solid Srate Phenom.. 182, 281. 1988.

18. Narayan, J. and Holland, 0. W., Characteristics of ion-implantation damage and annealing phenomena in semiconductors, J. Electrochem. SOC.. 13 I , 265 I , 1984.

19. Morgan, D. V. and Eisen, F. H., Ion implantation and damage in GaAs, in Gallium Arsenide, Howes, M. J. and Morgan, D. V., Eds., John Wiley & Sons, New York, 1985, 161.

120. Ryssel, H. and Ruge, I., Ion implantation, John Wiley & Sons, Chichester. 1986. 16ff, 229ff.

12 1. Poate, J. M. and Williams, J. S., Arnorphization and crystallization of semiconductors, in Ion Implantation and Beam Processing, Williams. J. S. and Poate, J. M., Eds.. Academic Press, Sydney, 1984. 14.

122. Feng, G. F., Zallen, R., Epp, J. M., Dillard, J. G., Raman-scattering and optical studies of argon-etched GaAs surfaces, Phys. Rev. B, 43, 9678. 1991.

123. Rubart, W. S., Jones, K. S., Seiberling, L., and Sandana, D. K., Low-energy implantation of Si and Sn into GaAs, Muter. Res. SOC. Symp. Proc., 157,677, 1990.

124. Pang, S. W., Geis, M. W., Efremow, N. N., and Lincoln, G. A., Effects of ion species and adsorbed gas on dry etching induced damage in GaAs, J. Vac. Sci. Technol.. B3. 398, 1985.

185

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

125. Pang, S. W., Goodhue, W. D., Lyszczarz, T. M., Ehrlich, D. J., Goodman, R. B.,and Johnson, G. D., Dry etching induced damage on vertical sidewalls of GaAs channels, J. Vuc. Sci. Technol.. B6. 1916, 1988.

126. Kawabe, M., Kanzaki, N., Masuda, K., and Samba, S., Effects of ion etching on the properties of GaAs. Appl. Opt., 17, 2556, 1978.

127. Bert, N. A., Pogrebitskii, K. Yu., Soshnikov, I. P., and Yur'ev, Yu. N., Sputtering of GaAs (001 ) by Ar+ ions with energies of 1-9 keV, Sov. Phys. Tech. Phj.s.. 37. 449, 1992.

128. Konomi, I. and Kawano, A., and Kido, Y.. Damage profiling of Ar+-irradiated Si( 100) and GaAs ( 100) by medium energy ion scattering, Surf. Sci.. 207. 427, 1989.

129. Sekino, Y., Owari, M., Kudo, M., and Nihei, Y., Estimation of low-energy ion bombardment damage on GaAs (001) surface by X-ray photoelectron diffraction, Jpn. J. Appl. Phys.. 25, 538, 1986.

130. Owari, M., Kudo, M., Nihei, Y., and Kamada, H., Estimation of surface crystal regularity by utilizing X-ray photoelectron diffraction (XPED) effects, Jpn. J. Appl. Phys., 19, 1203, 1980.

131. Jindal, S. K., Lin, S. S., Brown, L. D., IMarcus, H. L., and Schmerling, M. A., A study of low energy Ar+ ion bombardment induced lattice damage in ( 100) n-GaAs by channeling,Scan. Microsc.. 2, 1879. 1988.

132. Kang, H. J., Moon, Y. M., Kang, T. W., Leem, J. Y., Lee, J. J., and Ma, D. S., Surface composition and structure changes in GaAs compounds due to low-energy Ar+ ion bombardment. J. Vac. Sci. Technol.. A7. 3251, 1989.

133. Alberts, H. W., Caigher, H. L., and Friedland, E., Pulsed-electron-beam annealing of ion implanted GaAs, Muter. Sci. Eng., B9. 331, 1991.

134. Alberts, H. W., Gaigher, H. L., and Bredell, L. J., Channeling and TEM investigations of pulse electron beam annealed GaAs implanted with Pb, Nucl. Instrum. Methods, B80/8 I , 5 19, 1993.

135. Gaigher, H. L. and Alberts, H. W., TEM study of ion implanted GaAs after pulsed electron beam annealing, Radiat. Efl., 125. 373, 1993.

136. Vieu, C., Schneider, M., Launois, H., and Descouts, B., Damage generation and annealing in Ga' implanted GaAs/(Ga,At)As quantum wells, J. Appl. Phys., 71, 4833, 1992.

137. Williams, J. S., Jagadish, C., Clark, A., Li, G., and Larsen, C. A., Damage accumulation and amorphization in GaAs-APGaAs structures, Nucl. Instrum. Methods, B74, 80, 1993.

138. Desnica, U. V., Wagner, J., Haynes, T. E., and Holland, 0. W., Raman and ion channeling analysis of damage in ion-implanted GaAs: dependence on ion dose and dose rate, J. Appl. Phys.. 71. 2591, 1992.

139. Kido, Y. and Kawamoto, J., Projected range and damage distributions in ion-implanted A!, Si, A1,0,, and GaAs, J. Appl. Phys., 58, 3377, 1985.

140. Heyd, A. R., An, I., Collins, R. W., Cong, Y., Vedam, K., Bose, S. S., and Miller, D. L., Ion beam etching

of GaAslAPGaAs heterostructures probed in real time by spectroscopic ellipsometry, J. Vuc. Sci. Technol.. A9. 810, 1991.

I4 I . Oetzmann, H., Feuerstein, A., Crahmann, H., and Kalbitzer, S., Range parameters of heavy ions in amorphous targets at LSS-energies of 0.0006 I E I 0.3. P h y . Lett. A. 55. 170, 1975.

142. Kalbitzer, S. and Oetzmann, H., Ranges and range theories, Rudiat. Eff., 47. 57, 1980.

143. Schiott, H. E, Approximations and interpolation rules for ranges and range stragelings. Radiur. Efl., 6.107. 1970.

144. Tognetti, N. P., Elliman, R. C., and Carter, G., Disorder saturation in heavy ion implanted semiconductors, Vacurmi. 33. 165, 1983.

145. Winterbon, K. B., loti Implantation Runge and En- ergy Deposition Distributions. Vol. 2. Plenum Press, New York, 1975.

146. Kieslich, A., Straka. J., and Forchel, A., Investiga- tion of random and channeling Ar+ irnplantation-induced damage in AE( In)GaAs/GaAs quantum wells. Jpn. J. Appl. Phjs.. 3 I , 4428, 1992.

147. Larelle, F., Bagchi, A., Tsuchiya, M., Merz, J., and Petroff, P. M.. Focused ion beam channeling effects and ultimate sizes of GaAPAslGaAs nanostructures. Appl. Phys. Lett.. 56. 156 I , 1990.

148. Pearton, S. J., Poate, J. M., Sette, F., Gibson, J. M., Jacobson, D. C., and Williams, J. S., Ion implantation in GaAs. Nucl. Instruni. Methods. B 19R.O. 369, 1987.

149. Immorlica, A. A. and Eisen, F. H., Capless annealing of ion-implanted GaAs, Appl. PhTs. Lett.. 29. 94. 1976.

150. Hara, T.. Suzuki, H., Suga, A., Terada, T., and Toyoda, N.. Radiation damage of gallium arsenide induced by reactive ion etching, J. Appl. Phys.. 62, 4109. 1987.

151. Kanber, H., Feng, M., and Whelan. J. M., Charac- terization of ion implantation damage in capless annealed GaAs. Muter. Res. Soc. Symp. Proc.. 21. 365, 1984.

152. Wesch, W., Ion implantation in 111-V compounds, Nucl. Instrum. Methods. B68. 342, 1992.

153. Pearton, S. J.. Ion implantation doping and isolation of 111-V semiconductors, Nucl. Instrum. Methods. B59/60, 970, I99 1.

154. Wendler, E. and Wesch, W., Defect nucleation in weakly damaged ion implanted GaAs. Nucl. Instrum. Methods. 873. 489, 1993.

155. Kozanecki, A., Sealy, B. J., Cwilliam, R.,and Kidd, P., X-ray and channeling analysis of ion implanted gallium arsenide. Nucl. Instrum. Methods. B80/8 I , 798, 1993.

156. Cries, W. H., private communication, 1993. 157. Johnson, S. T., Elliman, R. G., and Williams, J. S.,

Ion beam induced epitaxy of (100) and ( I 1 I ) GaAs, Nucl. Instrum. Methods. B39, 449, 1989.

158. Komarov, F. F., Self-annealing in ion-implanted Si and GaAs, Vucuum. 42, 101, 1991.

159. Jones, K. S. and Santana, C. J., Arnorphization of elemental and compound semiconductors upon ion implantation. J. Marer. Res.. 6, 1048. 1991.

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

160. Poate, J. M. and Williams, J. S., Amorphization and crystallization of semiconductors, in Ion Implantation and Beam Processing, Williams, J. S . and Poate, J. M.. Eds.. Academic Press, Sydney, 1984. 13.

161. Hayes, T. R., Dry etching of In-based semiconductors, in Indium Phosphate and Related Materials: Processing Technology and Devices, Katz. A.. Ed., Artech House, Boston, I99 1, 277.

162. Callec, R., Poudoulec, A., Salvi, M., L'Haridon, H., Favennec, P. N., and Gauneau, M., Ion implantation damage and annealing in GaSb, Nucl. Instrum. Meth- ods, B80181. 532, 1993.

163. Callec, R. and Poudoulec, A., Characteristics of implantation-induced damage in GaSb, J. Appl. Phys.. 73,4831, 1993.

164. Gauneau, M., Chaplain, R., Rupert, A., Toudic, Y., Riviere, D., and Callec, R., Swelling of GaSb at low energies, Nucl. Instrum. Methods. B8018 I , 543, 1993.

165. Gauneau, M., Chaplain, R., Rupert, A., Toudic, Y., Callec, R., and AndrC, E., Secondary ion mass spectrometry generates swelling of GaSb: depth resolution and secondary ion yields, J. Appl. Phys.. 73, 205 I , 1993.

166. McHargue, C. J. and Williams, J. IM., Ion implantation effects in silicon carbide, Nucl. Instrum. Meth- ods. B80181, 889, 1993.

167. Ruttensperger, B., Krotz, G., Miiller, G., Derst, G., and Kalbitzer, S., Crystalline-amorphous contrast formation in thermally crystallized Sic, J. Non-Cryst. Solids. 1371138, 635, 1991.

168. Fohl, A., Emrick, R. M., and Carstanjen, H. D., A Rutherford backscattering study of Ar- and Xe-implanted silicon carbide, Nucl. Instrum. Methods. B65, 335. 1992.

169. Bentley, J., Romana, L. J., Horton, L. L., and McHargue, C. J., Distribution and characterization of iron in implanted silicon carbide, Mater. Res. SOC. Symp. Proc.. 235, 363, 1992.

170. Horton, L. L., Bentley, J., Romana, L., Perez, A., McHargue, C. J., and McCallum, J. C., Microstruc- tural characterization of iron ion implantation of silicon carbide, Nucl. Instrum. Methods. B65, 345, 1992.

17 1. McHargue, C. J. and Williams, J. M., Structure and properties of silicon carbide implanted with chromium, in Metastable Materials Formation by Iron Implanta- tion. Picraux, S . T. and Choyke, W. J., Eds., Elsevier, Amsterdam, 1982. 303.

172. McHargue, C. J., White, C. W., Appleton, B. R., Farlow, G. C., and Williams, J. M., Ion beam modification of ceramics, Mater. Res. Soc. Symp. Proc., 27, 385, 1984.

173. McHargue, C. J., Farlow, G. C., White, C. W., Williams, J. M., Appleton, B. R., and Naramoto, H., The amorphization of ceramics by ion beams, Mater. Sci. Eng.. 69, 123, 1985.

174. McHargue, C. J., Lewis, M. B., Williams, J. M., and Appleton, B. R., The reactivity of ion-implanted Sic, Mater. Sci. Eng., 69, 391. 1985.

175. McHargue, C. J., Farlow, G. C., Begun, G. M., Williams, J. M., White, C. W., Appleton, B. R., Sklad, P. S., and Angelini, P., Damage accumulation in ceramics during ion implantation, Nucl. Instrum. Methods, B 16. 2 12, 1986.

176. Nakata, T., Mizutani, Y., Mikoda, M., Watanabe, M., Takagi, T., and Nishino, S., Evaluation of At ion implanted 6H-Sic single crystals, Nucl. Instrum. Methods. B74, 13 I , 1993.

177. Spitznagel, J. A., Wood, S., Choyke, W. J., Doyle, N. J., Bradshaw, J., and Fishman, S. G., Ion beam modification of 6H/I 5R S i c crystals, Nucl. Instrum. Methods. B16, 237, 1986.

178. Tulinow, A. F., Chechenin, N. G., Bourdelle, K. K., Makarov , V. N., a n d Suvorov, A. V., S i c amorphization as a result of Ga+ implantation, Nucl. Instrum. Methods, B33. 788, 1988.

179. Williams, J. M., McHargue, C. J., and Appleton, B. R., Structural alterations in S i c as a result of Cr+ and N+ implantation, Nucl. Instrum. Methods, 2091 210, 317. 1983.

180. Bohn, H. G., Williams, J. M., McHargue, C. J., and Begun, G. M., Recrystallization of ion-implanted a-Sic , J . Mater. Res., 2, 107. 1987.

181. Chechenin, N. G., Bourdelle, K. K., and Suvorov, A. V., A channeling study of ion-produced disorder in silicon carbide, Nucl. Instrum. Methods, B48, 235, 1990.

182. Chechenin, N. G., Bourdelle, K. K., Suvorov, A. V., and Kastilio-Vitloch, A. X., Damage and aluminum distributions in S i c during ion implantation and annealing, Nucl. Instrum. Methods, B65, 341, 1992.

183. Pezoldt, J., Kalnin, A. A., Moskwina, D. R., and Savelyev, W. D., Polytype transitions in ion implanted silicon carbide, Nucl. Instrum. Methods, B80/81,943, 1993.

184. Pezoldt, J., Kalnin, A. A., and Savelyev, W. D., Application of nonequilibrium phase transition to heteropolytype structure creation, Nucl. Instrum. Meth- ods, B65, 361, 1992.

185. Du, H., Libera, M., Yang, Z., Lai, P., Jacobson, D. C., Wang, Y. C., and Davis, G. D., Effect of self- implantation on structure and oxidation behavior of single crystal p-Sic. Appl. Phys. Lett., 62,423, 1993.

186. Itoh, A., Hioki, T., Kawamoto, J., Tribological properties of ion implanted SIC ceramics, Nucl. Instrum. Methods. B37/38, 692, 1989.

187. Sigmon, T. W., Ion implantation in 11-VI compound semiconductors, Nucl. Instrum. Methods. B7/8, 402, 1985.

188. Kerkow, H., Quang, V. X., and Selle, B., Post-range implantation damage in XI-VI compounds due to elec- tronically stimulated diffusion, J . Cryst. Growth, 1 17, 677, 1992.

189. Friedland, E., Annealing of radiation damage in MgO single crystal after krypton implantation, Nucl. Instrum. Methods, B85, 316, 1994.

190. Bowman, R. C., Jr., Alt, R. L., Adams, P. M., Knudsen, J. F., Jamieson, D. N., and Downing,

187

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

R. G., Optical and structural characterization of heavily boron-implanted CdTe, J. Cryst. Growth. 86, 768, 1988.

191. Leo, G. and Ruault, M.-O., In situ transmission electron microscopy study of dislocation loops induced by P implantation in CdTe single crystals. J. Appl. Phys.. 73, 2234, 1993.

192. Wilson, R. G., Ion implantation and SIMS profiling of impurities in 11-VI materials (HgCdTe and CdTe), J. Cryst. Growth, 86, 735, 1988.

193. Leo, G., Drigo, A. V., and Traverse, A., Specific behaviour of CdTe ion implantation damage, Muter. Sci. Eng., B16, 123, 1993.

194. Leo, G., Traverse, A., Ruault, M. O., and Drigo, A. V., Ion channeling study of P implantation damage in CdTe, Nucl. Instrurn. & Methods, B63. 41. 1992.

195. Traverse, A., Leo, G., Alves, J. C., Almeida, P. M., Da Silva, M. F., and Soares, J. C., Disorder creation and annealing in Hg implanted CdTe, Nurl. Instrunt. Methods. B80/8l, 938. 1993.

196. Villegas, I. and Stickney, J. L., Formation of ordered, nonreconstructed CdTe ( I I I ) surfaces electro- chemically, J . Electrochem. Soc., 138, 1310. 1991.

197. Lu, Y.-C., Feigelson, R. S., and Route, R. K., A study of polar CdTe ( I 1 I ) surfaces using angle-resolved X-ray photoelectron and Auger electron spectroscopy and low-energy electron diffraction, J . Appl. Phys., 67,2583. 1990.

198. Chu, M., Bube, R. H., and Gibbons, J. F., Elec- tronic properties of As- and P-implanted cadmium telluride, J . Electrochem. Soc.. 127. 483, 1980.

199. Ueng, H. Y. and Hwang, H. L., Defects and doping effects in CdTe and CuInS, by phosphorus ion implantation and pulsed electron beam annealing. Solid State Phenom.. In. 343, 1988.

200. Bubulac, L. O., Defects, diffusion and activation in ion implanted HgCdTe, J. Cryst. Growth. 86, 723, 1988.

201. DestCfanis, G. L., Electrical doping of HgCdTe by ion implantation and heat treatment. 1. Ci-yst. Growth. 86,700, 1988.

202. DestCfanis, G. L., Ion implantation in Hg,-,Cd,Te, Nucl. Instrurn. Methods 209/210, 567. 1983.

203. Bahir, G. and Finkman, E., Ion beam milling effect on electrical properties of Hg,-,Cd,Te. J . Vuc. Sci. Technol.. A7, 348, 1989.

204. Bubulac, L. O., Tennant, W. E., Riedel, R. A., and Magee, T. J., Behavior of implantation-induced defects in HgCdTe,J. Vuc. Sci. Technol., 21,25 1. 1982.

205. Carey, G. P., Cole, S., Yamashita, T., Silberman, J. A., Spicer, W. E., and Wilson, J. A., TEM investigation of the differences in ion milling induced damage of Hg,-,Cd,Te and CdTe heterojunctions, J. Vuc. Sci. Technol., A3,255, 1985.

206. Schaake, H. F., Ion implantation damage in Hg,,,Cd,,,Te, J . Vuc. Sci. Technol., A4, 2174, 1986.

207. Bahir, G. and Kalish, R., Structure of ion-implanted and annealed Hg,_,Cd,Te, J . Appl. Phys., 54, 3129. 1983.

208. Bahir, G., Bernstein, T., and Kalish, R., Damage and lattice location studies in Hg implanted Hg,-,Cd,Te. Rudiut. Eff., 48, 247. 1980.

209. Uzan-Saguy, C., Richter, V., Shaanan, M., Kalish, R., and Tribolet, R., Comparison of ion-implantation- induced damage in narrow-gap (0.1 eV) Hg,Cd,-,Te and Hg,Zn,-,Te,J. Vuc. Sci. Technol..AlO, 3246.1992.

210. Baars, J., Hurrle, A., Rothemund, W., Fritzsche, C. R., and Jakobus, T., Boron ion implantation in Hg,-,Cd,Te. J. Appl. P h y . , 53. 1461, 1982.

2 I 1 . Ivanov-Omskii, V. I., Mironov, K. E., and Mynbaev, K. D., HgCdTe doping by ion-beam treatment, Semicond. Sci. Technol.. 8, 634, 1993.

212. Bubulac, L. O., Lo, D. S., Tennant, W. E., Edwall. D. D., and Robinson, J. C., Boron and indium ion- implanted junctions in HgCdTe grown on CdTe and CdTe/Al,O,, J. Vuc. Sci. Technol.. A4. 2 169, 1986.

213. DestCfanis, G. L., Indium ion implantation i n Hg,,,,Cd,, ,,Te/CdTe. J . Vuc. Sci. Techno/.. A3, I7 1, 1985.

214. Friedland, E., le Roux, G., and Malherbe, J. B., Deep radiation damage in copper after ion implantation, Rudiur. Efl. Lerr.. 87, 281, 1986.

215. Friedland. E., .Malherbe, J. B., Alberts, H. W., Vorster, R. E., and Prins, J. F., Damage profiles in copper single crystals after ion implantation, S. Afr. J . Phys.. 9. 135. 1986.

216. Friedland, E., Alberts, H. W., and Fletcher, M., Temperature dependence of damage ranges in some metals after argon implantation, NM/. Instrum. MerA- ods. B45, 492. 1990.

217. Friedland, E. and Alberts, H. W., Radiation damage in nickel and iron after ion implantation. Nucl. Instrum. Methods, B35, 244. 1988.

218. Friedland, E. and Fletcher, M., Structure dependence of radiation damage depths after ion implantation. Nucl. l/istrirm. Methods. B64, 242. 1992.

2 19. Elkind, J. L., Ion mill damage in n-HgCdTe, J . Vuc. Sci. Technol.. BIO, 1460, 1992.

220. Blanchard, C., Barbot, J. F., Cahoreau, M., Desoyer, J. C., le Scoul, D., and Dessus, J. L., Study of A1 ion implantation in Hh,Cd,,Te. Nucl. /mtrirm. Methods, B47. 15. 1990.

22 I . Uzan, C., Marfaing, Y., Legros, R., Kalish, R.. and Richter, V., Damage study and physical properties of ion implanted C4,,Hg,, ,Te processed by furnace and rapid thermal annealing, J. Cryst. Growth. 86, 744, 1988.

222. Dimiduk, K. C., Opyd, W. G., Gibbons, J. F., Sigmon, T. W., Magee, T. J., and Ormond, R. D., Annealing of Hg,-,Cd,Te: Hg loss rates and annealing of ion implantation damage, J. Vuc. Sci. Technol.. A I , 1661, 1983.

223. Gerasimenko, N. N., Nesterov, A. A., and Vasil’ev, V. V., Ion implantation and rapid thermal annealing of Au-Cd,Hg,-,Te structures, Nucl. Instrurn. Meth- ods. B80/81, 919, 1993.

224. Uzan, C., Legros, R., and Marfaing, Y., A compara- tive study of laser and furnace annealing of P+ implanted CdTe, J. Cryst. Growth, 72, 252, 1985.

188

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

225. Magel, L. K. and Sigmon, T. W., Dose rate effects in indium-implanted Hg,_,Cd,Te, J. Cryst. Growth, 86, 756, 1988.

226. Baars, J., Seelewind, H., Fritzsche, Ch., Kaiser, U., and Ziegler, J. , Arsenic ion implantation i n Hg,_,Cd,Te, J. Cryst. Growth. 86, 762. 1988.

227. Wu, S. Y., Choyke, W. J., Takei, W. J., Noreika, A. J., Francombe, M. H., and Irwin, R. B., Effects of implantation and annealing temperatures on implantation induced damage in HgCdTe. J. Vac. Sci. Technol., 21. 255, 1982.

228. Ghenim, L., Robert, J. L., Bousquet,C., Raymond, A., and Destefanis, G. L., Evidence, by transport measurements under hydrostatic pressure, of resonant level induced by implantation defects in epitaxial layers of Hg,-,Cd,Te. J. Cryst. Growth, 72.448, 1985.

229. Bubulac, L. O., Tennant, W. E., Riedel, R. A., Bajaj, J., and Edwall, D. D., Some aspects of Li behavior in ion implanted HgCdTe. J. Vac. Sci. Technol.. A l , 1646, 1983.

230. Bubulac, L. O., The role of epitaxy and substrate on junction formation in ion-implanted HgCdTe, J. Cryst. Growth. 72, 478, 1985.

231. Cotton, V. A. and Wilson, J. A., Effects of ion implantation on deep electron traps in He,,Cd,,Te, J. Vac. Sci. Technol., A4, 2177, 1986.

232. Shastov, K. V., Depth distribution of radiation defects in implanted Hg,Cd,_,Te crystals, Phys. Status Solidi A. 130, 293, 1992.

233. Belas, E.. Hoschl, P., Grill, R., Franc, J., Moravec, P., Lischka, K., Sitter, H., and Toth, A., Deep p-n junction in Hg,Cd,_,Te created by ion milling, Semicotid. Sci. Techno/.. 8, 1695. 1993.

234. Bahir, G., Kalish, R., and Nemirovsky, Y., Electri- cal properties of donor and acceptor implanted Hg,_,Cd,Te following cw COz laser annealing. Appl. Phgs. Lett.. 41. 1057, 1982.

235. Kalish, R. and Bahir, G., Formation of p on n pho- todiodes in Hg,_,Cd,Te by ion implantation and cw CO? laser annealing, J . Ctyst. Growth, 72.474, 1985.

236. Kao, T. M. and Sigmon, T. W., B+ implantation and postimplant anneal studies in Hg,-,Cd,Te. Nucl. Instrum. Methods. B2 1, 578, 1987.

237. Conway, K. L., Opyd, W. G., Greiner, M. E., Gibbons, J. F., Sigmon, T. W., and Bubulac, L. O., Thermal pulse annealing of boron-implanted HgCdTe, Appl. Phvs. Lett., 41, 750, 1982.

238. Syz, J., Beck, J. D., Orent, T. W., and Schaake, H. F., Thermal annealing studies on boron-implanted HgCdTe diodes, J. Vac. Sci. Technol.. A7.396, 1989.

239. Betz, G. and Wehner, G. K., Sputtering of multicomponent materials. in Sputtering by Particle Bom- bardment 11, Behrisch, R.. Ed., Springer-Verlag. Ber- lin, 1984. 1 I .

240. Kelly, R. and Harrison, D. E., A summary of the theory of the preferential sputtering of alloys, Mater. Sci. Eng., 69, 449, 1985.

241. Kelly, R. and Oliva, A., The theory of the preferential sputtering of alloys, including the role of Gibbsian segregation, in Erosion and Growth of Solids Stimu-

lated by Atom and Ion Beams, Kiriakidis, G.. Carter, G.. and Whitton. J. L., Eds., Martinus Nijhoff, Dordrecht, 1986, 41.

242. Lam, N. Q., Physical sputtering of metallic systems by charged-particle impact, Scan. Microsc. Suppl., 4. 311, 1990.

243. Sigmund, P., Alloy and isotope sputtering, Mat. Fiz. Medd. Dan. Vid. Selsk, I . 1993.

244. Kelly, R., Surface compositional changes by particle bombardment. in Chemistry and Physics of Solid Sur- faces V , Vanselow, R. and Howe, R., Eds., Springer- Verlag. Berlin, 1984. chap. 7.

245. Andersen, H. H., Ion-bombardment-induced composition changes in alloys and compounds, in Ion Im- plantation and Beam Processing, Williams. J. S . and Poate. J. M., Eds., Academic Press, Sydney, 1984, chap. 6.

246. Malherbe, J. B., Hofmann, S., and Sanz, J. M., Preferential sputtering of oxides: a comparison of model predictions with experimental data, Appl. Sutf. Sci., 27, 355, 1986.

247. Haff, P. K., A model for surface layer composition changes in sputtered alloys and compounds, Appl. Phys. Lett., 3 I , 259. 1977.

248. Sigmund, P., Sputtering by ion bombardment: theoretical concepts, in Sputtering by Particle Bombard- ment I , Behrisch, R., Ed., Springer-Verlag. Berlin, 1981. chap. 2.

249. Kelly, R., An attempt to understand preferential sputtering, Nucl. Instrum. Methods, 149. 553, 1978.

250. Kelly, R., On the problem of whether mass or chemical bonding is more important to bombardment-induced compositional changes in alloys and oxides, SurJ Sci.. 100. 85, 1980.

251. Pauling, L., The Nature of the Chemical Bond, 3rd ed., Cornell University Press, Ithaca. NY, 1960.

252. Brongersma, H. H., Spaarnay, M. J., and Buck, T., Surface segregation in Cu-Ni and Cu-PI alloys: a comparison of low-energy ion-scattering results with theory, Sutf. Sci., 71. 657, 1978.

253. Williams, F. L. and Nason, D., Binary alloy surface compositions from bulk thermodynamic data, Sutf. Sci., 45. 377, 1975.

254. Johnson, R. A. and Lam, N. Q., Solute segregation under irradiation, J. Nucl. Muter., 69/70, 424, 1978.

255. Lam, N. Q., Leaf, G. K., and Wiedersich, H., Sput- ter-induced surface composition changes in alloys, J. Nucl. Mater., 88, 289, 1980.

256. Lam, N. Q., Okamoto, P. R., and Wiedersich, H., Effects of solute segregation and precipitation on void swelling in irradiated al1oys.J. Nucl. Mater.. 74. 101. 1978.

257. du Plessis, J., Surface segregation, Diffus. Defect Data Part B , 11. 1, 1990.

258. du Plessis, J., van Wyk, G. N., and Taglauer, E., heferential sputtering and radiation enhanced segregation in palladium-platinum, Sutf. Sci., 220. 381, 1989.

259. Lam, N. Q. and Wiedersich, H., Dynamical behavior of the subsurface region in alloys under ion bom-

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

bardment at high temperatures, in Metastable Materi- als Formation by Ion Implantation. Picraux, S. T. and Choyke. W. J., Eds.. Elsevier, 1982, 35.

260. Rehn, L. E., Surface modification and radiation-induced segregation, in Metastable Materials Forma- tion by Ion Implantation, Picraux, S. T. and Choyke. W. J., Eds., Elsevier, 1982, 17.

261. Marwick, A. D., Piller, R. C., and Sivell, P. bl., Mechanisms of radiation-induced segregation in di- lute nickel alloys, J. Nucl. Mater.. 83. 35. 1979.

262. Lam, N. Q. and Leaf, C. K., Ion implantation at elevated temperatures, Mater. Res. Soc. Symp. P roc., 51. 467, 1985.

263. Yacout, A. M., Lam, N. Q., and Stubbins, J. F., Radiation-induced precipitation at the alloy surface during ion bombardment, Nucl. Instrum. Methods. B71, 148, 1992.

264. Wynblatt, P. and Ku, R. C., Surface segregation in alloys, in Interfacial Segregation. Johnson, W. C. and Blakeley, J. M., Eds., American Society of Metals, Metals Park, OH, 1979, I 15.

265. Schulz, F. and Wittmaack, K., Model calculation of ion collection in the presence of sputtering, Rudiat. Eff., 29, 3 I , 1976.

266. Malherbe, J. B. and Hofmann, S., Low energy nitrogen implantation profiles in cobalt using AES. Surf. Interface Anal., 2, 187. 1980.

267. Malherbe, J. B., Projected range and straggling rneas- urements of low energy nitrogen in silicon, Radiat. E#., 70, 261, 1983.

268. Malherbe, J. B., Implantation parameters of low energy nitrogen in copper, Nucl. Instrum. Metliods. B1, 774. 1984.

269. Wittmaack, K. and Blank, P., The influence of recoil implantation of absorbed oxygen on the entrap- ment of xenon in aluminum and silicon. in Ion Im- plantation in Semiconductors, 1976. Chemov. F.. Borders, J. A.. and Brice, D. K., Eds.. Plenum Press, New York. 1977, 363.

270. Wittmaack, K. and Blank, P., Radiation-enhanced outdiffusion of xenon implanted in aluminum. Appl. Phys. Lett.. 31, 21, 1977.

271. Wittmaack, K., Blank, P., and Wach, W., High fluence retention of noble gases implanted in silicon, Radiat. Ef., 39. 81, 1978.

272. Blank, P., Wittmaack, K., and Schulz, F., The influence of sputtering. range shortening and stress- induced outdiffusion on the retention of xenon implanted in silicon. Nucl. Instrum. Methods, 132.387, 1976.

273. Menzel, M. and Wittmaack, K., Comparison of the retention characteristics of low-energy xenon and cae- sium implanted in silicon, Nucl. Instrum. Methods. 191, 235, 1981.

274. Wach, W. and Wittmaack, K., Comparison of xenon retention in ion implanted silicon dioxide and oxygen-doped silicon, Nucl. Instrum. Methods. 149, 259, 1978.

275. Wach, W. and Wittmaack, K., Implantation and sputtering in the presence of reactive gases: bombard-

ment-induced incorporation of oxygen and related phenomena, J. Appl. Phys.. 52, 3341. 1981.

276. Kempf, J., AES investigations of Ar' ion retention in Si during Ar sputtering, Appl. Phys., 16, 43, 1978.

277. Kirschner, J. and Etzkorn, H. W., Sputtering of amorphous silicon films by 0.5 to 5 keV Ar' ions, Appl. Suif. Sci., 3. 251. 1979.

278. Carter, G., Armour, D. G., Donnelly, S. E., Ingram, D. C., and Webb, R. P., The injection of inert gas ions into solids: their trapping and escape, Radiat. E8. . 53, 143. 1980.

279. Cries, W. H., Quantitative ion implantation: theoretical aspects, Int. J. Mass Spectrom. Ion Phys.. 30.97, 1979.

280. Singer, 1. L., Murday, J. S., and Cooper, L. R., Surface composition changes in GaAs due to low- energy ion bombardment, Surf. Sci.. 108, 7, 198 I .

281. Tsang, J. C., Kahn, A., and Mark, P., Comparison of LEED and Auger data from cleaved and sputtered- annealed InP ( I 10) surfaces, Surf. Sci.. 97, I 19, 1980.

282. Barnard, W. 0.. Malherbe, J. B., and Myburg, G., AES study of InP surfaces subjected tochemical cleaning and Ar+ sputtering. s. Afi. J. PhFs.. 14, 22. 1991.

283. Malherbe, J. B. and van der Berg, N. G., Surface compositional changes of InP due to krypton ion bombardment, Surf. Interface Anal., in press.

284. Mayer, T. M. and Barker, R. A., Simulation of plasma-assisted etching processes by ion-beam techniques, J. Vac. Sci. Techno/.. 21, 757. 1982.

285. van Roijen, R., Kemp, M. B. M., Bulle-Lieuwma, C. W. T., van Ijzendoorn, L. J., and Thijssen, T. L. G., Surface analysis of reactive ion-etched InP. J. Appl. Phys., 70, 3983. 1991.

286. Gallet, D., Hollinger, G., and Santinelli. C., I n sitit

characterization of InP surfaces after low-energy hydrogen ion cleaning. J. Vuc. Sci. Teclv?ol.. B 10. 1267. 1992.

287. Olego, D. J., Schachter, R., Viscogliosi, M.. and Bunz, L. A., Optoelectronic and structural properties of sputter etched surfaces of InP, Appl. Phys. Lett.. 49. 719, 1986.

288. Dubreuil, B., Gibert,T., Barthe, M. F., and Debrun, J. L., RIMS study of low energy-laser sputtering of metal and semiconductor surfaces, Inst. Phys. Con& Ser.. 128. 265. 1992.

289. Bayliss, C. R. and Kirk, D. L., The effect of temperature on the surface structure and stoichiometry of (100) InP surfaces, Thin Solid Films, 29, 97, 1975.

290. Jones, C. J. and Kirk, D. L., Some considerations of the limitations of the Auger-depth-profiling technique as applied to silver metal and (100) surfaces of indium phosphide. J. Phys. D, 12, 837. 1979.

291. Stair, K. A. and Chung, Y.-W., Fermi-level pinning and growth characteristics of Au on InP(iii). Appl. Surf. Sci., 26, 381, 1986.

292. Clark, D. T., Fok, T., Roberts, G. C., and Sykes, R. W., An investigation by electron spectroscopy for chemical analysis of chemical treatments of the (100) surface of n-type InP epitaxial layers for Langmuir film deposition, Thin Solid Films, 70. 26 I , 1980.

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

293. Barcz, A., Croset, M., and Mercandalli, M. L., Quantitative analysis of Ill-V compounds, using low- energy ion-scattering spectrometry (ISS). Inst. Phys. Conf. Ser.. 54, 124. 1980.

294. Barcz, A., Croset, M., and Mercandalli, M. L., Quantitativity in 111-V compounds by low-energy ion scattering spectrometry (ISS), Surf. Sci., 95. 51 I . 1980.

295. Williams, R. S., Low energy Ar ion bombardment damage of Si, GaAs, and InP surfaces, Solid Sfare Comrnun., 41, 153, 1982.

296. Tu, C. W. and Schlier, A. R., Surface studies of InP by electron energy loss spectroscopy and Auger electron spectroscopy, Appl. Surf. Sci., 11/12, 355, 1982.

297. Shapira, Y., Brillson, L. J., and Heller, A., Investi- gation of InP surface and metal interfaces by surface photovoltage and Auger electron spectroscopies, J . Vac. Sci. Technol., A I , 766, 1983.

298. Gruzza, B., Achard, B., and Pariset, C., Surface composition of (1 0O)InP substrates bombarded by low energy Ar+ ions, studied by AES and EPES. Suf. Sci.. 162,202, 1985.

299. Gruzza, B., Pariset, C., and Abdellaoui, S., An ELS and AES study of InP( 100) surfaces during Sb atom condensation, Suif . Inreiface Anal., 16, 54, 1990.

300. Gruzza, B. and Pariset, C., Some aspects of AES, EELS and EPES application. Surf. Sci.. 247, 408, 1991.

301. Gruzza, B., Porte, A., Bideux, L., Jardin, C., and Miloua, J., Models for the interpretation of the different interfaces Sb/InP( 100) using quantitative results of AES and EELS, Acto Phys. Pol. A , 81,223. 1992.

302. Achard, B., Cruzza, B., and Pariset, C., Effect of ion bombardment at low energy on (100)InP surfaces, studied by Auger electron spectroscopy, S w f Sci.. 160, L519, 1985.

303. Jin, X., Yu, M., and Wang, X., The elimination of In islands on InP surfaces and its mechanism, Chin. Phys. Lett.. 2. 345. 1985.

304. Lau, W. M., Sodhi, R. N. S., Flinn, B. J., Tan, K. H., and Bancroft, G. M., Photoemission study of sputter-etched InP surfaces, Appl. Phys. Lett.. 5 1, 177, 1987.

305. Sodhi, R. N. S., Lau, W. M., and Ingrey, S. I. J., Core-level binding energy shifts on InP surfaces induced by sputtering and annealing, Surf. Interface Anal., 12, 321, 1988.

306. Wang, X., Investigation of atomic and electronic smc- tures of InP polar surfaces, Appl. Suf. Sci., 33/34,88, 1988.

307. Abdellaoui, S., Gruzza, B., Pariset, C., Bouslarna, M., Jardin, C., and Robert, D., Study of Sb condensation on InP( 100) substrates previously cleaned by low energy Ar' ion beam, Surf. Sci., 208, L21, 1989.

308. Pearton, S. J., Chakrabarti, U. K., Perley, A. P., and Jones, K. S., Ion milling damage in InP and GaAs, J. Appl. Phys., 68, 2760. 1990.

309. Schaefer, J. A., Frankel, D. J., and Lapeyre, G. J., Active sites of adsorption on cleaved and sputtered indium phosphide surfaces, Z. Phys. B . 79,259.1990.

3 10. Valeri, S. and Lolli, M., AES, EELS and XPS study of ion-induced GaAs and InP (1 10) surface and subsurface modifications, Surf. Interface Anal., 16. 59, 1990.

3 1 1. Valeri, S., Lolli, M., and Sberveglieri, P., AES and EELS study of alkali-metal adsorption kinetics on either cleaved or sputtered GaAs and InP ( I 10) surfaces, Surf. Sci.. 238, 63, 1990.

312. Malherbe, J. B. and Barnard, W. O., Preferential sputtering of I n P an AES investigation. Sutf. Sci., 255. 309, 1991.

3 13. Malherbe, J. B., Ion angle dependence in the preferential sputtering of InP, Appl. Sutf. Sci., 70/71, 322, 1993.

314. Morais, J., Fazan, T. A., and Landers, R., Ion induced composition changes on (1 10) InP cleaved surface, Phys. Status Solidi, 141. K19, 1994.

315. Nelson, G. C., Search for preferential sputtering in Ag/Au alloys, J. Vac. Sci. Technol., 13, 974, 1976.

3 16. Holloway, P. H. and Bhattacharya, R. S., Preferen- tial sputtering of RSi. Nisi, and AgAu, J . Vac. Sci. Technol.. 20, 444, 1982.

317. Hofmann, S. and Zalar, A., Electron beam effects during the sputter profiling of thin Au/Ag films analysed by Auger electron spectroscopy, Thin Solid Films, 56, 337, 1979.

318. Santorius, B., Schlak, M., Rozenzweig, M., and Parschke, K., Thermal degradation effects in InP, J . Appl. Phys.. 63, 4677, 1988.

319. Hofmann, S., Charging and charge compensation in AES analysis of insulators, J. Electron. Spectrosc. Relat. Phenom., 59. 15, 1992.

320. Kernpf, J.,Nonnenmacher, M.,and Wagner,H. H., Electron and ion beam induced heating effects in solids measured by laser interferometry, Appl. Phys., A56. 385, 1993.

321. Klekarnp, A., Snowdon, K. J., and Heiland, W., Real time study of the change in surface composition of insulators under electron bombardment, Radial. Ef., 108. 241, 1989.

322. Wiedmann, L., Ganschow, O., and Benninghoven, A., Contamination of clean metal surfaces associated with electron bombardment in conventional AES analysis, J. Electron. Spectrosc. Relat. Phenom.. 13, 243, 1978.

323. Roland, C., Baumann, H., and Bethge, K., RBS analysis of GaAs and InP after electron beam annealing, Nucl. Instrum. Methods, B50. 145, 1990.

324. di Bona, A., Facchini, A., Valeri, S., Ottaviani, G., and Piccirillo, A., Structural and compositional modifications of 111-V ternary and quaternary compounds induced by ion bombardment, Mater. Res. SOC. Symp. Proc., 223, 197, 1991.

325. Valeri, S., di Bona, A., and Nava, E., Diffraction effects in Auger quantitative analysis on III-V compounds, Appl. Sutf. Sci., 70P1, 20, 1993.

326. Valeri, S. and di Bona, A., Auger electron spectroscopy for structural changes, Riv. Nuovo Cimento, 16, 1. 1993.

191

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

327. Malherbe, J. B., Barnard, W. O.,Strydom, 1. Le R., and Louw, C. W., Preferential sputtering of GaAs, Surf. Interface Anal., 18, 491, 1992.

328. Taglauer, E., Surface modifications due to preferential sputtering, Appl. Surf. Sci., 13, 80, 1982.

329. Baretzky, B. and Taglauer, E., Does preferential sputtering depend on the angle of incidence?, Slf$.

Sci.. 162, 996, 1985. 330. Baretzky, B., Moiler, W., and Taglauer, E., Colli-

sional processes in preferential sputtering of tantalum oxide, Nucl. Instrum. Methods. B18. 496. 1987.

331. van Wyk, G. N., du Plessis, J., Bangwei, Z., and Taglauer, E., Preferential sputtering in amorphous CuTi alloys. S. Afr. J. Phys., 12, 58c, 1989.

332. Baretzky, B., Moller, W., and Taglauer, E., Colli- sion dominated preferential sputtering of tantalum oxide, Vacuum, 43. 1207, 1992.

333. Nagai, Y., Nishimura, C., and Toshima, T., Primary Ar' ion bombardment effect on Ni-Fe film composition formed by ion beam sputtering, J. Vac. Sci. Technol.. A3, 2147, 1985.

334. Corvini, P., Kahn, A., and Wagner, S., Surface order and stoichiometry of sputter-cleaned and annealed CuInSe,, J . Appl. Phys., 57. 2967, 1985.

335. Moses, A. J., The Practicing Scientist's Handbook. Van Nostrand Reinhold, New York. 1978, chap. 9.

336. Biersack, J. P., Calculation of projected ranges - analytical solutions and a simple general algorithm, Nucl. Instrum. Methods, 182/183, 199, 1981.

337. Ziegler, J. F., TRIM: The Transport of Ions in Matter. IBM Res. 28-0, Yorktown. NY.

338. For this calculation, the average composition for the top five layers (=AES composition) was taken to represent the composition of the sputtered samples of References 298 to 30 I , viz., InPo,82.

339. Kerr, J. A., Strength of chemical bonds, in Handbook of Chemistry and Physics, Lide, D. R., Ed., CRC Press, Boca Raton. FL. 1992. Table 2. 9-135.

340. Allred, A. L., McGraw-Hill Encyclopedia of Science and Technology. Vol. 6, 6th ed., McGraw-Hill, New York, 1987, 210.

341. Andersen, H. H., Computer simulations of atomic collisions in solids with special emphasis on sputtering, Nucl. Instrum. Methods, B18, 321, 1987.

342. Sakalas, A. and Zhukauskas, S., GaAs surface low- dose irradiation pretreatment related reduction of Schottky barrier height. Real or effective?, Solid Stare Commun., 70. 363. 1989.

343. van Oostrom, A., Application of AES to the study of selective sputtering of thin films. J. Vac. Sci. Technol.. 13. 224, 1976.

344. Valeri, S. and di Bona, A., Alkali metals adsorption kinetics on sputtered and cleaved GaAs ( 1 10) surfaces, Surf. Sci., 251/252, 995. 1991.

345. Wang, Y. X. and Holloway, P. H., Summary abstract: effect.of ion sputtering on the interface chemistry and electrical properties of Au/n-GaAs (100) Schottky contacts, J. Vac. Sci. Techno/.. A2.567,1984.

346. Wang, Y. X. and Holloway, P. H., Effect of ion sputtering on interface chemistry and electrical prop-

erties of Au/n-GaAs (100) Schottky contacts, J. Vac. Sci. Technol., B2, 613. 1984.

347. McGuire, G. E., Effects of ion sputtering on semiconductor surfaces, Sug. Sci.. 76. 130, 1978.

348. Bussing, T. D., Holloway, P. H., Wang, Y. X., Moulder, J. F., and Hammond, J. S., Composition depth profiles of oxidized silicon and sputtered GaAs from angle-resolved X-ray photoelectron spectroscopy, J. Vac. Sci. Technol.. B6, 1514, 1988.

349. Barcz, A., Bugajski, M., Croset, M., and Mercan- dalli, L. M., Stoichiometry changes in 111-V compounds under ion bombardment, Nucl. Instrum. Meth- ods. 209/210.621. 1983.

350. Jacobi, K. and Ranke, W., Oxidation and annealing of GaP and GaAs ( 1 1 I ) - faces studied by AES and UPS, J. Electron Spectrosc. Relat. Phenom.. 8. 225, 1976.

351. Ludeke, R., Esaki,L.,and Chang, L. L., Ga,_,AI,As superlattices profiled by Auger electron spectroscopy, Appl. Php. Lett.. 24. 417, 1974.

352. Mizokawa, Y., Iwasaki, H., Nishitani, R., and Nakamura, S., ESCA studies of Ga. As. GaAs. Ga,O,. As,O, and As@,. J. Electron Spectrosc. Relar. Phenoni.. 14, 129. 1978.

353. Vosylius, J. J., Pranevicius, L. J., Puras, R. D., and Sakalas, A. P., Analysis of the element composition on the surface of heated GaAs under ion bombardment, Litov. f i: . Sh.. 23. 66, 1983.

354. Chang, C. C., Citrin, P. H., and Schwartz, B., Chemical preparation of GaAs surfaces and their characterization by Auger electron and X-ray photoemission spectroscopies. J. Vac. Sci. Techitol.. 14. 943. 1977.

355. Schwartz, G. P., Gualtieri, G. J., Kammlott, G. W., and Schwartz, B., An X-ray photoelectron spectroscopy study of native oxides on GaAs. J. Electrochem. Soc.. 126. 1737, 1979.

356. Holloway, P. H., Progress in instrumentation, data reduction. and depth profiles in Auger electron spectroscopy. Appl. S w f . Sci.. 26. 550, 1986.

357. Rabinzohn, P., Gautherin, G., Agius, B., and Cohen, C., Cleaning of Si and GaAs crystal surfaces by ion bombardment in the 50-1500 eV range, J. Electrochem. Soc.. 13 I , 905, 1984.

358. Feng, G. F., Holtz, M., Zallen, R., Epp, J. M, Dillard, J. G., Cole, E., Johnson, P., Sen, S., and Burton, L. C., Optical. chemical and electrical characterization of ion-etched gallium arsenide surfaces, Mate).. Res. SOC. Symp. Proc.. 93. 381, 1989.

359. Orrman-Rossiter, K. C., Al-Bayati, A. H., and Armour, D. G., Low energy ion-scattering of the GaAs (001) surface, Surf. Sci.. 225, 341. 1990.

360. Valeri, S., Lancellotti, M. G., di Bona, A., Granozzi, G., and Rizzi, G. A., Crystalline effects on Auger and photoelectron emission from clean and Cs-covered GaAs (110) surfaces, Appl. Surf. Sci.. 56. 205, 1992.

36 1. Valeri, S. and di Bona, A., Alkali metals adsorption kinetics on sputtered and cleaved GaAs (1 10) surfaces, Surf. Sci.. 251. 995, 1991.

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

362. DeLouise, L. A., Investigation of the mechanism of Ar+ ion-assisted CI, etching of GaAs I 110): role of ion-induced charge acceptor states, J. Appl. Phys., 70, 1718, 1991.

363. Vaseashta, A., Elshabini-Riad, A., and Burton, L. C., Effects of low energy Ar‘ ion bombardment on GaAs, Muter. Sci. Eng.. B9. 489, 1991.

364. Wright, S. L., Marks, R. F., and Savoy, R. J., I n situ measurement of the composition of molecular-beam epitaxial (At , Ga)As by Auger electron spectroscopy, J. Vuc. Sci. Techno/., B6. 1105, 1988.

365. Epp, J. M. and Dillard, J. G., Effect of ion bombardment on the chemical reactivity of gallium arsenide (100). Chem. Muter., 1, 325, 1989.

366. Cossu, G., Ingo, G. M., Mattogno, C.,Padeletti, G., and Viticoli, S., Auger sputtering profiling of an AI,,,Ga,, ,As/GaAs superlattice grown by molecular beam epitaxy, Appl. Suq. Sci., 56-58, 708, 1992.

367. DeLouise, L. A., Reactive Nf ion bombardment of GaAsl 110): a method for GaN thin film growth. J. Vuc. Sci. Techno/., AIO. 1637. 1992.

368. DeLouise, L. A., Nitridation of GaAs ( 1 10) using energetic N+ and N; ion beams, J. Vuc. Sci. Techno/.. A l I . 609, 1993.

369. Singer, I. L., Murday, J. S., and Cooper, L. R., Abstract: composition changes in GaAs due to low- energy ion bombardment, J. Vuc. Sci. Techno/.. 15, 725. 1978.

370. Singer, 1. L., Murday, J. S., and Comas, J., Prefer- ential sputtering from disordered GaAs. J. Vuc. Sci. Techno/.. 18, 161, 1981.

371. Basol, B. M., Ou, S. S., and Stafsudd, 0. M., Type conversion, contacts, and surface effects in electro- plated CdTe films, J. Appl. Phys., 58, 3809, 1985.

372. Anthony, T. C., Fahrenbruch, A. L., and Bube, R. H., Low resistance contacts to p-type cadmium telluride, J. Electron. Muter.. 11 , 89, 1982.

373. Kahn, A., Semiconductor surface structures. Sir$

Sci. Rep., 3, 193, 1983. 374. Kahn, A., Atomic geometries of compound semicon-

ductor surfaces and interfaces, J . Vuc. Sci. Techno/.. A l , 684, 1983.

375. Duke, C. B., Paton, A., and Kahn, A.. Atomic ge- ometry of GaSb ( I 10): determination via elastic low- energy electron diffraction intensity analysis. Phys. Rev. B. 27, 3436, 1983.

376. Kiibler, B., Ranke, W., and Jacobi, K., LEED and AES of stoichiometric and arsenic-rich GaAs ( 1 10) surfaces prepared by molecular beam epitaxy, Si41-5 Sci., 92. 519, 1980.

377. Miller, D. L., Preparation and characterization of ultrathin Pb and In films on PbTe and Te, J. Vuc. Sci. Techno/.. 13, 1138, 1976.

378. Schweitzer, M. O., Leibsle, F. M., Jones, T. S., McConville, C. F., and Richardson, N. V., An STM study of the InSb (100)-c(8 x 2) surface, Surf. Sci., 280, 63, 1993.

379. Chang, S.-C. and Mark, P., LEED analysis of the principal nonpolar ( I 120) and (lOi0) surfaces of cadmium sulfide, J . Vuc. Sci. Techno/., 12. 624, 1975.

380. Kohiki, S., Nishikura, K., Wada, T., and Hirao, T., Characterization of single crystalline CdTe surface, Appl. Suq. Sci., 59. 39, 1992.

381. Solzbach, U. and Richter, H. J., Sputter cleaning and dry oxidation of CdTe. HgTe and H&,,C&,,Te surfaces, Surf. Sci., 97, 191. 1980.

382. Brillson, L. J., Surface photovoltage and Auger spectroscopy studies of ( I 120) CdS surface, J. Vuc. Sci. Technol., 12, 249. 1975.

383. Morris, C. C. and Wood, B. J., Quantitative surface analysis of 11-VI compounds, Muter. Forum, 15.44, 1991.

384. Chieh, E. K. and Munir, Z. A., AES investigation of the composition of polar (OOOI ) surfaces of CdS, J . Muter. Sci., 26. 4268. 1991.

385. Ebina, A., Asano, K., and Takahashi, T., Surface properties of clean, and with adsorbed oxygen, surfaces of CdTe ( 1 10). [ 1 1 1 I , and (100) and of CdSe [ 0001 I studied by electron-energy-loss spectroscopy and Auger-electron spectroscopy, Phys. Rev. B. 22, 22, 1980.

386. Ricco, A. J., White, H. S., and Wrighton, M. S., X-ray photoelectron and Auger electron spectroscopic study of the CdTe surface resulting from various surface pretreanents: correlation of photoelectrochemical and capacitance-potential behavior with surface chemical composition,J. Vuc. Sci. Techno/.. A2.910, 1984.

387. Feldman, R. D., Opila, R. L., and Bridenbaugh, P. M., Auger electron spectroscopic study of the etching of cadmium telluride and cadmium manganese telluride, J. Vuc. Sci. Technol., A3, 1988, 1985.

388. Lu, Y.-C., Stahle, C. M., Feigelson, R. S., and Morimoto, J., Etching on polar (1 1 1 ) surfaces of CdTe crystals studied with Auger electron spectroscopy, J. Appl. Phys., 62, 4453, 1987.

389. Lu, Y.-C., Stahle, C. M., Morimoto, J., Bube, R. H., and Feigelson, R. S., Determination of crystallographic polarity of CdTe crystals with Auger electron spectroscopy, J. Appl. Phys.. 61, 924. 1987.

390. Zelaya, O., Sanchez-Sinencio, F., Gonzalez- Hernandez, J., Petia, J. L., Farias, M. H., Cota- Araiza, L., Hirata, G. A., and Rabago, F., Charac- terization of CdTe polycrystalline films by X-ray photoelectron and Auger spectroscopies, J. Vuc. Sci. Techno/., A7. 245, 1989.

391. Stahle, C. M. and Helms, C. R., Ion sputter effects on HgTe, CdTe, and HgCdTe. J. Vuc. Sci. Technol., A10, 3239. 1992.

392. Carin, R., Deville, J. P., and Werckmann, J., An XPS study of GaN thin films on GaAs, Surf. lnterfoce Anal., 16. 65, 1990.

393. Liau, Z. L., Brown, W. L., Homer, R., and Poate, J. M., Surface-layer composition changes in sputtered alloys and compounds, Appl. Phys. Letr., 30, 626, 1977.

394. Bondarchuk, A. B., Go-a, S. N., Koval’, I. F., Mel’nik, P. V., and Nakhodkin, N. G., Interaction of moderate-energy argon ions with a (1 1 l)B surface of gallium phosphide. Sov. Phys. Solid Stare, 28. 1577, 1986.

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

395. Yurasova, V. E. and Linnik, S. P., Sputtering of two-component alloys and compounds, 2. Phys. Chem.. 264, 1079. 1983.

396. Linnik, S. P., Buleev, M. A., Yurasova, V. E., Zaporozhchenko, V. I.,and Chernysh, V. S., Selec- tive sputtering of single crystals of binary semiconductor compounds, Rudiut. Efl.., 52, 191. 1980.

397. Maksimov, S. K., Egorov, V. L., Kryuk, V. V., Piskunov, D. I., Pitirimova, E. A., and Veselov, V. F., Effect of stoichiometric disturbances in AirlBV compounds during ion implantation, Phys. Status Solidi A, 73. K283, 1982.

398. Jones, R. G., Singh, N. K., and McConville, C. F., The c(4 x 4) clean surface reconstruction of InSb (001) formed by UHV chemical etching using iodine and chlorine, Surf. Sci.. 208. L34, 1989.

399. Bettini, M. and Richter, H. J., Oxidation in air and thermal desorption on PbTe, SnTe and Pb,,,Sno,2Te surfaces, Sutf. Sci., 80, 334. 1979.

400. Zagorenko, A. I. and Zaporozchenko, V. I., Quan- titative Auger electron spectroscopy of binary systems: matrix correction, Surf Interface A n d . . 14,438, 1989.

401. Roth, J., Bohdansky, J., and Martinelli, A. P., Low energy light ion sputtering of metals and carbides, Rudiut. Eff.. 48. 213. 1980.

402. Kaplan, R., Surface studies of epitaxial p-Sic on Si (100). J. Appl. Phys., 56, 1636. 1984.

403. Fukuda, S., Mohri, M., and Yamashina, T., Com- positional changes in surface region of SIC single- crystal due to hydrogen and helium ion irradiation. Fusion Technol.. 6, 420. 1984.

404. Bellina, J. J., Ferrante, J., and Zeller, M. V., Sur- face modification strategies for (100) 3C-SiC. J . Vuc. Sci. Technol.. A4, 1692, 1986.

405. Bellina, J. J. and Zeller, M. V., Stoichiometric changes in the surface of (100) cubic S ic caused by ion bombardment and annealing, Appl. Surf. Sci.. 25, 380, 1986.

406. Laidani, N., Capelletti, R., Elena, M., Guzman. L., Mariotto, G., Miotello, A., and Ossi, P. M., Spectro- scopic characterization of thermally treated carbon- rich Si,_,C, films, Thin Solid Films, 223, 114. 1993.

407. Laidani, N., Bonelli, M., Miotello, A., Guzman, L., Calliari, L., Elena, M., Bertoncello, B., Glisenti, A., Capelletti, R., and Ossi, P., Chemical and compositional changes induced by N+ implantation in amorphous S i c fi1ms.J. Appl. Phys., 74.2013, 1993.

408. Miotello, A., Calliari, L., Kelly, R., Laidani, N., Bonelli, M., and Guzman, L., Composition changes in Ar' and e--bombarded Sic: an attempt to distin- guish ballistic and chemically guided effects, Nucl. Instrum. Methods, B80, 931, 1993.

409. Castle, J. E., Liu, H. D., Watts, J. F., Zhang, J. P., Hemment, P. L. F., Bussrnann, U., Robinson, A. K., Newstead, S. M., Powell, A. R., Whall, T. E., and Parker, E. H. C., An investigation of Sio,,Geo,, alloy oxidation by high dose oxygen implantation, Nucl. Instrum. Methods, B55, 697, 1991.

410. Greene, J. E., Barnett, S. A., Cadien, K. C., and Ray, M. A., Growth of single crystal GaAs and metastable (GaSb),_,Ge, alloys by sputter deposition: ion- surface interaction effects, J . Cryst. Growth, 56. 389, 1982.

41 I . Nitz, H. M., Ganschow, O., Kaiser, U., Wiedmann, L., and Benninghoven, A., Quasisimultaneous SIMS, AES, XPS, and TDMS study of preferential sputtering, diffusion, and mercury evaporation in Cd,Hg,_,Te, Surf. Sci.. 104, 365, 1981.

4 12. Procop, M, XPS data for sputter-cleaned Iq,&a,, r7As. GaAs, and InAs surfaces, J. Electron. Specrrom. Relur. Phenom.. 59, R1. 1992.

413. Morgen, P. and Wilson, J. A., Preferential chemical reactions during noble gas ion sputtering of Hg,-,Cd,Te and its oxides, Nucl. Instrum. Methods, B26.585, 1987.

414. Valeri, S., di Bona, A., Engeli, E., Bordiga, S., and Piccirello, A., Auger electron spectroscopy study of cleaved and sputter-etched In, &a,, , 7 A ~ surfaces, Thin Solid Films, 197. 179. 1991.

415. Morgen, P., Silberman, J. A., Lindau, I., Spicer, W. E., and Wilson, J. A., Oxidation of Hg,_,Cd,Te studied with surface sensitive techniques, J . Electrori. Muter.. I I , 597. 1982.

416. Morgen, P., Silberman, J. A., Lindau, I., Spicer, W. E., and Wilson, J. A., AES sputter profiles of anodic oxide films on (Hg. Cd)Te,J. Vuc. Sci. Technol.. 21, 161, 1982.

417. Neff, H., Lay, K. Y., Abid, B., Lange, P., Lucovsky, G., and Bachmann, K. J., Sputtering and native oxide formation on ( I 10) surfaces of Cd,-,Mn,Te, J. Appl. PhTs., 60, 151, 1986.

418. Neff, H., Lay, K. Y., Su, M. S., Lange, P., and Bachmann, K. J., Sputter induced near surface electronic defects in group 11-VI compound semiconductor alloys, Surf. Sci.. 189, 661, 1987.

419. Rufus, I. B., Ramakrishnan, V., Viswanathan. B., and Kuriacose, J. C., Surface characterization of CdS,,,Se, 3R by X-ray photoelectron spectroscopy, J . Muter. Sci. Lett.. 11, 252, 1992.

420. Seelmann-Eggebert, M., Brandt, G., and Richter, H. J., Comments on the composition of anodic films on Hg,&d,,2Te, J . Vuc. Sci. Technol.. A2. 11, 1984.

421. Davis, G. D., Sun, T. S., Buchner, S. P., and Byer, N. E., Anodic oxide composition and Hg depletion at the oxide-semiconductor interface of Hg,-,Cd,Te, J. Vuc. Sri. Technol.. 19, 472. 1981.

422. Davis, G. D., Buchner, S. P., Beck, W. A., and Byer, N. E., Use of surface behavior diagrams to compare anodic and photochemical oxides of Hg,,&&,,Te, Appl . Surf. Sci.. 15, 238, 1983.

423. Sun, T. S., Buchner, S. P., Byer, N. E., Chen, J. M., Oxygen uptake on an epitaxial PbSnTe (1 1 1) surface, J. Vuc. Sci. Technol., 15. 1292, 1978.

424. Sun, T. S., Buchner, S. P., and Byer, N. E., Oxide and interface properties of anodic films on Hg,-,Cd,Te, J. Vuc. Sci. Technol.. 17, 1067. 1980.

425. Grant, R. W., Pasko, J. G., Longo, J. T., and Andrews, A. M., ESCA surface studies of Pb,-,Sn,Te devices, J. Vuc. Sci. Technol., 13, 940. 1976.

194

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

426. Chen, W. D., Bender, H., Demesmaeker, A., Vandervorst, W., and Maes, H. E., Quantitative Auger electron spectroscopy of AI,Ga,-,As layers and superstructures grown by MBE. Surf. Inregace Anal., 12, 156, 1988.

427. Christie, A. B., Sutherland, I., and Walls, J. M., Studies of the composition, ion-induced reduction and preferential sputtering of anodic oxide films on Hg,,,C&,,,Te by XPS, Sug. Sci., 135,225, 1983.

428. Sander , P., Altebockwinkel, M., Storm, W., Wiedmann, L., and Benninghoven, A., A versatile spectrometer system for quantitative surface and in- depth analysis with secondary ion and secondary neu- tral mass spectroscopy, Auger electron and X-ray photoelectron spectroscopy, J. Vuc. Sci. Techno/.. A7. 3305, 1989.

429. Chakarov, I. R., Todorov, S. S., and Karpuzov, D. S., Dynamic Monte-Carlo simulation of compositional change and atomic redistribution in multicomponent targets under ion bombardment. Nucl. Insrrum. Methods. 869, 193, 1992.

430. Bolkhovityanov, Yu. B., Bolkhovityanova, R. I., Vaulin, Yu. D., Katkov, A. V., and Ol’shanetskii, B. Z., Study of InGaAsP/GaAs( 11 l)A heterostructures with the aid of layer-by-layer Auger analysis, Inorg. Muter.. 25, 902, 1989.

43 I . Arthur, J. R. and LePore, J. J., Quantitative analysis of AI,Ga,-,As by Auger electron spectroscopy, J. Vac. Sci. Technol., 14, 979. 1977.

432. Wilson, J. A. and Cotton, V. A., (HgCd)Te-SIO, interface structure, J. Vac. Sci. Technol., Al. 1719. 1983.

433. Chen, W. D. and Cui, Y. D., Quantitative analysis of In,Ga,..,As by Auger electron spectroscopy, J . Vac. Sci. Technol.. A l l . 2379. 1993.

434. Morais, J., Machado, A. M., Sacilotti, M. A., and Landers, R., AES sputter depth profiles applied to interface analysis of GaInAs/InP grown by atmos- pheric pressure MOCVD. Appl. Surf. Sci.. 44, 161. 1990.

435. Olivier, J., Padeletti, G., Ingo, G. M., Mattogno, G., Bosacchi, A., and Franchi, S., Quantitative analysis of AI,Ga,-,As/GaAs multiquantum wells by means of AES depth profiling and small area XPS. Appl. Sug. Sci., 70/71, 89, 1993.

436. Padeletti, G. and Ingo, G. M., Small-area XPS investigation on ion-induced chemical modifications during depth-profiling of an Ae,Ga,-,As/GaAs structure, Surf. Itireface Anal., in press.

437. Erickson, L. P. and Phillips, B. F., Examination of MBE GaAs/AP,.,Ga,,,As superlattices by Auger electron spectroscopy,/. Vac. Sci. Technol.. B I , 158, 1983.

195

Dow

nloa

ded

by [

Uni

vers

itaet

s un

d L

ande

sbib

lioth

ek]

at 0

0:27

10

Dec

embe

r 20

13

Documents

Sputtering of compound semiconductor surfaces. II. Compositional changes and radiation-induced topography and damage