Chalcogen modification ofGaAs(100) surfaces and

metal/GaAs(100) contacts

von der Fakultät für Naturwissenschaften

der Technischen Universität Chemnitz

genehmigte

Dissertation

zur Erlangung des akademischen Grades

Doctor rerum naturalium

(Dr. rer. nat.)

vorgelegt

von Diplom-Physiker Stefan Hohenecker

geboren am 16.05.1968 in Hückeswagen

eingereicht am 5. Oktober 2000

Gutacher: Prof. Dr. Dietrich R. T. Zahn

Prof. Dr. Thomas Chassé

Prof. Dr. Iggy McGovern

Tag der Verteidigung: 3. Mai 2001

2

Bibliografische Beschreibung

Verfasser: Stefan Hohenecker

Titel: Chalcogen modification of GaAs(100) surfaces and metal/GaAs(100)

contacts

Erscheinungsort: Chemnitz

Erscheinungsjahr: 2002

Anzahl Seiten: 173

Referat

In dieser Arbeit wird der Einfluss der Modifikation der technologisch relevanten

GaAs(100) Oberfläche durch Chalkogene, i.e. Selen, Schwefel und Tellur

untersucht. Die modifizierten Grenzflächen werden dann mit Metallen

unterschiedlicher Reaktivität und verschiedenen Elektronegativitäten bedeckt.

Die Bandbreite dieser Eigenschaften wird durch die Metalle Indium und Silber,

das Alkalimetall Natrium, das Erdalkalimetall Magnesium und das Halbmetall

Antimon abgebildet. Die Untersuchung des Einflusses der Chalkogene auf die

chemischen Eigenschaften und die Barrierenhöhe der Metall/GaAs(100)

Grenzfläche bilden einen weiteren Schwerpunkt. Als experimentelle Techniken

werden Photoemissionsspektrokopie, Raman Spektroskopie und Strom-

Spannungsmessungen verwendet.

Schlagworte

GaAs(100), Sulphur, Selenium, Schottky contacts, Passivation, Schottky barrier,

Metal induced gap states (MIGS), Band bending, Surface, Photoemission

spectroscopy.

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1 Introduction .................................................................................................52 Models for the formation of the Schottky barrier .........................................9

2.1 Introduction..............................................................................................92.2 Schottky-Mott model................................................................................92.3 Bardeen model........................................................................................102.4 Advanced Unified Defect Model ............................................................122.5 Metal induced gap states.........................................................................13

3 Experimental..............................................................................................173.1 Sample preparation.................................................................................173.2 Soft X-ray photoemission spectroscopy..................................................213.3 Raman spectroscopy...............................................................................283.4 Transport measurements.........................................................................333.5 Low energy electron diffraction (LEED) ................................................35

4 Characterisation of clean and modified GaAs surfaces...............................394.1 Clean GaAs(100) surfaces ......................................................................394.2 Se modified GaAs(100) surfaces ............................................................464.3 S modified GaAs(100) surfaces ..............................................................604.4 Te modified GaAs(100) surfaces ............................................................75

5 Characterisation of chalcogen modified metal/ GaAs(100) contacts ..........835.1 In contacts on chalcogen modified GaAs(100) surfaces..........................835.2 Mg on chalcogen modified GaAs(100) surfaces .....................................925.3 Na on Se modified GaAs(100) surfaces ................................................1035.4 Sb on chalcogen modified GaAs(100) surfaces.....................................1105.5 Ag on chalcogen modified GaAs(100) surfaces....................................1175.6 Electronic properties of chalcogen modified metal/GaAs(100) contacts ................................................................................................127

6 Summary .................................................................................................141Acknowledgement..........................................................................................145Literature........................................................................................................147Selbständigkeitserklärung...............................................................................168Thesen ............................................................................................................169Curricula vitae ................................................................................................173

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1 IntroductionThe semiconductor GaAs plays an important role in optoelectonics and thefabrication of fast electronic devices. Especially the GaAs(100) surface isimportant for the commercial technology, because wafers grown in (100)direction are mainly used for the fabrication of devices. Therefore this surface isinvestigated here.Due to the ongoing miniaturisation the interfaces in complex devices are gettingmore and more important. Especially the metal/semiconductor interface is ofgreat importance, since all devices have to be contacted by a metal to apply avoltage. Between metal and semiconductor a potential barrier is build, calledSchottky barrier. The height of this barrier influences the character of themetal/semiconductor contact. If the barrier height is low the contact shows anohmic behaviour and if it is high the contact shows an rectifying behaviourwhen a voltage is applied. The latter, which is called Schottky contact, is ofgreat importance to predict the properties of the metal/semiconductor contact.Therefore there are several models to explain the origin of the Schottky barrier.The main models are described in chapter 2.In many cases the metal deposition on semiconductor surfaces causes a chemicalreaction of the metal with the semiconductor. That yields to new electronicproperties of the semiconductor surface and the metal/semiconductor interface,which makes it more difficult to predict the properties of the Schottky barrier bya model. A suitable technique for the observation of the interface reactions is thephotoemission spectroscopy, because of its high surface sensibility. Thephotoemission spectroscopy can also be applied for the determination of thebarrier height up to some monolayers of adsorbate thickness. In chapter 3 theexperimental techniques used in this study are described. These are besidesphotoemission spectroscopy also Raman spectroscopy for the determination of achange in band bending and current-voltage measurements for the determinationof the barrier height in technological relevant metal/semiconductor contacts.Much of the success of the silicon-based semiconductor devices depends on theeasy fabrication of Si:SiO2 interfaces, which has been optimised to give only alow density of electrically active interface states. That means that the SiO2 actsas an ideal passivant for Si [Gre93]. On GaAs the native oxides do not lead to achemically stable and defect free interface [Spic80b, Thu80]. Furthermore thelarge density of states at the interface lead to strong Fermi level pinning aroundmidgap. These states also act as electron-hole recombination centers in theactive region, which leads to a reduced device performance [Gre93].Therefore the idea arose to use other chalcogens like sulphur or selenium tosatisfy the dangling bonds on the GaAs(100) surface and to reduce the density ofstates in the bandgap. Sandroff et al. were the first who used an aqueous solutionof NaS2, in which the GaAs was dipped to passivate the surface [San87]. Thisyielded to a significant enhancement of the photoluminescence intensity and an

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improvement of a factor 60 of the current gain of a heterojunction bipolartransistor. Thereafter great attention has been paid to the passivation usingsulphur containing solutions and gases such as Na2S [Ali96, Bes96a-d, Far87,Has88, Shin89], S2Cl2 [Cao97, Li94, Wan96], (NH4)Sx [Ali96, Cow89, Ha95,Lu93, Shin89, Sug91, Yua97], [(NH4)S+Se] [Bel96], SeS2 [Kur96a, Kur96b],H2S [Con96, Con97, Gei90, Shin89, Xin96] and by molecular beam epitaxy(MBE) of S2 in ultrahigh vacuum (UHV) [Koe88, Mor94]. Additionally,attention has been paid to the passivating properties of other chalcogens likeselenium, which was deposited onto the GaAs surface via metal organicchemical vapour deposition (MOCVD) [Cha90, Cha91], wet chemicalpreparation with a solution of Se in NH4OH [Tsu95] and by MBE of elementarySe in UHV [Cai92, Hag96, Lid94, Miw96, Pas94, Sci92, Sci93].The various methods of preparation and different experimental circumstancesmake it difficult to compare the results and partly they seem to be inconsistent.The application of photoemission spectroscopy has shown that the chalcogentreatment results in a clear decrease of band bending, which was stated from ashift in the binding energy of core levels. But there are some discrepancies in theamount of the reduction. Takatani et al. [Tak92a] reported nearly flat bandconditions while this was not achieved by Scimeca et al. [Sci92]. Thepassivating effect is mainly attributed to the heterovalent exchange reaction ofAs and the chalcogen, which strongly depends on the temperature. Even at roomtemperature a partially As-chalcogen exchange takes place, but it is alsoreported that As-Se [Sci92] and As-S [Sug91] bonds are formed, respectively.Only at higher temperature the excess As and the arsenic chalcogenides vanish,which results in the clear reduction of band bending [Sci92].Up to now there is no generally accepted model for the structure of thechalcogen modified GaAs(100) surface, although different models have beenproposed [Bie94, Gun98, Hir98, Pas94]. Among the structural models proposedto describe the chalcogen modified GaAs surface there are some which involvesurface As [Bie94, Ber92, San89] or Ga-chalcogen dimers on the surface[Wan94].The exchange between As and Se has been limited to the first three layers byScimeca et al., while Chambers et al. [Cha90, Cha91] found out an exchange inthe top five or seven layers. Nevertheless the formation of a Ga2Se3 [Men91,Tak92a, Tak92b] or Ga2S3 [But96, Wan96] like compound on top of the surfaceis favoured. But there is still a discrepancy in the interpretation of the differentcomponents in the core level spectra of Se 3d and S 2p [Sci92, Tak92a, Con96,Mor94].When prepared under UHV conditions with subsequent annealing the Se and Smodified surfaces exhibit a (2x1) LEED pattern, which is generally accepted[Mae93, Hag96, Miw96, Pas94, Con96, Mor94]. The (2x1) LEED patter hasbeen found independent from the fact, whether the reconstruction of the cleanGaAs surface was As-rich c(4x4) or Ga-rich (4x2).

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In chapter 4 of this study the chalcogen passivation of n- and p-dopedGaAs(100) surfaces is investigated. An in situ dry chemical process is used forthe preparation of the surface, because it has the advantage to deposit a welldefined amount of chalcogens onto the GaAs surface under well definedconditions in contrast to a wet chemical treatment.Since the chalcogen modification of the GaAs(100) surface has attracted a lot ofinterest up to now, there are only a few studies dealing with the interaction ofmetals with these modified surfaces. Among these are the metals Au [Wal85a],Pd [Berr92, Osh93] and Al [Osh93].The influence of S and Se on the interface of Na, Mg, In, Sb and Ag contacts onGaAs(100) are explored in chapter 5 of this study.

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2 Models for the formation of the Schottky barrier

2.1 IntroductionWhen metals and semiconductors are brought together they may exhibitinteresting properties. This is the case for some metal-semiconductor interfacesthat, depending on the sign of the applied bias, let the current flow or not. Theserectifying properties were explained by Schottky [Scho38] by the presence of aspace-charge layer on the semiconductor side of the interface that is depleted ofits carriers. For these reason these metal-semiconductor interfaces were calledSchottky contacts.

2.2 Schottky-Mott modelA model proposed for the metal-semiconductor interfaces with rectifyingproperties was proposed by Mott [Mot38] and Schottky [Scho40]. It is assumed

that no state exist in the band gap of the semiconductor so that no interface

Figure 2.1: Energy band diagram to describe the Schottky model of a contact between ametal and a n-type semiconductor. EF is the Fermi level, EVac the vacuum level, CBM theconduction band minimum, VBM the valence band maximum, Egap the gap width, Vn thedifference between the CBM and the Fermi level in the bulk, eVBB the band bending onthe surface, ΦBn the Schottky barrier height, ΦM the metal work function, ΧA the electronaffinity of the semiconductor, w the width of the depletion zone, I the ionisation energy,QM the interface charge in the metal and QSC the space charge in the semiconductor.

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dipole can be induced. On the semiconductor side there exists only a spacecharge. At equilibrium, the Fermi levels must align in the metal and thesemiconductor. This results in a charge transfer, which creates an accumulationcharge QM on the metal surface and a space charge QSC in the semiconductor(figure 2.1). In the metal the charges remain on the interface because of the highdensity of electrons (≈ 1022 cm-3), which screen the charge. In contrast the spacecharge on the semiconductor is due to uncompensated ionised donors, whosedensity depends on the doping concentration, typically between 1014 and 1018

cm-3. This is much lower than on the metal so that the charge extends into thesemiconductor and creates a band bending. The length scale for the bandbending amounts to approximately 100 to 1000 nm. Because the whole systemis electronically neutral, formula 2.1 is valid.

SCM QQ −= (2.1)

From figure 2.1 it can be deduced that the Schottky barrier height dependslinearly on the metal work function according to formula 2.2.

Χ−Φ=Φ MBn (2.2)

This model is far too simple since it neglects the role of surface and interfacestates. This has already been shown by Schottky [Scho39] by experiments onmetal-selenium contacts. He concluded that new models have to take these statesinto account.

2.3 Bardeen modelThe Schottky model ignores the role of surface states at the metal-semiconductor interface. Since it is well known that a clean semiconductorsurface may show states in the gap, Bardeen [Bar47] proposed a model wherethese states play a major role. These states are filled up to a level Φ0 and theycarry a charge QSS (figure 2.2 a.). Because the charge neutrality of the system,the charge in the surface states QSS must be compensated by a space chargeinside the semiconductor QSC leading to equation 2.3.

0=+ SCSS QQ (2.3)

This space charge is associated with a band bending and depends on the densityof the surface states. If the semiconductor surface is brought into contact with ametal (figure 2.2 b.) the charge of the surface states is compensated by the samecharge of opposite sign on the metal side. In the case of a high density of surfacestates the semiconductor surface behaves like a metal, whose Fermi levelcoincides with Φ0. Therefore all the charge transferred from the metal can

11

accommodate in the surface states without noticeably moving the Fermi levelfrom Φ0. The charge transfer between the metal and the semiconductor creates adipole ∆ which leads to an alignment of the Fermi level in the metal and Φ0.According to this the Schottky barrier is given by formula 2.4.

0Φ−=Φ gapBn E (2.4)

Figure 2.2: Band diagrams of a.) a clean metal and a clean semiconductor surface andb.) a metal in contact to a n-type semiconductor. The band diagram explains the Fermilevel pinning by states in the semiconductor band gap. QSS is the charge in the surfacestates.

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The space charge region in the semiconductor remains unchanged due to thecontact to a metal. This leads to the conclusion that the Schottky barrier is anproperty of the semiconductor and completely independent of the metal.Therefore EF is pinned near Φ0 because of the high density of states at thesurface.Bardeen is the originator of the idea that the Fermi level is pinned by states inthe semiconductor band gap.

2.4 Advanced Unified Defect ModelSpicer proposed a model where the Fermi level is pinned by defect levels in theband gap. This has been called the Unified Defect Model (UDM) [New86,Spic79, Spic80a, Spic80b, Spic81, Spic85] and after the identification of thedefect levels, which are responsible for the pinning of the Fermi level in GaAs,advanced unified defect model (AUDM) [Spic88, Spic89, Spic93]. This modelcan also explain the pinning of the Fermi level on clean (110) surfaces of III-Vsemiconductors, which usually show no surface states in the band gap. It iscalled unified defect model, because it applies to III-V-insulator and III-V-metalinterfaces.

If ΦD is the defect level in the semiconductor band gap, relative to the top of thevalence band, and if the density of defects is high enough, the charge transferbetween the metal and the defect states creates a dipole that aligns the Fermilevel EF and the defect level ΦD. This leads to a pinning of the Fermi level atthese states according to formula 2.5.

Figure 2.3: a.) The energy diagram for the advanced unified defect model shows theAsGa antisite double donor with the two distinct energy levels and the compensatingacceptor, which probably can be attributed to GaAs antisites. The energy levels of theacceptors lies below 0.5 eV. Both defects are located in the same spatial region near thesurface [Spic88]. b.) Energies of the AsGa antisite related to the valence band determinedby photospin resonance at 8 K [Web82, Web83].

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DgapBn E Φ−=Φ (2.5)

Studies of the GaAs system showed that antisites are the main defects [Spic88].The key defect in GaAs is the AsGa antisite, which is a double donor with levelsat 0.52 and 0.75 eV above the valence band maximum (figure 2.3). Thesedonors are partly compensated by acceptors, most likely GaAs antisites.Therefore the Fermi level depends on the doping and the ratio R between theAsGa and GaAs antisite. This ratio mostly exceeds 1, because the GaAs crystalsare grown from the As-rich side of the phase diagram.In this model, As-rich interfaces are expected to be pinned near the midgaplevel, whereas less As-rich interfaces should exhibit a pinning near the lowerlevel. In conclusion the advanced unified defect model explains why the Fermilevel can be pinned at different levels in the gap of n- and p-type semiconductorssince it involves donor and acceptor defects.

2.5 Metal induced gap statesA microscopic description of interface states was first given by Heine [Hei65].The interface states consist of virtual states in the band gap of the complex bandstructure of a semiconductor. These states are called virtual induced gap states(VIGS). A complex wave function with real energy Eigenvalues makes onlysense if it is located at the surface. This wave functions decrease periodicallyinto the semiconductor while they decrease exponentially into the vacuum[Mön89]. They can be normalised and represent a state, which is located at thesurface. The states which are derived from the valence band exhibit a donor likecharacter while the states derived from the conduction band exhibit an acceptorlike character. The point where the influence of both bands is equal is called thebranch-point. This point corresponds to a local neutrality of the surface andtherefore it is also called charge neutrality level (CNL). If a semiconductor isbrought in contact to a metal, the conduction band of the metal overlaps theenergy region between the valence band maximum and the Fermi level of thesemiconductor. The metal wave functions which leaks into the semiconductorcan be derived from the VIGS. These metal induced gap states (MIGS) are anintrinsic property of the semiconductor. The continuum of MIGS is occupied upto the Fermi level and empty above it. The states in the middle of the band gapexhibit the largest damping. In analogy to the VIGS the MIGS also possesses acharge neutrality level. If the Fermi level is positioned above (below) the CNLthe MIGS are negatively (positively) charged.The charge on the semiconductor side consists of the charge in the depletionlayer QSC and the charge in the MIGS QMIGS, which is compensated by a chargein the metal QM. The space charge in the depletion layer is much smaller thanthe charge in the MIGS and thus may be neglected in the followingconsideration [Mön96a].

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Interface charge neutrality demands that the charge densities on thesemiconductor and metal side are of equal magnitude but have opposite sign.Therefore a dipole layer of the dipole strength ∆ is formed at the interface andthe subsequent band diagram is shown in figure 2.4. According to Pauling[Pau60] the charge transfer across the interface may be described by thedifference of the metal ΧM and semiconductors ΧS electronegativities. Thebarrier height ΦBn of the semiconductor-metal contact can be calculatedaccording to formula 2.6 after the predictions of the MIGS model [Mön87].

( )SMXCNLBn S Χ−Χ⋅+Φ=Φ (2.6)

The slope parameter SX = δΦBn/δΧM was estimated from a semi-empirical rule[Mön86b] and also justified theoretically [Mön96a]. It correlates with theelectronic part of the static dielectric constant (formula 2.7).

( )211,01 −⋅=− ∞εXX SA (2.7)

The parameter AX accounts 0.86 for the use of Miedema’s electronegativities and1.79 for the use of Pauling’s electronegativities [Mön95].The MIGS model is relevant for an ideal and abrupt metal-semiconductorcontact. The predictions of the MIGS model can be influenced by defects or

Figure 2.4: Energy band diagram to explain the Schottky barrier of a contact between ametal and a n-type semiconductor. CNL is the charge neutrality level of the metalinduced gap states.

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additionally deposited atoms at the interface. These can induce interface statesor dipoles. It is found that metallic adatoms induce surface states, which act asdonors on semiconductor surfaces [Mön88], while other adatoms induceacceptor like states [Mön95]. These states can only influence the barrier heightif they persist after the deposition of a metal on the modified semiconductorsurface.The Schottky barrier height at the surface can also be influenced by interfacedipoles. The change in the barrier height ∆φB can be described according toformula 2.8 [Top27].

ad

ad

i

iB N

N

pe⋅

+⋅±=∆Φ ⊥

2/30

0

91 αεε (2.8)

e0: elementary charge

ε0: dielectric constant of vacuum

εi: dielectric constant of the semiconductor at the interface

⊥ip : dipole moment normal to the interface

α: polarizability of the adatoms

Nad: surface density of the adsorption-introduced dipoles

The sign of ∆φB is determined by the orientation of the adsorption-introduceddipoles or, in other words, by the charge distribution between adsorbate andsubstrate atom. The factor 9 in the denominator of formula 2.8 takes thegeometry into account [Top27]. The dipole moment normal to the interface canbe calculated according to formula 2.9.

⊥⊥ ⋅∆⋅= iii dqep 0 (2.9)

∆qi: charge of the interface dipole

⊥id : length of the dipole normal to the interface

The charge of the interface dipole depends on the electronegativity difference ofthe participating atoms, which can be calculated after an empirical formula fromPauling, which was modified by Hanney [Han46]. This formula is valid forsingle bonds in diatomic molecules (formula 2.10).

2035.016.0 BABAi XXXXq −⋅+−⋅=∆ (2.10)

XA: electronegativity of atom A

XB: electronegativity of atom B

16

17

3 ExperimentalThe following chapter gives an overview of the different experimental methodsapplied in this study. The theoretical background of these is briefly describedand also the reason why the application of this method is convenient for thepurpose of this investigation and what can be learned from it.

3.1 Sample preparation

3.1.1 Preparation of clean samples

In this study GaAs(100) samples with different type of doping (n and p) and adoping concentration between 1⋅1016 cm-3 and 5⋅1018 cm-3 are investigated. Toexamine the properties of an atomically clean surface an UHV with a basepressure of less than 2⋅10-10 mbar is necessary, because otherwise the surfacewill be soon covered by other atoms. The number of atoms, which is necessaryto build up a monolayer strikes the surface every second at a pressure of 10-6

mbar [Lüt95]. For the preparation of a clean GaAs(100) surface in UHV theGaAs epilayers are covered with a thick As cap layer after growth. This layerprotects the GaAs surface from contamination and oxidation. The typicalthickness of the cap is approximately 10 µm. These specially prepared samplesare provided by the group of D. Westwood from the Cardiff College of theUniversity of Wales. They put epiready GaAs(100) wafers in a MBE chamberwhere a homoepitaxial approximately 2 µm thick GaAs layer with a certaindoping concentration is grown on top of the wafers. At the end of the growthstep the Ga source is closed by a shutter, while As is still evaporated onto theGaAs layer to produce the cap [Res94, Res96]. When the capped GaAs sampleis introduced into the UHV the As cap can be removed by gentle annealing,which leads to the evaporation of the excess As on top of the sample. Thedifferent possible reconstructions of the GaAs(100) surface can be tuned by theannealing temperature [Mön95]. In this study mainly the As rich c(4x4) surfaceis prepared, which is achieved by annealing the sample to approximately 380 °C.The reconstruction of the surface can be identified by LEED and PES, whichwill be described later (see chapter 4).

3.1.2 Evaporation of chalcogens

Three different species of chalcogens are used in this study: Sulphur (S),selenium (Se) and tellurium (Te).The easiest to handle is the Te, because its gas pressure is high enough (133 Paat 520 °C) to evaporate it in UHV from the solid phase. A Knudsen cell likeoven is used for its deposition. The deposition rate is monitored by a quartzcrystal microbalance and controlled by the current from a stabilised powersupply.

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The evaporation of Se can not be done that easy, because Se has a vapourpressure of 0.13 Pa at 200 °C [Ber68], which corresponds to the bake outtemperature of the UHV system. Therefore the compound SnSe2 is used for theevaporation of Se to avoid that the Se gets lost during the bake out. Thiscompound decomposes at approximately 340 °C [Shi92] as follows:

SnSe2 → SnSe + Se ↑ . (3.1)

At a temperature of more than 400 °C a second stage of decomposition starts atwhich also SnSe3 is evaporated. This would lead to an unwanted contaminationof the samples with Sn. Therefore the temperature of the Knudsen cell ismonitored by a Ni-Ni/Cr thermocouple to avoid this. The deposition rate of Sewhich can be derived from a quartz crystal microbalance is controlled by theheating current.The S source is similar to the Se source as it uses SnS2, because the evaporationof S causes the same problems as Se. The vapour pressure of S is even lowerthan that of Se. To avoid the problems of evaporating S during the bake out, onecan use the compound SnS2 similar to SnSe2 which is used for the evaporationof Se. According to Shimada [Shi92] SnS2 decomposes at approximately at 550°C as follows:

SnS2 → SnS + S ↑ . (3.2)

Applying temperatures higher than 740 °C leads to the formation of SnS3 similarto the SnSe2 compound. Therefore it is important to keep the SnS2 in the righttemperature range, which is also monitored by a Ni-Ni/Cr thermocouple.

3.1.3 Evaporation of metals

The chalcogen modified semiconductor surfaces serve as a starting point formetal deposition. Five different kinds of metals were used to form metalcontacts on the modified GaAs(100) surfaces: Na, Mg, In, Sb and Ag. They arechosen for this study, because they represent a wide spectrum of different workfunctions and electronegativities (compare to table 3.1).The formation of metallic contacts consisting of Na is difficult, because pure Nais immediately oxidised as soon as it is exposed to air. To avoid these problemsNa dispensers of the firm SAES Getters are used. These consist of a metalmatrix, which contains elemental Na. To get pure Na an heating current ofapproximately 6 A has to be applied to the dispensers to activate them. Butbefore Na can be evaporated the dispensers have to be outgased by applying alower current for several hours. To avoid contamination with oxygen and otherelements a vacuum of less than 2⋅10-10 mbar has to be achieved in the UHVchamber. An additional step to avoid contamination of the Na is to keep thedispensers always on temperature after the activation.The Mg and Ag source each consist of a piece of wire, which is wrapped by athin tungsten wire of 0.1 mm diameter. This wire is always connected to an

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electrical feedthrough. The elements are evaporated by applying a constant

current to the tungsten wire. This simple setup is possible since tungsten doesnot build any alloys with Mg or Ag. No cooling was necessary, because theevaporation temperatures are not high enough to cause outgassing from thevessel or other components within the vacuum.For the evaporation of Sb granules of the pure material are filled into a littleenvelope out of tungsten foil. This envelope is punctured several times andwrapped by a thin tungsten wire (∅ 0.1 mm). To evaporate Sb the wire isresistively heated by a constant current. Sufficiently high evaporation rates canbe reached by temperatures, which are not so high that a cooling of theevaporation cell is necessary.

Na Mg In Sb AgVapour

pressure/kPa at 133 Pa

58.7 82.7 261.3(at 53.3 kPa)

128.0(at 1.3 kPa)

174.7

Atomic weight/amu

22.99 24.31 114.82 121.75 107.87

Atomic radius(Goldschmidt)/

nm

0.192 0.160 0.157 0.161 0.144

Work function/eV

2.12 3.66 4.12 4.1 4.7

Density/g cm-3 0.97 1.74 7.3 6.68 10.5Melting point/

° C97.8 649 156.6 2310 961.9

Boiling point/ ° C

883 1090 2080 3900 2212

Purity 6N 3N 5N 6N 4NElectro-

negativity afterPauling[Pau60]

0.93 1.31 1.78 2.05 1.93

Electro-negativity after

Miedema[Mie73, Mie80]

2.70 3.45 3.90 4.40 4.45

Table 3.1: Properties of the metals, which are examined in this study to investigate theirbehaviour on chalcogen modified GaAs(100) surfaces.

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Because of its high vapour pressure the In source needs to be cooled during theevaporation. The In granules are put into cell, which has an integrated watercooling to avoid the warm up of the environment.All evaporation sources are carefully outgassed before metallic contacts areformed on the semiconductor surfaces. Therefore the pressure in the vacuumchamber never exceeds 5⋅10-9 mbar during the metal evaporation.The thickness of the evaporated metal films is evaluated from the frequencychange of a quartz crystal microbalance, which is located near thesemiconductor sample or near the evaporation source by formula 3.3 [Sot92].

2

21

1

22

2

⋅

−−

−⋅⋅=

l

lNd qqq

m

qq ν

ννννν

ρρ

(3.3)

d: Thickness of the evaporated film

ρq: Density of the quartz (2.20 g cm-3)

ρm: Density of the evaporated film

νq: Resonant frequency of the uncoated crystal (6 MHz)

ν1: Resonant frequency of the crystal before the evaporation

ν2: Resonant frequency of the crystal after the evaporation

lq: Distance between source and crystal

l: Distance between source and substrate

Nq: Frequency constant for an quartz crystal vibrating in the

thickness shear mode (1.668∗ 105 cm/s)

The Tooling-factor (lq/l)2 represents the different geometrical arrangement. It is

valid if the substrate and the quartz crystal have a different distance from theevaporation source.

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3.2 Soft X-ray photoemission spectroscopyAlmost all core level spectra shown in this study were recorded using thesynchrotron radiation provided by the ”Berliner Speicherringgesellschaft fürSynchrotronstrahlung mbH” (BESSY). This is a synchrotron radiation source ofthe second generation, which means that it is build only for the application ofthe radiation predominantly generated in the bending magnets which hold theelectron beam on its course. The high energy limit of the synchrotron radiationis given by the maximum energy of the electron beam in the storage ring. Since

BESSY works with electrons up to 800 MeV it is thus dedicated to the Soft X-ray and Ultraviolet spectral region from 5 to 4000 eV.To monochromatise the synchrotron radiation different types ofmonochromators are used. The monochromator where all measurements of thisstudy were done was a toroidal grating monochromator (TGM 2). Thismonochromator provides three different gratings for certain energy regions(grating 1: 10-30 eV, grating 2: 20-100 eV, grating 3: 80-200 eV) [Bra86]. Allcore level spectra of As 3d, Ga 3d, In 4d, Mg 2p, Na 2p, Sb 4d and Se 3d shown

Figure 3.1: Escape depth of photoelectrons in different semiconductors depending on thekinetic energy of the photoelectrons [Ber88].

22

in this study were recorded by using grating 2. The S 2p core level has to beexcited by light from grating 3 because of its high binding energy.The beamline is equipped with a UHV chamber, which consists of threedifferent vacuum chambers. The vacuum system is pumped by turbo molecularpumps, an ion getter pump and a titanium sublimation pump. A base pressure of1⋅10-10 mbar is achieved after baking the system at 180 °C.

The first chamber includes a load lock to transfer new samples into the vacuum.The second chamber is used for analysis of samples. It contains an VG ADES400 photoelectron spectrometer, which can be rotated in the polar and azimuthplane. The third chamber serves for the sample preparation. It has severalflanches, where different Knudsen cells for the evaporation of certain elementscan be added. All chambers can be separated by valves. The samples aremounted on a manipulator, on which they can be moved in all three directions ofspace and rotated in the polar and azimuth plane. Additionally the manipulator isequipped with a heater, which achieves a temperature of approximately 600 °Cat maximum. The temperature is measured using a Ni-Ni/Cr thermocouple,

30 31 32 33 34 35 36 37

d3/2d5/2

measured points sum of components Se1 component Se2 component Background Residuum

Photon energy = 88.2 eVIn

tens

ity (a

rb. u

nits

)

Kinetic energy (eV)Figure 3.2: Fitted photoemission spectrum of the Se 3d core level with two spin orbitsplit components. The measured points agree quite well with the sum of the twocomponents Se1 and Se2 and the background. The remaining residuum on the bottomdoes not show any features, which could be identified as an additional component.

23

which is mounted near the sample (for more information about the vacuumsystem compare to [Now95]).The excitation energy for the core level spectra was always selected to providemaximum surface sensitivity, which means that the kinetic energy of thephotoelectrons emitted from the sample was between 30 and 40 eV (compare tofigure 3.1). This corresponds to a minimum escape depth of approximately 0.5nm in GaAs(100) [Ber88]. The small escape depth of the photoelectrons isimportant for the measurement of the position of the Fermi level at the surface,because the band bending decreases exponentially into the bulk of thesemiconductor.The photoelectrons are emitted due to the irradiation of the examined samplewith monochromatic synchrotron radiation The binding energy of the elementscan be calculated from formula 3.4.

EBin = hν − EKin − Φ (3.4)

EBin: Binding energy

hν : Photon energy EKin: Kinetic energy

Φ : Work function of examined sample

But the binding energy of an element is not only an intrinsic property, it alsodepends on the surrounding elements. Diatomic, heteropolar moleculesgenerally exhibit dipole moments. These dipole moments can be described bythe partial ionic character of the covalent bond between the atom which form thedipole. The ionic character of single bonds in diatomic molecules was correlatedwith the difference of the atomic electronegativities of the concerned atoms[Pau60]. Formula 3.5 represents a revised version [Han46] of the originalrelation from Pauling giving the charge transfer in a diatomic molecule.

∆q = 0.16 XA − XB + 0.035 XA − XB 2 (3.5)

XA,B: Electronegativity of element A,B

∆q: Transferred charge

In a simple model the more electronegative atom gets charged by −∆qe0 whilethe less electronegative atom is charged by +∆qe0. Therefore the core level of anatom which is bonded to a more electronegative element than another, will occurat higher binding energy in the photoemission spectrum.

24

A typical photoemission spectrum is shown in figure 3.2. To extract the physicalinformation out of the measured spectrum, these were fitted using spin-orbitsplit Voigt profiles [Arm67, Kie73, Wer74, Whi68] and a Shirley background[Shir72], which consists of a third order polynom and an overlaying step

function, which is proportional to the signal height of the core level. In the fitroutine (program BFIT [Hemp93]) the values for the spin orbit splitting, thebranching ratio and the Lorentzian widths were kept constant for all spectra,whereas the intensities, the binding energy shifts and the Gaussian widths werevariable. The same Gaussian width was used for all components in one certainspectrum. For metallic overlayers an additional Doniach-Sunjic Parameter α[Doni70] has to be taken into account due to the many particle effects. Thefitting of the spectra followed the principles laid down by Joyce [Joy89].To determine the absolute binding energy of the core levels a spectrum of theFermi edge on a clean peace of metal in contact with the sample was measuredbefore moving the monochromator to the next photon energy. The Fermi levelwas evaluated by fitting three lines in the photoemission spectrum and thenadjusting the density of the unfilled states below the Fermi level and the densityof the filled states above it. The density of states is always proportional to thearea of the triangles T1 and T2 shown in figure 3.3.The nominal photon energy from the control of the beamline does notcorrespond to the measured photon energy from the Fermi level. Therefore it hasto be determined very carefully for each measured spectrum, because it directlyinfluences the calculated binding energy (table 3.2). The width of the Fermilevel where it decreases from 90 to 10% of its height is also a measure for theresolution of the whole system including the monochromator and the usedADES 400 spectrometer (VG Ltd.). The values obtained for the photon energiesused in the experiments are given in table 3.2. These values vary a little bit,because the resolution depends on many factors in the set up of the beamline,esp. the beam position.

Nominal photonenergy / eV

Photon energydeterminedfrom Fermilevel / eV

Energyresolution / eV

Valence band 26 23.93 0.19 – 0.42Ga3d 65 60.30 0.20 – 0.33As3d 85 78.80 0.29 – 0.48

Table 3.2: Nominal photon energy from the control of the monochromator, photonenergy as measured by the determination of the Fermi level and Energy resolution of themonochromator at different photon energies.

25

The photoemission spectroscopy is also applied to measure the band bending onthe semiconductor surface during the modification of the GaAs(100) surface.Therefore the escape depth of the photoelectrons must be much smaller than thethickness of the depletion layer. This is fulfilled, because the escape depth of thephotoelectron in the observed kinetic energy range amounts to approximately0.5 nm while the typical width of the depletion region in a doped semiconductoramounts to 100 nm. From figure 3.4 it is obvious that the band bendinginfluences not only the valence band maximum (VBM) but also the bindingenergy of the Ga 3d core level, which originates from a Ga atom in the depletionregion. Therefore the energetic position of the Fermi level with respect to thevalence band maximum is determined from the spectra of the Fermi level on aclean peace of metal which is in contact with the sample and the valence bandspectrum from the sample.With this value the binding energy of the Ga 3d core level with respect to theVBM can be determined. As it can be seen from figure 3.4 this value does notdepend on the band bending. The same value can be determined for the As 3dcore level. If the surface of the sample is modified by chalcogen treatments forexample, the position of the Fermi level can be determined from the energyposition of the Ga 3d and As 3d core levels with respect to the Fermi level. It is

59,5 60,0 60,5 61,0

EF

T2

T1

Kinetic energy / eVFigure 3.3: Determination of the Fermi level from the photoemission spectrum of a cleanpeace of metal. Because the signal from the photoelectron analyser is corrected by thework function of the analyser, the Fermi level corresponds to photon energy accordingto formula 3.4.

26

important that the core level spectra are decomposed into the contributions fromthe GaAs bulk and from the reacted compounds at the surface to get the rightposition of the binding energy.In this study always highly doped GaAs samples have been used (n, p = 1⋅1018

cm-3) to avoid surface photovoltage [Alo90].

The growth mode of evaporated materials on the semiconductor surface can alsobe determined by photoemission spectroscopy. The intensity of thephotoemission signal of the elements of the bulk has to be followed as a functionof metal coverage. In the case of a laminar growth the signal intensity decreasesexponentially with the thickness of the growing film on the surface, which canbe approximated by formula 3.6.

I = I0 ⋅ exp ( − d / λGaAs (E) ⋅ cos θ ) (3.6)

I: Intensity of the substrate signal with a grown film of thickness d I0: Intensity of the clean substrate

d: Thickness of the grown film

Figure 3.4: Determination of the position of the Fermi level above the valence bandmaximum from the position of the Ga 3d core level relative to the Fermi level.

27

λGaAs (E): Escape depth of photoelectrons in GaAs at kinetic energy E θ : Angle relative to the surface normal

In figure 3.5 the different possible growth modes are shown. In the case of aVolmer-Weber growth islands grow on the surface of the substrate. Therefore

0 10 20 30 40 500.00001

0.0001

0.001

0.01

0.1

1

10

100

λ(EKin): escape depth of photoelectrons

I=A*exp(d/λ(EKin))c)

b)

a)

λ=5

λ=30

λ=30

λ=5lo

g(In

tegr

ated

inte

nsity

of i

/ (a

.u.))

Metal coverage d / a.u.

Figure 3.5: Characterisation of different growth modes and the observed photoemissionsignal from elements of the substrate for a) Volmer-Weber growth (islands), b) Stranski-Krastanov growth (layer and islands) and c) Frank-van der Merwe growth (layer bylayer) [Lüt95]. In the figure the escape depth for photoelectrons of GaAs at maximumsurface sensitivity, which amounts to approximately 5 ngstrm, is used.

28

big areas of the surface are not covered by the evaporated elements whichresults in an attenuation length, which is much bigger than the experimentaldetermined attenuation length in GaAs [Ber88]. In the case of a Stranski-Krastanov growth the deposited material forms several complete layers on top ofthe substrate before an island growth starts. This results in a photoemissionintensity of the substrate, which firstly decreases exponentially with anattenuation length comparable to the escape depth of the photoelectrons beforethe attenuation length increases due to the formation of islands. In the case of aFrank-van der Merwe growth the evaporated material forms layer by layer onthe substrate. Therefore the intensity decreases exponentially with thephotoelectron escape depth as attenuation length. The course of thephotoemission intensity from the substrate thus gives a hint on the growth modeof the deposited materials on the investigated semiconductor surfaces.

3.3 Raman spectroscopyRaman spectroscopy is a powerful method to study the effect of chalcogens onthe band bending at GaAs(100) surfaces. The Raman measurements are carriedout with a Dilor XY Raman spectrometer equipped with a CCD camera formultichannel detection. The 488.0 nm line of an Ar+-Laser is used for excitation.The general experimental set-up for measuring the Raman scattering from

0 2 4 6 8 10 12 14 16 18 200

100

200

300

400

500

600

700

800

ωTO

ωLO

1 x 1017 4 x 1017 4 x 10182 x 1018n / cm-3

1 x 1018

ωp

Ω+

Ω-

Wav

enum

ber

/ cm

-1

n1/2 x 10-8 / cm3/2

Figure 3.6: Doping dependence of the eigenfrequencies of the coupled plasmon-LO-phonon modes Ω+ and Ω− of n-doped GaAs for q = 0. The wavenumbers are calculatedwith the formula 3.8.

29

samples can be found elsewhere [Bau96]. To study the effects of chalcogens onclean GaAs(100) surfaces in situ, it is necessary to combine the Ramanspectrometer with an UHV chamber. The chamber contains several Knudsencells for the evaporation of chalcogens and a manipulator with an heater toprepare the GaAs(100) samples. A schematic viewgraph of the experimental set-up at the TU Chemnitz can be found elsewhere [Emd96].There are several features in the Raman spectrum which can be affected by bandbending:• The LO intensity due to deformation-potential scattering from the depletion

layer.• The intensity of the coupled plasmon-LO-phonon (PLP) modes from the bulk

region.• The LO intensity due to electric field induced Raman scattering (EFIRS),

induced by the E-field in the depletion layer.The PLP modes occur only in doped semiconductors. They arise from acoupling between the LO phonon and the plasmon through their macroscopicelectric fields [Var65]. Due to this coupling two new eigenmodes Ω+ and Ω−occur, whose frequencies depend on the doping level [Bur67]. These frequenciesare mainly influenced by the LO (ωLO) and TO (ωTO) phonon and the plasmafrequency (ωp), which can be calculated according to formula 3.7.

*

0

2

m

enp ⋅⋅

⋅=∞εε

ω (3.7)

ωp: Plasma frequency n: Free-carrier concentration

e: Elementary charge

ε0: Dielectric constant

ε∞: Contribution of valence electrons to dielectric function m*: Effective mass

The eigenfrequencies Ω+ and Ω− of the PLP modes at q = 0 can be estimated byformula 3.8.

( ) ( ) ( )

−⋅⋅+−±+=Ω±

222222222 42

1TOLOpLOpLOp ωωωωωωω (3.8)

ωLO: LO phonon frequency

ωTO: TO phonon frequency

30

100 200 300 400 500 600

20

30

40

50

60a.)

Ω-

ωLO

Raman shift / cm-1

Cou

nts

/ arb

.uni

ts

20

30

40

50

60b.)

ωLO

Cou

nts

/ arb

. uni

ts

Figure 3.7: Raman spectra of clean GaAs(100) samples excited by the 488.0 nm laserline of an Ar+ laser with 60 mW laser power in backscattering geometry on a.) n-doped(ND = 1⋅1018 cm-3) and b.) p-doped (NA = 5⋅1018 cm-3) GaAs.

31

From the data in figure 3.6 it can be concluded that the doping concentrationmust be in a certain range to separate the frequencies of the PLP modes (Ω+ andΩ-) from the frequency of the LO phonon. This is given for the GaAs(100)samples, which were used in this study. The concentration of donators in the n-type GaAs samples amounts to 1⋅1018 cm-3. A typical Plasmon-LO-Phononspectrum is shown in figure 3.7.For p-doped GaAs(100) the PLP modes are not very pronounced because of thehigher damping which is caused by the higher effective mass of the holes (figure3.7). This is also reported by Cardona and Olego who did not find any Ω+ and Ω-

peaks at small wave vectors [Ole81]. Therefore the PLP modes are onlyapplicable to determine the band bending on n- type GaAs.The band bending on doped semiconductors is determined by the intensity ratioof the LO and the PLP modes. The PLP mode arises only in region with asufficient concentration of free charge carriers. Therefore it is necessary tochoose the correct laser line in sense of the penetration depth of the light. ThePLP modes do not derive from the depletion layer, which is the origin of the LOphonons.The width of the depletion layer depends on the band bending on the surface andthe concentration of charge carriers according to formula 3.9.

D

BBsD Ne

Vd

⋅⋅⋅⋅

= 02 εε(3.9)

dD: Thickness of depletion layer

εs: Static dielectric constant

ε0: Dielectric constant

VBB: Band bending

e: Elementary charge

ND: Donator concentration

From equation 3.9 the band bending can be calculated, if the thickness of thedepletion layer is known. This leads to equation 3.10:

( )0

2

2 εε ⋅⋅⋅⋅=s

DDBB

dNeV . (3.10)

To determine the width of the depletion layer from a Raman spectrum, oneneeds the intensity ratio of the LO phonon and the PLP mode (formula 3.11[Far87, Geu93]).

32

( )

( ) ( )[ ]12exp0

0−⋅⋅⋅

== DS

BBPLP

DPLO

PLP

DPLO d

VI

I

I

I α (3.11)

ILODP: LO intensity

IPLP: PLP intensity

ILODP(0): LO intensity of an undoped sample

IPLP(VBBS = 0): PLP intensity of a doped sample without

band bending α (E)= 1/dp: Attenuation coefficient (dp: Light penetration depth)

To determine the change in band bending due to the chalcogen modification of aclean GaAs(100) sample equation 3.11 is transformed to the two equations 3.12and 3.13. These equations represent the Raman intensity ratios before and afterthe chalcogen modification.

( )[ ]12exp0 −⋅⋅⋅= iDi dII α (3.12)

( )[ ]12exp0 −⋅⋅⋅= fDf dII α (3.13)

Ii: Quotient of LO and PLP intensity of the unmodified sample If: Quotient of LO and PLP intensity of the chalcogen modified sample

I0: Quotient of ILODP(0) and IPLP(VBB

S = 0)

Supposing the band bending of the clean GaAs(100) is determined by anothermethod (here: Photoemission), the width of the depletion layer on the cleansample can be calculated according to equation 3.9. Then the depletion layerwidth after the chalcogen modification can be estimated from the transformationof the quotient of equations 3.12 and 3.13. Due to this method the calibration ofI0 in equations 3.12 and 3.13 is not necessary. For the determination of theRaman intensity ratios it is only necessary that the conditions of the Ramanmeasurements before and after the chalcogen modification coincide. Then theband bending after the chalcogen modification can be calculated according toequation 3.14.

( )[ ]2

0

112expln2

1

2

+−⋅⋅⋅⋅

⋅⋅

⋅⋅⋅= i

Di

f

s

DfBB d

I

INeV α

αεε(3.14)

33

3.4 Transport measurementsCurrent-Voltage measurements are a useful method to determine the Schottkybarrier height of a metal-semiconductor contact. There are different

mechanisms, which can attribute to the transport of charge over the Schottkybarrier (figure 3.8).If the voltage V applied to a Schottky contact is bigger than 3kBT/e0, the currentI due to the thermoionic emission of charge carriers over the Schottky barriercan be described by formula 3.15.

Tk

Ve

BeII0

0 ⋅= (3.15)

with TkB

Bn

eTSAIΦ−

⋅⋅⋅= 2*0

(3.16)

and3

2*0* 4

h

kmeA Be ⋅⋅⋅⋅= π , effective Richardson constant (3.17)

e0: Elementary charge

kB: Boltzmann constant

T: Temperature

Figure 3.8: Band diagram to explain the different transport mechanisms over theSchottky barrier of a metal-semiconductor contact: Thermoionic emission over thebarrier (a), tunnelling through the barrier (b), recombination in the space charge region(c) and injection of minority charge carriers into the neutral part of the semiconductor(d). Black and white arrows represent electrons and holes, respectively.

34

S: Area of the Schottky contact

me*: Effective mass of electrons

h: Planck constant

It is found that the barrier height depends on the applied voltage, which islowered by the image force. This influence is represented by the so calledideality factor n. For high applied voltages the serial resistance RS of the metal-semiconductor contact mainly influences the current-voltage characteristic of theSchottky contact. That leads to the modified formula 3.18 for the description ofthe Schottky contact.

( )

Tkn

RIVe

B

S

eII ⋅⋅⋅−⋅

⋅=0

0(3.18)

The ideality factor for the current-voltage characteristic of the Schottky barriershould not exceed the value 1.1. For small barrier heights the current-voltagecharacteristic shows an ohmic behaviour, because high currents flow even at lowvoltages. This can be excluded by doing the measurements at low temperatures,

-0,4 -0,2 0,0 0,2 0,4 0,6 0,8 1,00,0001

0,001

0,01

0,1

1

10

100

1000

I0

I / 1

0-6 A

U / V

Figure 3.9: I-V characteristic of a real diode. I0 is the intersection of the extrapolatedfitted line at a voltage of 0 V, which is used to determine the Schottky barrier height.

35

which lowers the current at low voltages so that the quasi ohmic behaviour iscancelled.Evaluating the other contributions to the current-voltage characteristic of theSchottky barrier leads to the following conclusions. The tunnelling (figure 3.8b.) through the barrier is negligible in this study, because the measurements arecarried out at room temperature. The temperature limit at which the tunnellinggets sensible amounts to 255 K. This limit is calculated by applying the dopingconcentration of the used GaAs samples (ND = 1018 cm-3), the effective mass ofelectrons in GaAs (me

* = 0.063 ⋅ m0) and the dielectric permitivity of GaAs (εb =10.9) after a formula given by Kampen [Kam95b].The contribution of the recombination current in the space charge region (figure3.8 c.) is high for high barrier heights, materials with a low lifetime of thecharge carriers, low temperatures and low applied voltage [Kam95b]. Theseproperties do not apply for the realised measurements on metal-GaAs(100)Schottky contacts, so that this contribution is negligible.The relation between the current due to thermoionic emission (figure 3.8 a.) anddue to injection of minority carriers into the neutral part of the semiconductor(figure 3.8 d.) increases proportional to the temperature, the dopingconcentration and the barrier width and antiproportional to the barrier height[Kam95b]. Regarding the parameters of measurements, which have been carriedout in this study, it can be concluded that the thermoionic emission remains theonly process, which is relevant in the current-voltage characteristic of theSchottky barrier height.The measured I-V characteristics (figure 3.9) were fitted by an exponentialfunction according to formula 3.15. The fitting was done with an extra program.The Schottky barrier height and the ideality factor can be determined from theformulas 3.16 and 3.18.

⋅⋅⋅⋅=Φ0

2*

lnI

TSATkBBn (3.19)

( )I

V

Tk

en

B ln0

∂∂⋅

⋅= (3.20)

3.5 Low energy electron diffraction (LEED)The diffraction or elastic scattering of electrons is a standard technique to obtainstructural information about surfaces. The accelerating energy of the electrons isin the range of 20 - 200 eV, which corresponds to a wavelength of 0.27 – 0.09nm. The wavelength is in the same range as the interatomic distances in a solid.Due to the low kinetic energy this method is also very surface sensitive (figure3.10). The elastically back scattered electrons give rise to diffraction or Braggspots which are imaged on a phosphorus screen. The difference compared to X-

36

ray scattering is, that the scattering vector component must equal a vector of thetwo dimensional reciprocal lattice G||. By relaxing the third Laue condition,which is valid for vectors perpendicular to the surface, the Ewald constructioncan be applied to the two dimensional problem (figure 3.10).Therefore the discrete reciprocal lattice points of the three dimensional problemare substituted by rods normal to the surface. The distance between two of theserods equals a reciprocal lattice vector. In a real LEED experiment the electronbeam penetrates several atomic layer and therefore the third Laue condition getmore and more valid with increasing penetration depth. This is indicated bythicker regions in the rod of the two dimensional Ewald construction in figure3.10. The observable scattered beams are now constructed by positioning theend of the wave vector, which represents the incident electron beam, to the (0,0)reciprocal lattice point. Around the starting point of this wave vector a Ewaldsphere with the radius |k| is constructed. Where the sphere crosses a twodimensional reciprocal lattice rod a scattered electron beam is observed. If thesphere crosses a thicker region of the rod a strong reflex is observed while onlya weak reflex is observed when crossing the thinner region. Therefore the strong

Figure 3.10: Ewald construction for elastic scattering on a quasi-twodimensional surfacelattice. Not only the scattering from the topmost layer but also from some underlyingplanes is taken into account. The thicker regions of the rods arise from the third Lauecondition, which cannot be completely neglected. Correspondingly the (20) reflex has ahigher intensity than the ( 20) reflex [Lüt95].

37

reflexes represent scattering from the bulk while the weak reflexes can beattributed to the surface.The LEED pattern must exhibit sharp spots with high contrast and lowbackground intensity. A broadening of the spots or an increase in backgroundintensity is a sign for random defects or crystallographic imperfections.Therefore the LEED techniques is an easy and powerful method to check thesurface reconstruction of semiconductor surfaces and to get a hint on theordering and in some cases also on the contamination of the surface.

38

39

4 Characterisation of clean and modified GaAs surfaces

4.1 Clean GaAs(100) surfacesFor the investigation of modified GaAs(100) surfaces a well understood surface

is needed as a starting point, which is identified by reconstruction. The geometryof the surface is determined by the reconstruction and with the geometry also the

47 46 45 44 43 42 41 40 39

a.)

As6 As5

As4

Inte

nsity

/ a

rb.

units

Binding energy / eV

b.)

Inte

nsity

/ a

rb.

units

As3

As1

As2

Figure 4.1: As 3d core level spectra excited with 79 eV photon energy: a.) As 3dspectrum of the As capped GaAs(100) surface before decapping; b.) after the decappingat 380 °C which exhibits a c(4x4) reconstruction.

40

electronic properties of the surface are determined. Another important point isthe easy preparation of the surface reconstruction. Therefore GaAs(100) waferswith a h1omoepitaxial n- and p-type GaAs(100) layer grown by molecular beamepitaxy (MBE) were taken for the experiments.

Figure 4.2: Model of the GaAs(100) c(4x4) reconstructed surface. Above is a view onthe top of the sample where the As dimers can be observed. The different sites of atomscorresponds to the shifted components in the figure 4.1 and 4.3. Below is view on fromthe side of the sample (from [Bie90, Fal92]).

41

These wafers are capped with a thick amorphous As layer which is immediatelydeposited after the growth of the homoepitaxial layer to protect the samplesfrom contamination.These As capped GaAs surfaces were examined by photoemission. To separate

the contributions of chemical shifted components from the bulk component themeasured photoemission spectra were fitted using spin-orbit split Voigt profilesapplying the method given by Joyce [Joy89]. In the decomposition of the corelevel spectra the spin-orbit splitting, the branching ratio and the Lorentzianwidth were always kept constant throughout the fits of the differently modifiedGaAs(100) surfaces. These values were derived from many series ofphotoemission measurements of clean GaAs(100) surfaces (table 4.1) and agreequite well with the values evaluated by others [Lar83, Lari94, Lay91,Vit92].The other parameters such as binding energy, intensity, background (Shirleybackground) and Gaussian width were variable during the fits.In figure 4.1 a. a photoemission spectrum of the As capped GaAs(100) surface isshown. This viewgraph exhibits three different contributions to the core levelemission. The As4 and As5 components are assigned to amorphous arsenic. Thebinding energy of the dominant As structure in the spectrum (As4) amounts to41.81 eV for As 3d5/2. The contribution of the As5 component is shifted by 0.60eV towards higher binding energy. The Gaussian width of the As6 component isapproximately 0.37 eV broader than that of the other two, which may indicatethat it does not have an homogeneous surrounding. It is shifted by 2.88 eVtowards higher binding energy. This agrees quite good with the binding energyshifts of As2O3, which amounts to approximately 3.0 eV related to the bulk peakof GaAs [Ber88, Lan84, Lu93, Mil 84, Pan97 and references therein].The As cap is removed from the surface of the samples by gentle annealing to380 °C in ultra high vacuum (UHV). This leads to a c(4x4) reconstructedsurface, what can be stated from the LEED pattern. The dots of the LEEDpattern were always very weak, which has also been reported by others [Dra78,Dus92].

As 3d Ga 3dLorentzian width /eV

0.1 0.1

Branching ratio 1.50 1.68Spin-orbit splitting /eV

0.44 0.69

Table 4.1: Fit parameters for the core level spectra, which are held fixed during thefitting procedure.

42

The c(4x4) reconstruction can also be stated from the shape of the As 3d corelevel spectrum, because of the high intensity on the higher binding energy side.The fitting of this spectrum shows three different components (figure 4.1 b.).The As1 component is assigned to bulk As while As2 and As3 are surfacecomponents. The binding energy of the bulk component amounts to 41.17 eV(all binding energies are given for the d5/2 components of the spin orbit split corelevels relative to the Fermi level), while the surface components are shifted by0.65 eV towards higher (As2) and 0.40 eV towards lower binding energy (As3).These shifts are in good agreement with the data given by [Lari94]. Accordingto the literature [Lar83, Lay91, Nea83, Vee84, Vit92] the As2 component isassigned to As in an pure As surrounding as it occurs for the As dimers on thec(4x4) reconstructed GaAs(100) surface. The As3 component is attributed to 3-fold coordinated As which is caused by the missing dimer rows of this surface[Lay91, Eas80]. This is illustrated in a model of the c(4x4) reconstructedGaAs(100) surface in figure 4.2.The Ga 3d core level spectrum of the decapped sample exhibits only a bulk(Ga1) and one surface component (Ga2) (figure 4.3). The binding energy of thebulk and surface component are given in table 4.2. An explanation for the Gasurface component may be that it originates from Ga at the interface of the bulk

22 21 20 19 18 17

Ga2

Ga1

Inte

nsity

/ a

rb.

units

Binding energy / eVFigure 4.3: Ga 3d core level spectrum excited with 60 eV photon energy of the cleanAs-rich GaAs(100) after the decapping at 380 °C, which lead to a c(4x4) reconstructedsurface.

43

to the As covered surface where it is bound to the 3-fold coordinated As atomsin the second layer [Lay91].

The statistic values of all the samples (n- and p-type GaAs(100)) measured inthis study are given in table 4.2. These values coincide quite well with thevalues, which have been measured by others for the different binding energy

0 10 20 30 40 50 60 700,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

As1/Astot

As2/Astot

As3/Astot

a.)

Take off angle / °

Rel

. in

tens

ity

0,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

Ga1/Gatot

Ga2/Gatot

b.)

Rel

. int

ensi

ty

Figure 4.4: Intensities of the different contributions in the a.) As 3d and b.) Ga 3dphotoemission spectra relative to the total intensity of the measured photoelectrons.

44

shifts of the surface components for the c(4x4) reconstructed GaAs(100) surface[Lay91, Lari94, Lar83, Nea83].The assignment of the different components in the spectra of the cleanGaAs(100) surface is supported by the measurement of As 3d and Ga 3d corelevel spectra with increasing angles relative to the surface normal. For these

measurement the photoelectron analyser is turned away from the direction ofnormal emission. In figure 4.4 a. the change in intensity of the different Ascomponents relative to the cumulated emission of the As 3d core level as afunction of the take off angles is shown. Up to 30 ° the relative intensities arenot affected by the angle. At angles higher than 30 ° the relative intensity of theAs1 component decreases dramatically as it is expected for a contribution fromthe GaAs bulk. At the same time the relative intensity of the As2 componentincreases while the relative intensity of As3 remains mainly constant during thevariation of the angles. From this behaviour it is concluded that thephotoelectron emission of the As2 component originates from the first layer ofAs which consists of As dimers. Because the As3 component of the spectra doesnot show this extensive surface sensitivity, it is not assigned to the first layer ofthe GaAs(100) surface but to the second layer. This is in good agreement withthe assignment given above.In the case of the Ga 3d photoemission spectra (figure 4.4 b.) there is also asmall variation in the relative intensity up to 30 ° take off angle which can beexplained by the same argument as in the case of As. For higher take off anglesit is obvious that the Ga2 component exhibits an increasing relative intensity asa function of the take off angle while the Ga1 component decreases in intensity.This leads to the interpretation that the emission responsible for the Ga2

As1 As2 As3 Ga1 Ga2EBin(n-type) / eV

41.14 ±0.08

+ 0.62 ±0.03

− 0.50 ±0.04

19.26 ±0.07

+ 0.46 ±0.12

rel.Intensity(n-type)

− 0.37 ± 0.11 0.20 ± 0.04 − 0.07 ± 0.02

EBin(p-type) / eV

40.98 ±0.12

+ 0.65 ±0.04

− 0.48 ±0.07

19.10 ±0.10

+ 0.48 ±0.13

rel.Intensity(p-type)

− 0.33 ± 0.11 0.17 ± 0.05 − 0.07 ± 0.04

Table 4.2: Binding energies of As 3d5/2 and Ga 3d5/2 for the bulk components relative tothe Fermi level and binding energy shifts of the surface components. Additionally theintensities of the surface components relative to the bulk component are given.

45

component originates from Ga atom which are located nearer to the surface thanthe atoms responsible for the Ga1 photoemission contribution. That coincideswith the assignment of the components in figure 4.2.Another important point in the analysis of the clean GaAs(100) surface is thedetermination of the Fermi level with respect to the valence band maximum.

Therefore the valence band is measured on the GaAs(100) surface at 24 eVphoton energy (figure 4.5). In the next step a linear background is subtractedfrom the valence band spectrum, which is determined from the high energy tailof the spectrum. Thereafter the top of the valence band is fitted by a linear fit.The valence band maximum is determined from the intersection of the linear fitand the x=0 axis. Additionally the Fermi level of a clean peace of molybdenum,

which is in contact with the semiconductor, is measured at the same photonenergy. The position of the Fermi level is determined by the method, which has

22,0 22,5 23,0 23,5 24,0 24,5

Fermi level Valence band Fit for valence

band maximum

0,51 eV

0

Inte

nsity

/ ar

b. u

nits

Kinetic energy / eV

Figure 4.5: Evaluation of the Fermi level position relative to the valence band maximumon clean n-doped GaAs(100).

n-doped GaAs p-doped GaAsFermi level − VBM /

eV0.49 ± 0.07 0.34 ± 0.10

Table 4.3: Average Fermi level position relative to the valence band maximum (VBM)on n- and p-doped GaAs(100) from all samples examined in this study.

46

been described in chapter 3. As shown in figure 4.5 the difference in the kineticenergy of the Fermi level position and the valence band maximum represents theposition of the Fermi level in the band gap of the GaAs(100) surface withrespect to the valence band maximum.The statistic values of the Fermi level position are given in table 4.3. It can bestated that the position of the Fermi level on p-doped GaAs is generally closer tothe valence band maximum than the Fermi level of the n-doped GaAs samples.A comparison of the quotient of the integrated intensities of the As 3d and Ga 3dcore level emissions with the Fermi level position does not show any relation.Therefore it is concluded that the As to Ga ratio on the surface does not affectthe position of the Fermi level with respect to VBM. It is much more probablethat the Fermi level position depends strongly on the formation of kinks anddefects on the surface [Hir97a, Hir97b, Ish97].

4.2 Se modified GaAs(100) surfacesThe clean GaAs(100) surface, which was prepared as described in chapter 4.1, isnow modified by the evaporation of Se. To study the effect of Se on the As richsurface a small amount of Se was evaporated at room-temperature (RT) in a firststep.Each step of the sample modification was analysed by taking photoelectronspectra of the Ga 3d, As 3d and Se 3d core levels. The spectra were decomposedby curve fitting using Voigt functions for the different components.The clean surface in figure 4.6 a. exhibits the three As 3d components, whichhave been described in the last chapter. After evaporation of a small amount ofSe the As 3d photoemission spectrum changes dramatically. In figure 4.6 b. anew component As5 on the higher binding energy side is observed. Thiscomponent is shifted by + 1.14 eV. The shift of the As5 component to higherbinding energy compared to the GaAs bulk component agrees quite well withthe higher electronegativity difference between As and Se. The As5 componentis assigned to the formation of As2Se3 [Sci92, Tak92a, Tsu95]. This reactiontakes place between the excess As on the As-rich GaAs(100) surface and theevaporated Se. It is energetically favoured because it has a heat of formation of−103 kJ/mol, which is larger than the heat of formation of GaAs (−74 kJ/mol).Annealing the sample to 225 °C leads to the evaporation of the As2Se3 from thesurface of the sample, which has also been observed by others [Sci92, Tak92a].This can be concluded from the decreasing intensity of the As5 component(figure 4.6 c.). After annealing the sample to 380 °C the As5 componentdisappears (figure 4.6 d.). This is interpreted in the way that the As2Se3 isvolatile at temperatures between 225 °C and 380 °C. Additionally As from theAs dimers vanishes from the surface which can be concluded from a decrease inintensity of the As2 component in the photoemission spectrum.

47

c.) f.)

Inte

nsity

/ ar

b. u

nits

As5

As2 As1

b.) e.)

x 4

45 44 43 42 41 40 39

As2

As1

As3

Binding energy / eV

a.)

45 44 43 42 41 40 39

Binding energy / eV

d.)

Figure 4.6: As 3d core level spectra excited with 74 eV photon energy of a.) the cleanc(4x4) reconstructed GaAs(100) surface after decapping, b.) after evaporation of a smallamount of Se, c.) after annealing to 300 °C, d.) after annealing to 480 °C, e.) afterevaporation of a great amount of Se and f.) after annealing to 480 °C.

48

c.) f.)

Ga4

Ga3

Ga1

b.)

Inte

nsity

/ ar

b. u

nits

x 4

e.)

22 21 20 19 18 17

Ga2 Ga1

a.)

Binding energy / eV

22 21 20 19 18 17

d.)

Binding energy / eV

Figure 4.7: Ga 3d core level spectra excited with 51 eV photon energy of a.) the cleanc(4x4) reconstructed GaAs(100) surface after decapping, b.) after evaporation of a smallamount of Se, c.) after annealing to 300 °C, d.) after annealing to 480 °C, e.) afterevaporation of a great amount of Se and f.) after annealing to 480 °C.

49

c.)

b.)

Inte

nsity

/ ar

b. u

nits

e.)

58 57 56 55 54 53 52 51

Se1Se2

a.)

Binding energy / eV

58 57 56 55 54 53 52 51

/ 4

d.)

Binding energy / eV

Figure 4.8: Se 3d core level spectra excited with 88 eV photon energy of a.) theGaAs(100) surface after evaporation of a small amount of Se, b.) after annealing to 300°C, c.) after annealing to 480 °C, d.) after evaporation of a great amount of Se and e.)after annealing to 480 °C.

50

After the second Se treatment only two components are observed which areassigned to As in As2Se3 on the surface (As5) and in the bulk (As1) of the GaAs(figure 4.6 e.).Any contribution of the As dimers has vanished. The subsequent annealing to380 °C removes nearly all the arsenic selenide on the sample surface and onlythe As bulk component remains in the spectrum (figure 4.6 f.).

Figure 4.9: Model of the Se modified GaAs(100) (2x1) reconstructed surface. Above is aview on the top of the sample where the Se dimers can be observed. The different sitesof atoms corresponds to the shifted components in figures 4.6, 4.7 and 4.8. Below is aside view of the sample (from [Pas94]).

51

The Ga 3d spectra are interpreted correspondingly. The clean GaAs(100) surfaceexhibits with exception of the bulk component only one additional component(figure 4.7 a.). Upon deposition of a small amount of selenium onto the surfacetwo new Ga components Ga3 and Ga4 are observed (figure 4.7 b.). Both areshifted to higher binding energy by 0.40 eV (Ga3) and 1.00 eV (Ga4),respectively. This is in the same range as already observed by others [Cha90].The Ga3 component can be attributed to Ga bonded to Se on the surface and theGa4 component to subsurface Ga. The formation of a Ga2Se3 like layer on top ofthe sample is supported by the high heat of formation of – 409 kJ/mol forgallium selenide.The observed additional components in the Ga 3d and Se 3d (see below) spectrado not correspond to the formation of GaSe, because GaSe exhibits only onecomponent in a photoemission spectrum per element [Amo98]. The formation ofa Ga2Se3 like layer on top of GaAs(100) due to the Se deposition is supported bymany other photoemission measurements [Mär95, Men91] as well as X-rayphotoemission diffraction (XPD) [Cha90], Auger electron spectroscopy (AES)[Izu96], electron energy loss spectroscopy (EELS) [Tu85] and extended X-rayabsorption fine structure (EXAFS) [Tak92b]. The annealing of the sample to225 °C does not affect the Ga 3d photoemission spectrum (figure 4.7 c.). Afterannealing to 380 °C the arsenic selenide is removed from the surface of thesample, which causes an increase of intensity for all Ga components (figure 4.7d.).Further deposition of Se onto the surface leads to a strong decreasing of theintensity of the Ga 3d core level photoemission (figure 4.7 e.). It also changes

the intensity ratio of the different components in the spectrum. The intensity ofthe Ga4 component increases relative to the bulk component Ga1. This is a signfor the continuation of the As Se exchange reaction on the surface. Thisinterpretation is also supported by the high heat of reaction for Ga2Se3, whichwas mentioned above. During this second step of selenium deposition thegallium selenide like layer is completed, because further deposition of seleniumonto this surface does not change the intensities of the different componentsanymore. Further annealing to 380 °C does not change the ratio of the differentGa components in the photoemission spectrum (figure 4.7 f.). Only the total

Se 3dLorentzian width / eV 0.1Branching ratio 1.50Spin-orbit splitting / eV 0.86

Table 4.4: Fit parameters for the deconvolution of the Se 3d core level spectra, which areheld fixed during the fitting procedure. The values correspond quite well with the valuesmeasured by others [Mae93, Sci92, Tak92a].

52

intensity of the whole Ga 3d photoemission signal increases, because the excessAs and the arsenic selenide is removed from the surface.In figure 4.8 a. the photoemission spectrum of Se 3d after deposition of a smallamount of Se is shown. Two different chemical shifted components can beobserved, which are separated by 0.81 eV. The estimated values for theLorentzian width, the branching ratio and the spin-orbit splitting of the d5/2 andd3/2 component are given in table 4.4. The two Se components Se1 and Se2 areattributed to surface and subsurface components, respectively, with the Se1component being the most dominant one. Annealing the sample to 225 °C(figure 4.8 b.) results in an exchange reaction of As and Se where seleniumdiffuses below the surface. This is concluded from the fact that both componentscontribute an equal intensity to the Se 3d photoemission. The chemicalsurrounding of both Se components gets more homogenous and as aconsequence the energy difference between the two components increases to0.91 eV. This value does not change at the following treatments of the sampleand it coincides quite well with the binding energy difference which has beenobserved earlier [Mär95]. By further annealing to 380 °C the Se2 componentincreases strongly and finally exceeds the intensity of the Se1 component (figure4.8 c.).

Figure 4.10: (2x1) LEED pattern of a GaAs(100) surface, which has been modified bySe at high temperatures, as a negative image. The kinetic energy of the primary electronbeam amounts to 54 eV.

53

The deposition of a large amount of selenium yields an amorphous seleniumlayer as revealed by a single Se3d component shifted to higher binding energy(figure 4.8 d.). This agrees with the fact that selenium is the element with thehighest electronegativity in this system. Annealing the sample to 380 °C resultsin a loss of the amorphous selenium layer on top of the GaAs(100) sample(figure 4.8 e.). The remaining selenium on the surface may be arranged indimers for stabilisation purposes. This surface is selenium saturated which canbe concluded from the fact that further evaporation and annealing results in thesame spectra and no further change can be observed.A model for the Se modified GaAs(100) surface has been proposed by Pashley[Pas94]. The components which are observed in the photoemission spectra areassigned to atoms on different sites in the model (figure 4.9). On top of thesurface is a layer which consists of Se dimers. This is concluded from STMmeasurements. The second layer consists of Ga atoms while Se forms the thirdlayer. That is explained by an exchange reaction between As and Se, which hasbeen observed by many others [Cai 92, Cha90, Cha91, Lid94, Sci 92, Sci93,Tak92a, Tu85]. In the forth layer vacancies occur in the Ga layer. Thesevacancies are proposed to satisfy the electron counting rule [Pas89]. Thesevacancies are also supported by theoretical calculations, because Ga vacanciesnear Se atom makes the system extremely stable [Nar97].The deposition of Se on the GaAs(100) surface at room temperature withsubsequent annealing does not lead to a well reconstructed surface. This isconcluded from the fact that the LEED pattern derived from this surface is notvery sharp. Therefore the GaAs(100) surface is modified at higher temperatureto promote the As-Se exchange reaction. The Se modification at temperaturesbetween 340 °C and 380 °C yields to a perfect (2x1) LEED pattern (figure 4.10).

This (2x1) reconstruction has also been observed by others for the Se modifiedGaAs(100) surface [Bie94, Mae93, Miw96, Tak91, Tak92a].Photoemission spectra of the GaAs(100) surface which is modified by Se athigher temperature are shown in figure 4.11.

EBinn / eV EBin

p / eVGa1 19.57 ± 0.08 19.24 ± 0.04Ga3 + 0.37 ± 0.04 + 0.39 ± 0.04As1 41.38 ± 0.09 41.14 ± 0.08Se2 + 0.92 ± 0.04 + 0.93 ± 0.03

Table 4.5: Binding energies and binding energy shifts of the different components in thephotoemission spectra of the As 3d5/2, Ga 3d5/2 and Se 3d5/2 core levels for theGaAs(100) surface, which is modified by Se at elevated temperatures.

54

44 43 42 41 40 39

As1

a.)

Inte

nsity

/ ar

b. u

nits

Binding energy / eV

22 21 20 19 18

Ga3

Ga1b.)

Inte

nsity

/ ar

b. u

nits

58 57 56 55 54 53

Se2Se1

c.)

Inte

nsity

/ ar

b. u

nits

Figure 4.11: Photoemission spectra of the Se modified GaAs(100) surface at highertemperature of a.) the As 3d core level excited with 79 eV photon energy, b.) the Ga 3dcore level excited with 60 eV photon energy and c.) Se 3d excited with 88 eV photonenergy.

55

Comparing the spectrum of the As 3d core level measured on the GaAs(100)surface, which is modified by Se at room temperature with subsequent annealing(figure 4.6 f.) and the As 3d core level Se modified at higher temperature (figure4.11 a.) shows that the As4 component does not occur on the latter. This iscaused by the fact that the surface temperature lies above the temperature whenthe arsenic selenides evaporate from the Se modified GaAs(100) surface. SomeAs 3d spectra of the Se modified surfaces at higher temperature show also anadditional component at lower binding energy. This may originate from the Asatoms below the surface, which are located next to the Ga vacancies (figure4.9).The Ga 3d photoemission spectrum of the Se modified GaAs(100) surface atelevated temperature exhibits beside the bulk component only one additionalcomponent Ga3 (figure 4.11 b.). A second additional component Ga4 at higherbinding energy was found on some samples which were modified at highertemperature but always the intensity was not sufficient to establish a newcomponent in the fit of the photoemission spectra. A reason for this may be thechanging resolution of the synchrotron beamline. This is a difference to theGaAs(100) surface which was Se modified at room temperature andsubsequently annealed (figure 4.7 f.).The photoemission spectrum of the Se 3d core level of the Se modifiedGaAs(100) surface at elevated temperature (figure 4.11 c.) does not differ verymuch from the spectrum of the surface which was Se modified at roomtemperature and subsequently annealed (figure 4.8 e.). The only differenceoccurs in a slightly higher binding energy difference between the two Secomponents (table 4.5). The difference may be caused by the differenthomogeneity of the differently prepared GaAs(100) surfaces.To check whether the reconstruction of the Se modified GaAs(100) surface atelevated temperature is reproducible, the total intensities of the Ga 3d and As 3dphotoemission spectra after the Se modification were estimated. Thereafter therelation between the As 3d and Ga 3d intensity was calculated for all samples,which were Se modified. It was found that this relation is very stable over allsamples (table 4.6) independent of the type of doping.

The position of the Fermi level relative to the valence band maximum can bedetermined from the shift of the bulk components of the As 3d and Ga 3d core

Se modified n-GaAs(100)

Se modified p-GaAs(100)

IAs / IGa 0.30 ± 0.04 0.31 ± 0.02

Table 4.6: Intensity relation of the total photoemission intensity of the Ga 3d and As 3dcore level spectra after the Se modification of the GaAs(100) surface at elevatedtemperature.

56

levels as described in chapter 3. The values of the Fermi level position of theclean and Se modified sample on n- and p-doped GaAs(100) are given in table4.7. The shift of the Fermi level amounts + 0.31 eV and + 0.15 eV on n- and p-doped GaAs, respectively. Therefore the band bending on n-doped GaAs isreduced, while it is increased on p-doped GaAs due to the Se modification.

That leads to the interpretation that the intrinsic defects on the GaAs(100)surface such as antisites or kinks, which originally pin the Fermi level, areremoved by the Se modification. The Se modification forms new acceptor- anddonor-type states in the bandgap, which now pin the Fermi level at 0.80 eV and0.49 eV on n- and p-doped GaAs, respectively. That is in contrast to the ideathat Se passivates the GaAs surface by means of a removal of any kind of statesfrom the bandgap.Another interesting value, which can be determined from the energeticdistribution curves (EDC) measured by photoemission, is the change inionisation energy. The Se modification of the GaAs surface causes an averagechange in ionisation energy of (+ 1.15 ± 0.15) eV. This increase can beexplained by an adatom-induced surface dipole build by the Se-Ga bond[Mön95].The change in ionisation energy can be estimated theoretically. The Se-Ga bondon the Se modified GaAs surface exhibits a dipole moment, because both atomshave different electronegativities. The charge transfer ∆qi between the Ga andthe Se atom can be calculated according to formula 4.1.

2035.016.0 SeGaSeGai XXXXq −⋅+−⋅=∆ (4.1)

With the electronegativities of Ga (XGa = 1.81 [Pau60]) and Se (XSe = 2.55[Pau60]) the charge transfer ∆qi amounts to 0.138. By using a simple pointcharge model the dipole moment perpendicular to the surface may be written as

n-doped p-dopedEF – VBM of cleanGaAs(100) / eV 0.49 ± 0.07 0.34 ± 0.10

EF – VBM of Semodified GaAs(100) /eV

0.80 ± 0.14 0.49 ± 0.11

Table 4.7: Position of the Fermi level above the valence band maximum (VBM) on then- and p-doped GaAs(100) surface before and after the Se modification of the hotsurface determined by photoemission spectroscopy from the shift of the bulkcomponents of the As 3d and Ga 3d core levels.

57

⊥⊥ ⋅∆⋅= iii dqep 0 . (4.2)

The dipole moment normal to the surface can be calculated according to formula4.2. The distance between the Ga and Se atom normal to the surface is regardedto be quite similar to the distance of the Ga-S bond normal to the surface, whichamounts to 0.11 nm [Hir98, Sug97]. Using this value and the charge transfer ∆qi

from Ga to Se of 0.138 leads to a value of 2.43⋅10-30 Cm for the dipole momentnormal to the surface for the Ga-Se dipole.The change in the surface barrier height due to these dipoles can be calculatedaccording to the equation 4.3 given by Topping [Top27].

ad

ad

i NN

peI ⋅

+⋅±=∆ ⊥

2/30

0

91 αε (4.3)

With 2.43⋅10-30 Cm for the dipole moment of the Ga-Se dipole normal to thesurface, the elementary charge e0, the dielectric constant in vacuum ε0, 3.77⋅10-24

cm3 for the polarizability α of Se and 6.26⋅1014 cm-2 for the density of thedipoles at the surface, which coincides with the density of atoms on theGaAs(100) surface [Mön95], the change in ionisation energy amounts to 1.12eV. Because the Se atom of the surface dipole is negatively charged due to itshigher electronegativity while the Ga atom is positively charged, the ionisationenergy should be increased. This is in excellent agreement with theexperimentally derived value and thus supports the proposed occurrence of aSe-Ga double layer on top of the proposed model (figure 4.9).

58

100 200 300 400 500 600

15

25

35

45

55

Ω-

ωLO

a.)

Raman shift / cm-1

coun

ts /

a.u.

coun

ts /

a.u.

15

25

35

45

55

Ω-

ωLO

b.)

Figure 4.12: Raman spectra with an incident laser line of 488.0 nm and 60 mW in thegeometry (–100)[100;010](100) of the a.) clean and b.) Se modified GaAs(100) surface.

59

On n-type GaAs the change in band bending can also be observed by Ramanspectroscopy as described in chapter 3. The sample is excited by an Ar+ laserwith the 488.0 nm laser line with a power of light which amounts to 60 mW.The Raman spectrum which is recorded from the clean GaAs surface is shownin figure 4.12 a. Thereafter the GaAs(100) surface is modified by theevaporation of Se, while the surface is heated to 380 °C. A second Raman

spectrum is measured on this modified surface (figure 4.12 b.). The Ramanspectra are fitted by two Lorentzian peaks. The values derived from these fits areshown in table 4.8.The width of the depletion layer dD

i is directly related to the band bending onthe surface (formula 3.9). The band bending e⋅ VBB can be estimated from theband gap Egap and the position of the Fermi level relative to the valence band EF

− EVBM according to equation 4.4. This equation is valid for highly dopedsemiconductors, where k reveals the Boltzman constant, T the temperature, ND

the concentration of donators and NC the effective density of states in theconduction band. The last term in formula can be neglected, because theconcentration of donators in the investigated GaAs(100) sample (ND = 1⋅1018

cm-3) is much bigger than the effective density of states in the GaAs conductionband (NC = 4.7⋅1017 cm-3 [Lan82]). Therefore the Fermi level coincides with theconduction band minimum (CBM).

( )

−−−=⋅

D

CVBMFgapBB N

NkTEEEVe ln (4.4)

Intensity Raman shift /cm-1 Width / cm-1

ωLO of cleanGaAs(100)

294 291.3 5.5

Ω- of cleanGaAs(100)

435 272.0 18.7

ωLO of SemodifiedGaAs(100)

337 289.8 7.2

Ω- of SemodifiedGaAs(100)

787 270.9 16.6

Table 4.8: Parameters of the Lorentzian peaks, which are used to fit the Raman spectraof the clean and Se modified GaAs(100) surface.

60

The band gap Egap of GaAs at room temperature amounts 1.42 eV [Lan82].Together with the value for EF − EVBM (0.49 eV) for the clean n-GaAs(100)sample from table 4.7 the width of the depletion layer dD

i can be calculated to36.69 nm (formula 4.5).

D

BBsD Ne

Vd

⋅⋅⋅⋅

= 02 εε(4.5)

From the relation of the intensities of the LO and Ω- Raman peak before andafter the Se modification the change in band bending can be calculated.Therefore formula 4.6 was derived in chapter 3.

( )[ ]2

0

112expln2

1

2

+−⋅⋅⋅⋅

⋅⋅

⋅⋅⋅= i

Di

f

s

DfBB d

I

INeV α

αεε(4.6)

The change in band bending was calculated by using the following values:elementary charge e, doping density ND (1⋅1018 cm-3), static dielectric constant εs

for GaAs (13.1 [Bes98]), dielectric constant ε0 of vacuum, reciprocal of thelight penetration depth in GaAs α (1.1174⋅107 m-1 [Asp83]) for laser light with awave length of 488.0 nm, depletion layer width of the clean GaAs(100) surfacedD

i (36.69 nm), which can be calculated from the value for the band bending outof table 4.7 and the relations of the intensities of the LO and Ω- Raman peaksbefore Ii and after the Se modification If from table 4.8. Consequently the bandbending on the Se modified GaAs(100) sample amounts to 0.48 V. This impliesa reduction in band bending relative to the clean sample of 0.45 V, which isslightly higher than the value 0.31 V determined by photoemission. A reason forthis might be the difficulty to determine the right intensities of the LO and PLPmodes in the Raman spectra.In conclusion it seems that the Se treatment of the GaAs(100) surface leads to anew surface state distribution within the gap rather than a removal of surfacestates from the gap.

4.3 S modified GaAs(100) surfacesAnother chalcogen, whose influence on the GaAs(100) surface is investigated inthis study, is sulphur (S). Therefore the As rich GaAs surface is covered withincreasing amounts of S at room temperature from a Knudsen cell like oven.Thereafter the S modified GaAs(100) sample is annealed at increasingtemperature. The photoemission measurements were carried out to detectchemical reactions that occur at the GaAs(100) surface upon the deposition of Sand as a function of annealing temperature. The photoemission spectra weretaken after every step of modification. These spectra are fitted using Voigt

61

profiles for the different components, which contribute to the photoemissionsignal.Figure 4.13 a. shows the As 3d core level spectrum of the clean c(4x4)GaAs(100) surface containing three different components. As described inchapter 4.1 the As1 component is attributed to As in the four-fold coordinatedenvironment of the GaAs bulk and As2 and As3 to surface components. Theenergy shift amounts to 0.57 +/- 0.02 eV towards higher and to 0.49 +/- 0.02 eVtowards lower binding energy for As2 and As3, respectively. The interpretationof the surface components has already been done in chapter 4.1.Ten min of S deposition at RT changes the As 3d core level spectradramatically. In figure 4.13 b. the surface components As2 and As3 havedisappeared, whereas three new components (As4, As5, As6) occur in thespectrum. These components are distinct from the former surface components ascan be judged from the different binding energy shifts of 0.70 eV, 1.15 eV and1.63 eV towards higher binding energy for As4, As5 and As6, respectively.These values coincide quite well with the values of chemical shifted surfacecomponents, which have been observed by others due to the S modification ofthe GaAs(100) surface [Gei90, Mor94, Sak94, Yua97]. In some of thepublications the components between + 0.70 eV and + 1.1 eV relative to thebulk component are attributed to amorphous As on the surface, while thecomponent between + 1.6 eV and + 2.5 eV is attributed to the formation ofAs2S3. There were also components found with higher binding energy shifts,which are attributed to As2S5 or other arsenic sulphides (AsxSy). The existenceof the additional three components in the measured As 3d core level spectrum issupported by the spectrum displayed in figure 4.13 c., where S was evaporatedonto the sample for another 20 min. This additional S deposition increases theintensities of As5 and As6, whereas the intensity of As4 decreases slightly. Thissupports the conclusion that the deposition of S leads to the formation ofamorphous As on the GaAs(100) surface. This component is increased inintensity relative to the As2 component from the As dimers on the clean surface,which indicates that the S atoms also react with Ga in an exchange reaction withAs of the GaAs bulk. Additionally arsenic sulphides of different stochiometriesare formed, which are represented by the components As5 and As6. The shift ofthese components towards higher binding energy is in agreement with the highelectronegativity of S.

62

c.)

As4

Inte

nsity

/ ar

b. u

nits

As6

As5 As1

b.)

As7

e.)

45 44 43 42 41 40 39

As2

As1

As3

Binding energy / eV

a.)

45 44 43 42 41 40 39

Binding energy / eV

d.)

Figure 4.13: As 3d core level spectra excited with 79 eV photon energy a.) of the cleanAs-rich GaAs(100) surface after decapping at 380 °C, b.) after 10 min sulphurdeposition at room temperature, c.) after 30 min sulphur deposition, d.) after annealingto 180 °C and e.) after annealing to 480 °C.

63

The formation of these arsenic sulphides due to the sulphur modification of theGaAs surface by different treatments has also been observed before [Bes96b,But96, Gee96, Gei90, Kan96, Li94, Lu93, Mor94, Sak94, Sci91, Sug91,Yua97]. Subsequent annealing of the sample to 105 °C decreases the amount ofthe arsenic sulphides on the surface only slightly, whereas the As6 componenthas completely disappeared after annealing to 180 °C (figure 4.13 d.). Further

Figure 4.14: Model of the S modified GaAs(100) (2x1) reconstructed surface. Above is aview on the top of the sample, where the S dimers can be observed. The different sites ofthe atoms correspond to the components identified in the photoemission spectra from thesurface. Below is side view of the sample (from [Pas94]).

64

annealing to 480°C leads to the evaporation of the arsenic sulphides from thesurface, which has also been observed by Conrad et al. [Con97] and Xin et al.[Xin96].

22 21 20 19 18

Ga2

Ga1

a.)

Binding energy / eV

Ga4

Ga3

b.)

Inte

nsity

/ ar

b. u

nits

Figure 4.15: Ga 3d core level spectra excited with 60 eV photon energy a.) of the cleanGaAs(100) surface after decapping at 380 °C and b.) after 30 min evaporation of sulphurand annealing to 480 °C.

65

Therefore in figure 4.13 e.) only the bulk component (As1) and a component atlower binding energy (As7) are observed. The component As7, which is shiftedby 0.59 eV towards lower binding energy with respect to the As1 component, isassigned to As at the interface of the bulk GaAs to the gallium sulphide likelayer formed on top of the sample. The different components in the spectrum of

0,00,10,20,30,40,50,60,70,80,91,0

a.)

480 °C180 °C105 °C30 min Sclean

Ga1 Ga3 Ga4

Rel

ativ

e in

tens

ity

0 15 30 45 600,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7 b.)

Ga1 Ga3 Ga4

Rea

ltive

inte

nsity

Take off angle / °

Figure 4.16: Intensities of the different components of the Ga 3d core level spectrumrelative to the sum of the components: a.) as a function of the treatment of theGaAs(100) sample, b.) as a function of the angle of the photoelectron spectrometerrelative to the surface normal after 30 min S deposition and annealing to 480 °C.

66

the As 3d core level are attributed to As atoms at different chemicalenvironments. The different sites of the As atoms can be identified with the Asatoms in the model proposed by Pashley [Pas94] (figure 4.14). This model wasproposed for the (2x1) reconstructed GaAs(100) surface, which was modified bySe. But this model is also attributed to the S modified GaAs surface, because thephotoemission measurements lead to the same results as for the Se modifiedsurface.The Ga 3d spectrum in figure 4.15 a. of the clean GaAs(100) sample exhibitstwo different components. The Ga1 component originates from the GaAs bulk.The surface component (Ga2) is shifted by approximately 0.48 eV towardshigher binding energy. The interpretation of the surface component has beendescribed in chapter 4.1.The deposition of sulphur onto the GaAs(100) samples at RT causes a change inthe lineshape of the Ga 3d core level spectrum. From the curve fitting of thespectrum it is obvious that two new surface components (Ga3, Ga4) occur.These components are assigned to the formation of a gallium sulphide like layeron top of the sample. Both are shifted towards higher binding energy by 0.52 eV(Ga3) and 0.97 eV (Ga4), respectively.The formation of Ga-S bonds has also been observed before by different Smodification methods of the GaAs surface [Bes96b, Bes97a, Gei90, Hou96,Li94, Lu93, Mor94, Sak94, Sci 91, Sug91, Wan 96, Yua97]. The origin of theGa3 component is different to that of the Ga2 component of the clean surface,because it has a much higher intensity.In figure 4.16 a. the relative intensities of the Ga1, Ga3 and Ga4 component areshown as a function of the treatment of the sample. It is observed that therelative intensity of the Ga1 component decreases slightly due to the annealingwith increasing temperature. At the same time the Ga3 component increases inintensity with increasing annealing temperature, while the Ga4 componentremains mainly constant. This can be interpreted in the way that the Ga3component represents the formation of a gallium sulphide like layer, which isformed due to the increasing temperature which promotes the As-S exchangereaction. As this layer grows the intensity of the component from the GaAs bulk(Ga1) is more attenuated. The Ga4 component must originate from atoms at theinterface of the GaAs bulk to the gallium sulphide layer, because it remainsconstant during the different annealing cycles.In order to support the assumptions about the origin of these chemically shiftedcomponents angle resolved photoemission experiments were performed bymoving the analyser towards higher angles with respect to the surface normal.This increases the surface sensitivity of the photoemission measurements. It isobserved that the relative intensity of the bulk component (Ga1) decreases withincreasing angles, whereas the intensities of the chemically shifted componentsremain mainly constant (Ga4) or increase (Ga3) (figure 4.16 b.). Therefore it canbe concluded that the Ga3 component originates from Ga atoms located closer to

67

the surface than the Ga atoms responsible for the emission of Ga4. Applying the

163 162 161 160 159 158 157

S1 S2a.)

Binding energy / eV

b.)

Inte

nsity

/ ar

b. u

nits

Figure 4.17: S 2p core level spectra excited with 195 eV photon energy: a.) After 30 minsulphur deposition at room temperature and b.) after subsequent annealing at 480 °C.

68

model in figure 4.14 to the sulphur modified GaAs(100) surface the Ga3component is assigned to Ga atoms completely surrounded by sulphur, whereasGa4 originates from Ga atoms at the interface of this gallium sulphide like layerto the GaAs bulk which are bound to As and S atoms.

The lineshape of the S2p core level spectra (figure 4.17 a.) of the GaAs(100)surface which was exposed to an atomic sulphur beam at RT for 30 minindicates at least two different components. Due to the poor resolution of thebeamline at 195 eV photon energy (∆E ≈ 1 eV) the existence of morecomponents cannot be excluded. To analyse the different components in the S2p core level, it is fitted by two spin orbit split Voigt profiles. The fixedparameters used in the fitting process are given in table 4.9.

The two S components are separated by 0.69 eV. The S1 component is assignedto sulphur on top of the sample in an amorphous S environment and S2 toincorporated sulphur, which was exchanged with As in the top layer of the GaAssample [Con97]. Annealing the sample to 480°C enforces this S-As exchangereaction and leads to the reevaporation of sulphur from the surface, which can beconcluded from the increase of the S2 component relative to S1 (figure 4.17 b.).The separation in binding energy of the two S components is increased to 0.94

S 2pLorentzian width / eV 0.1Branching ratio 2.00Spin-orbit splitting / eV 1.18

Table 4.9: Fit parameters for the S 2p core level spectra, which are held fixed during thefitting procedure. The values correspond quite well with the values measured by others[Mor94].

EBinn / eV EBin

p / eVGa1 19.63 ± 0.12 19.24 ± 0.07Ga3 + 0.45 ± 0.02 + 0.48 ± 0.02Ga4 + 0.90 ± 0.04 + 0.93 ± 0.03As1 41.40 ± 0.09 41.08 ± 0.08As7 − 0.69 ± 0.07 − 0.76 ± 0.08S2 + 0.92 ± 0.06 + 0.94 ± 0.03

Table 4.10: Binding energies and binding energy shifts of the different components inthe photoemission spectra of the As 3d5/2, Ga 3d5/2 and S 2p3/2 core levels for theGaAs(100) surface, which is modified by sulphur at elevated temperatures.

69

eV, which can be interpreted that the chemical surrounding of the sulphur at thesurface has changed.After the annealing the surface also exhibits a weak (2x1) LEED pattern, whichis a hint for the formation of S dimers on top of the surface. The S1 componentis thus attributed to this surface dimers in accordance with the assignment of thetwo Se components observed for Se modified GaAs(100) surface (chapter 4.2).In order to achieve a better surface reconstruction the S modification of theGaAs(100) surface is carried out at 480 °C and not at room temperature withsubsequent annealing. The LEED pattern, which is measured at this surface(figure 4.18) is not as sharp as the one which has been measured on the Semodified GaAs(100) surfaces. This may be related to the fact that the Se atommatches the size of the As atom better than the S atom and therefore theexchange of As and S in the GaAs lattice induces some stress, which causes aworse long range order on the surface. But nevertheless the S modifiedGaAs(100) surface exhibits (2x1) reconstruction as has been observed before[Con96, Mor94].

Figure 4.18: (2x1) LEED pattern of a GaAs(100) surface, which has been modified bysulphur at high temperatures, as a negative image. The kinetic energy of the primaryelectron beam amounts to 62 eV.

70

44 43 42 41 40 39

As7As1

a.)

Inte

nsity

/ ar

b. u

nits

Binding energy / eV

22 21 20 19 18

Ga4

Ga3

Ga1b.)

Inte

nsity

/ ar

b. u

nits

163 162 161 160 159 158 157

S2S1

c.)

Inte

nsity

/ ar

b. u

nits

Figure 4.19: Photoemission spectra of the S modified GaAs(100) surface at highertemperature (> 400 °C): a.) As 3d core level excited with 79 eV photon energy, b.) Ga3d core level excited with 60 eV photon energy, c.) S 2p excited with 195 eV photonenergy.

71

In figure 4.19 photoemission spectra of the S modification of the GaAs(100)surface at high temperature (> 400 °C) are shown. For the As 3d spectrum(figure 4.19 a.) no difference is observed between the hot S modification and theS modification at room temperature and subsequent annealing.The same result is observed for the S 2p photoemission spectrum (figure 4.19c.), but the components of the hot S modified GaAs surface exhibit a Gaussianbroadening diminished by approximately 0.25 eV. This is interpreted in the waythat the S modification at higher temperature (> 400 °C) leads to a higherhomogeneity of the GaAs surface.

In comparison with the sulphur modification at RT and subsequent annealing thesulphur treatment at elevated temperature results only in a higher relativeintensity of the Ga3 and Ga4 components (figure 4.19 b.). This is related to thefact that sulphur exchanges with arsenic to a higher extent due to the highertemperature.

The sulphur modification of the GaAs(100) surface at high temperature alwayslead to the same surface reconstruction, which can be stated from the LEED andphotoemission measurements. The latter is supported by the binding energiesand the shifts in binding energy (table 4.10), which always exhibited the samevalues.

S modified n-GaAs(100)

S modified p-GaAs(100)

IAs / IGa 0.25 ± 0.04 0.24 ± 0.01

Table 4.11: Intensity relation of the total photoemission intensity of the Ga 3d and As 3dcore level spectra after the sulphur modification of the GaAs(100) surface at elevatedtemperature (> 400 °C).

n-doped p-dopedEF – VBM of cleanGaAs(100) / eV 0.49 ± 0.07 0.34 ± 0.10

EF – VBM of Semodified GaAs(100) /eV

0.90 ± 0.09 0.51 ± 0.17

Table 4.12: Position of the Fermi level above the valence band maximum (VBM) on then- and p-doped GaAs(100) surface before and after the Se modification of the hotsurface determined by photoemission spectroscopy from the shift of the bulkcomponents of the As 3d and Ga 3d core levels.

72

The good reproduction of the surface reconstruction due to the sulphurmodification is also supported by the constant ratio of the quotient of the totalintensities of the As 3d and Ga 3d core level spectra (table 4.11) of all samples,which were investigated. This is comparable to the observation made for theselenium modified GaAs(100) surfaces described in chapter 4.2.The position of the Fermi level relative to the valence band maximum isdetermined from the binding energy of the bulk components (As1, Ga1) of theAs 3d and Ga 3d core level spectra (table 4.12). From these values the shifts ofthe Fermi level due to the sulphur modification of the GaAs(100) surface can beestimated to + 0.41 and + 0.17 eV for n- and p-doped GaAs, respectively. Hencethe band bending on n-GaAs(100) is reduced, while it is increased for p-GaAs(100) due to the sulphur modification.This observation is compatible to what was found for the selenium modificationof the GaAs(100) surface. The S modification has removed the intrinsic defectstates (e.g. antisites, kinks), which pin the Fermi level on the clean GaAssurface. New acceptor- and donor-type states are generated in the bandgap,which pin the Fermi level at 0.90 eV and 0.51 eV on n- and p-doped GaAs,respectively. It is concluded that the S modification similar to the Semodification rather leads to the formation of new states in the bandgap than to apassivation by means of a removal of states.Also interesting is the change in ionisation energy caused by the S modificationof the GaAs surface. This can be determined from the width of the energydistribution curves (EDC), which are measured by photoemission spectroscopy.The change ionisation energy ∆I due to the S modification of the GaAs surfaceis determined to (+ 1.26 ± 0.28) eV. According to Mönch [Mön95] ∆I can beexplained by an adatom-induced dipole moment on top of the surface.From the model for the S modified GaAs surface in figure 4.14 it is known thata S-Ga double layer covers the surface. The S-Ga bond represents a dipole,because of the different electronegativites of the participating atoms. The chargetransfer ∆qi between the Ga and the S atom is calculated according to equation4.7 [Han46].

2035.016.0 SGaSGai XXXXq −⋅+−⋅=∆ (4.7)

With the electronegativities of Ga (XGa = 1.81 [Pau60]) and S (XS = 2.58[Pau60]) the charge transfer ∆qi amounts to 0.144. By using a simple pointcharge model the dipole moment perpendicular to the surface may be written by

⊥⊥ ⋅∆⋅= iii dqep 0 . (4.8)

The distance between the S layer on top of the sample and the Ga atoms in thesecond layer perpendicular to the surface amounts to 0.11 nm [Hir98, Sug97].

73

Using this value and the charge transfer ∆qi from Ga to S of 0.144 leads to avalue of 2.54⋅10-30 Cm for the dipole moment normal to the surface.The change in ionisation energy due to the Ga-S surface dipoles can becalculated according to the formula given by Topping [Top27].

ad

ad

i NN

peI ⋅

+⋅±=∆ ⊥

2/30

0

91 αε (4.9)

With 2.54⋅10-30 Cm for the dipole moment of the Ga-S dipole normal to thesurface, the elementary charge e0, the dielectric constant in vacuum ε0, 2.90⋅10-24

cm3 for the polarizability α of S and 6.26⋅1014 cm-2 for the density of the dipolesat the surface, which coincides with the density of atoms on the GaAs(100)surface [Mön95], the change in ionisation energy ∆I amounts to 1.28 eV(formula 4.9). Because the S atom of the surface dipole is negatively chargeddue to its higher electronegativity while the Ga atom is positively charged, theionisation energy should be increased. This values excellently agrees with theexperimentally determined value thus supporting the occurrence of a S-Gadouble layer on top of the S modified GaAs surface according to the proposedmodel (figure 4.14).

The reduction in band bending is also determined by Raman spectroscopy. Theclean c(4x4) reconstructed GaAs(100) surface is excited by the 488.0 nm line ofan Ar+ laser with a power of 60 mW.

Intensity Raman shift /cm-1 Width / cm-1

ωLO of cleanGaAs(100)

74 287.7 5.7

Ω- of cleanGaAs(100)

75 269.2 19.4

ωLO of SmodifiedGaAs(100)

105 287.6 6.0

Ω- of S modifiedGaAs(100)

223 269.3 23.6

Table 4.13: Parameters of the Lorentzian peaks, which are used to fit the Raman spectraof the clean and S modified GaAs(100) surface.

74

100 200 300 400 500 600

15

20

25

30

Ω-

ωLO

a.)

Raman shift / cm-1

coun

ts /

a.u.

coun

ts /

a.u.

15

20

25

30

Ω-

ωLO

b.)

Figure 4.20: Raman spectra with an incident laser line of 488.0 nm and 60 mW in thegeometry (–100)[100;010](100) of the a.) clean and b.) S modified GaAs(100) surface.

75

The Raman spectra from the clean and the S modified GaAs(100) surface, whichwas prepared in the same way as for the photoemission measurements, areshown in figure 4.20. The peaks of the LO and PLP mode are fitted by aLorentzian lineshape and the values derived by the fitting process are shown intable 4.13.The band bending on semiconductor surfaces can be calculated from thequotient of the LO and Ω− Raman peak before and after the sulphurmodification, which was derived in chapter 3 (formula 4.10).

( )[ ]2

0

112expln2

1

2

+−⋅⋅⋅⋅

⋅⋅

⋅⋅⋅= i

Di

f

s

DfBB d

I

INeV α

αεε(4.10)

The width of the depletion layer of the clean surface (dDi = 36.69 nm) has

already been calculated in chapter 4.2. Additionally the elementary charge e, thedoping density ND (1⋅1018 cm-3) of the sample, the static dielectric constant εs forGaAs (13.1 [Bes98]), the dielectric constant ε0 of vacuum, the reciprocal of thelight penetration depth in GaAs α (1.1174⋅107 m-1 [Asp83]) for laser light with awave length of 488.0 nm and the relations of the intensities of the LO and Ω-

Raman peaks before Ii and after the S modification If from table 4.13 are neededto calculate the band bending on the S modified GaAs(100) surface. Thus theband bending after the S modification amounts to 0.31 V, which implies areduction of the band bending by 0.62 V, which is slightly higher than the value0.41 V determined by photoemission. The value determined by the Ramanmeasurements is a little bit higher (0.62 eV), which may be caused by thedifficulties in determining the exact intensities of the peaks measured for the LOand PLP modes.In conclusion the sulphur treatment of the GaAs(100) surface leads, similar tothe selenium modification, to a new surface state distribution within the bandgap rather than a removal of all surface states.

4.4 Te modified GaAs(100) surfacesTo complete the investigations of chalcogen modifications on GaAs(100)surfaces, the influence of tellurium on GaAs is studied in this chapter. Studies ofthe Te adsorption on GaAs were originally motivated by investigating theheteroepitaxial growth of ZnTe and CdTe on the GaAs(100) surface [Coh86,Fau86, Fel86, Mar86, Sri87, Tat87]. Especially the orientation of the CdTegrowth was found to depend on the initial stochiometry of the GaAs(100)surface. Te has also been observed to act as a surfactant in the growth of InAson GaAs(100) [Gra92, Miw98].The As-rich c(4x4) reconstructed GaAs(100) surface was used to examine theinfluence of Te on its chemical and electronic properties. The As 3d core levelspectrum of the clean GaAs surface exhibits two surface components (figure

76

4.21 a., which have been described already in chapter 4.1. The peak, which isattributed to the bulk occurs at 41.34 eV binding energy relative to the Fermilevel (As 3d5/2), while the surface components are shifted by + 0.64 eV (As2)and − 0.39 eV (As3). The evaporation of a small amount of Te leads to apronounced change in the shape of the core level (figure 4.21 b.). The fitting ofthe core level spectrum by Voigt profiles shows two As and two Te components.The fitting parameters of the Te 4d core levels, which are held constant for allTe 4d components are shown in table 4.14.

In figure 4.21 b. the component As3 is no more visible, which indicates that thec(4x4) surface reconstruction is destroyed. The As2 component, whichrepresents As dimers on the c(4x4) reconstructed GaAs, remains in the spectrumat the same binding energy shift as on the clean GaAs(100) surface. But it isincreased in intensity relative to the bulk component As1. This may be a sign foran exchange between arsenic and tellurium in the GaAs bulk, thus increasing theamount of arsenic on the surface. But it can also be related to the formation ofAs-Te bonds, which should occur at the same binding energy shift, because theelectronegativities of tellurium and arsenic differ only by 0.08. The bindingenergy of the Te1 component amounts to 40.93 eV, while the second Tecomponent Te2 occurs at – 0.63 eV. How the separation between thecontributions of As 3d and Te 4d was achieved will be described later. Afterannealing the sample to 400 °C the arsenic on top of the sample is removed,which can be concluded from the fact that the As2 component is no longervisible in the spectrum (figure 4.21 c.). The Te1 component decreases inintensity while the intensity of Te2 increases. Therefore it can be concluded thatthe Te1 component originates from tellurium adsorbed on the surface of thesample, while the Te2 component represents incorporated tellurium. The Te1component may also represent the As-Te bonds established at the surface andthese arsenic tellurides are evaporated from the surface of the sample due to theannealing. This coincides quite well with the observations made for the seleniumand sulphur modified GaAs(100) surfaces described in chapter 4.2 and 4.3. Theattribution of the Te components is also supported by the spectrum taken afterthe deposition of a high amount of Te onto this surface (figure 4.21 d.).

Te 4dLorentzian width / eV 0.1Branching ratio 1.41Spin-orbit splitting / eV 1.46

Table 4.14: Fit parameters for the deconvolution of the Te 4d core level spectra, whichare held fixed during the fitting procedure.

77

c.) f.)

Te2

Te1

Inte

nsity

/ ar

b. u

nits

As2

As1

b.) e.)

44 43 42 41 40 39 38

As2

As1

As3

Binding energy / eV

a.)

44 43 42 41 40 39 38

Binding energy / eV

d.)

Figure 4.21: As 3d and Te 4d core level spectra excited with 69.7 eV photon energy ofthe a.) clean GaAs(100) surface, b.) after evaporation of a small amount of tellurium, c.)after subsequent annealing to 400 °C, d.) after evaporation of a high amount oftellurium, e.) after subsequent annealing to 400 °C and f.) after evaporation of telluriumonto the sample surface at 400 °C.

78

After this modification the As1 component is nearly not visible anymore,because the surface is covered with a thick layer of Te.The Te1 component is also increased in comparison to Te2, which coincideswith its attribution to surface tellurium. Annealing this surface to 400 °C causesthe evaporation of the excess Te from the surface as can be stated from thedecreasing intensity of the Te1 component in figure 4.21 e. At the same time theintensity of the As1 bulk component is increased again compared to figure 4.21d., which supports the conclusion that the Te layer at the surface is diminished.Comparing the two Te components in intensity shows that they have nearlysimilar intensity after the annealing process. The 4d5/2 core level of the Te1component occurs at 40.94, while the Te2 component is shifted by 0.83 eVtowards lower binding energy and the binding energy of the 3d5/2 core level ofthe As1 component amounts to 41.36 eV, which coincides with the bindingenergy of the clean GaAs surface. At least the sample is annealed to 400 °C andsimultaneously exposed to tellurium. The spectrum which was measured afterthis modification is shown in figure 4.21 f. It can be stated that this spectrumdoes not differ very much from the spectrum in figure 4.21 e. Therefore it isconcluded that the surface is Te saturated. The binding energies of the As1 andTe1 components did not change. Only the binding energy shift of the Te2component increases to − 0.90 eV compared to the Te1 component, whichseems to be the saturation value.To verify the attribution of the different components in the photoemissionspectrum of As 3d and Te 4d core levels after the evaporation of tellurium onthe GaAs surface with subsequent annealing the spectra were measured withdifferent photon energies from 60 eV to 160 eV. In figure 4.22 a. to e. it isobserved that the intensity of the Te 4d contribution to the photoemissionintensity decreases with increasing photon energy. At 140.3 eV photon energyonly the As1 component is visible in the spectrum. The reason for the decreasingintensity of the Te 4d photoemission is the Cooper minimum in the cross sectionof the Te 4d core level. The theoretical values for the photoemission crosssections, which are shown in figure 4.23 a., are calculated by Yeh and Lindau[Yeh85]. Comparing these theoretical derived values with the experimentalderived relative photoemission intensities in figure 4.23 b. shows the qualitativesimilarity in the lineshape of the curves. Therefore the cross section dependenceof the Te 4d core level on the photon energy is a good proof to check thecorrectness of the fit of the As 3d and Te 4d core level spectrum.The Ga 3d core level spectrum of the clean surface exhibits only one bulkcomponent and one surface component, which coincides with the observationsmade in chapter 4.1. The surface component Ga2 is shifted by + 0.63 eV relativeto Ga1 (figure 4.24 a.).

79

c.) f.)

Inte

nsity

/ ar

b. u

nits

b.) e.)

44 43 42 41 40 39 38

Te2As1

Te1

Binding energy / eV

a.)

44 43 42 41 40 39 38

Binding energy / eV

d.)

Figure 4.22: As 3d and Te 4d core level spectra of the tellurium modified GaAs(100)surface and subsequent annealing excited with different photon energies: a.) 62.7 eV, b.)69.7 eV, c.) 79.1 eV, d.) 110.1 eV, e.) 120.2 eV, f.) 140.3 eV.

80

After the tellurium treatment of the GaAs sample and subsequent annealing thissurface component vanishes and a new component Ga3 occurs in thephotoemission spectrum. This new component is shifted by + 0.9 eV relative tothe bulk component. The shape of the Ga 3d core level spectrum does notchange very much during further modification steps of the GaAs sample. TheGa3 component can be attributed to Ga-Te bonds which have formed during theAs-Te exchange process. Ga-Te bonds have also been found by others, who

40 60 80 100 120 140 160

1

10

a.)

Tetheo

Astheo

Cro

ss s

ectio

n (M

b)

Photon energy (eV)

0,1

1b.)

Rel

. int

ensi

ty As1 Te1 Te2

Figure 4.23: a.) Cross section of the As 3d and Te 4d core level dependent on the photonenergy [Yeh85], b.) relative intensities of the As1. Te1 and Te2 components dependenton the photon energy.

81

stated this founding from a component in the Te 3d spectrum, where the Te-Gabond related components are shifted by - 0.8 eV to - 1.2 eV relative to the Te-Tebond related components [Fel86]. This value coincides quite well with 0.9 eV,which was found for the binding energy difference of the two Te components inthe Te 4d core level spectrum. The fact that tellurium forms a compound on the

GaAs(100) surface was also found by others [Rod95] and some people evenproposed the formation of a Ga2Te3 like layer [Oht98].Comparing the Te 4d core level spectrum with the Se 3d and S 2p spectrum itreveals also two components in the photoemission spectrum. The one at higher

22 21 20 19 18

Ga2

Ga1

Binding energy / eV

Inte

nsits

y /

arb.

uni

tsGa3

Ga1

Figure 4.24: Ga 3d core level spectrum at 78.7 eV photon energy of the a.) cleanGaAs(100) surface and b.) the Te modified GaAs(100) surface with subsequentannealing to 400 °C.

82

binding energy is attributed to chalcogen on the surface while the one at lowerbinding energy is dedicated to incorporated chalcogen below the surface. Incontrast to the selenium and sulphur modified GaAs surface, on which thecomponent at higher binding energy is more than double in intensity than thecomponent at lower binding energy, the component at higher binding energyequals the component at lower binding energy. This indicates that there are lesschalcogen atoms on the surface than on the surfaces modified by selenium andsulphur. Similar to the other chalcogen modifications it is also observed thattellurium bonds to arsenic after the evaporation at room temperature and thesearsenic tellurides evaporate from the surface after the annealing of the sample.In the Ga 3d spectrum of the Te modified GaAs(100) surface only oneadditional component which originates from the Ga-Te bond is observed, whichis in contrast to the two components which have been observed for the otherchalcogens. The observed component is also very small in intensity andtherefore it is concluded that the exchange of As and Te does not take place tothe same extent as on the Se and S modified surface.The determination of the band bending of the surface shows that the telluriummodification does not influence the band bending of the investigated n-dopedGaAs(100) sample, which is in contrast to the observations on selenium orsulphur modified n-GaAs samples. Chambers et al. [Cha91] explained thisinability to passivate with the low As-Te exchange in the third layer, which islimited to 10 %.

83

5 Characterisation of chalcogen modified metal/GaAs(100) contacts

5.1 In contacts on chalcogen modified GaAs(100) surfacesThe growth of In on clean GaAs(100) shows a perturbation of the GaAs surface,because it bonds to As or exchanges with Ga atoms on the surface [Mao92]. Thegrowth of thick In layers at room temperature leads to islanding, because the Inatoms exhibit a strong clustering [Mao92, Spic90]. Hence it is interesting toinvestigate the influence of the chalcogen modification of the GaAs surface onthe chemistry at the In/GaAs interface and the growth mode of In. Furthermoreit is interesting to examine the influence of chalcogens on the Schottky barrierformation for In on GaAs, which will be done in chapter 5.6.The Se and S modified GaAs samples were covered with In of increasingthickness in several steps measured by a quartz crystal microbalance in thevicinity of the sample. After each modification step the samples werecharacterised by taking photoelectron spectra of the As 3d, Ga 3d, Se 3d, S 2pand In 4d core levels. The core level emission spectra were analysed with regardto surface or chemically shifted components. Therefore each contribution in thecore level spectra including the contribution from the bulk was fitted with aspin-orbit split Voigt profile, i.e. a Lorentzian function convoluted with aGaussian function. Spectra of the clean surface and those obtained at largecoverages of In were employed to derive values for the Lorentzian width, thebranching ratio, the spin-orbit splitting and the asymmetric factor for theDoniach-Sunjic lineshape of the photoemission spectra from metallic

components [Doni70]. These parameters were then kept constant in the furthercurve fitting of the spectra (table 5.1). The spin-orbit splitting of the In 4d corelevel agrees well with the values found by others [Mao92, Spi90]. The otherparameters, such as binding energy, intensity, background (Shirley background)and Gaussian width are variables.

In 4dLorentzian width / eV 0.1Branching ratio 1.5Spin-orbit splitting / eV 0.86Asymmetry factor 0.138

Table 5.1: Fit-parameters held constant during fitting of the spectra for the differentmodification steps of the GaAs(100) sample and for all In components occurring in thespectra.

84

4.57 nm In

x4

c.) 3.93 nm In

x4

f.)

0.29 nm In

Inte

nsity

/ ar

b. u

nits As4

As1

b.) 0.5 nm Ine.)

44 43 42 41 40

As1

Binding energy / eV

a.)

44 43 42 41 40

S:GaAsSe:GaAs

As4

As1

Binding energy / eV

d.)

Figure 5.1: As 3d core level spectra at different modification steps from GaAs(100)surfaces excited with 79 eV photon energy: a.) After Se modification at 330 °C, b.) afterevaporation of In with a nominal coverage of 0.29 nm and c.) 4.57 nm, d.) after Smodification at 490 °C, e.) after evaporation of In with a nominal coverage of 0.50 nmand f.) 3.93 nm.

85

4.57 nm Inc.) 3.93 nm Inf.)

0.29 nm In

Inte

nsity

/ ar

b. u

nits

In1 In2In3

b.) 0.50 nm In

In3

In1

e.)

22 21 20 19 18 17 16

Ga3

Ga1

Binding energy / eV

a.)

22 21 20 19 18 17 16

S:GaAsSe:GaAs

Ga4

Ga3Ga1

Binding energy / eV

d.)

Figure 5.2: Ga 3d and In 4d core level spectra at different modification steps fromGaAs(100) surfaces excited with 60 eV photon energy: a.) After Se modification at330 °C, b.) after evaporation of In with a nominal coverage of 0.29 nm and c.) 4.57 nm,d.) after S modification at 490 °C, e.) after evaporation of In with a nominal coverage of0.50 nm and f.) 3.93 nm.

86

In the following the evolution of the core level spectra for n-type substrates isillustrated as an example since p-type samples revealed a very similar behaviour.

0.29 nm In

Inte

nsity

/ ar

b. u

nits

b.) 0.50 nm Ind.)

58 57 56 55 54 53 52 51

Se1Se2

Binding energy / eV

a.)

163 161 159 157

S:GaAsSe:GaAs

S1S2

Binding energy / eV

c.)

Figure 5.3: Se 3d and S 2p core level spectra at different modification steps fromGaAs(100) surfaces excited with 89 eV and 193 eV photon energy, respectively: a.)After Se modification at 330 °C, b.) after evaporation of In with a nominal coverage of0.29 nm, c.) after S modification at 490 °C, d.) after evaporation of In with a nominalcoverage of 0.50 nm.

87

The clean GaAs(100) surface is modified by the evaporation of chalcogens ontothe surface at elevated temperature. The core level spectra of As 3d after the Seand S modification are quite similar as seen in figure 5.1. a. and d.. The onlydifference occurs in the small component As4 on the S modified GaAs. Thiscomponent has not been observed for the Se modified surface, which may becaused by its lower intensity so that the spectrum can be sufficiently fitted withone As component (compare to chapter 4).The new components in the Ga3d core level spectra after the chalcogenmodification (figure 5.2 a. and d.) are assigned to Ga bonded to Se and S,respectively.The formation of a gallium selenide and gallium sulphide like layer is furthersupported by the analysis of the Se 3d and S 2p emission clearly revealing twocomponents (figure 5.3 a. and d.). In agreement with chapter 4 these componentsare assigned to chalcogen dimers on the surface (Se1 and S1) and subsurfacechalcogens (Se2 and S2) from the As-Se exchange reaction in the topmost GaAslayers.The formation of a gallium selenide and gallium sulphide like layer isfurther supported by the analysis of the Se 3d and S 2p emission clearlyrevealing two components (figure 5.3 a. and d.). In agreement with chapter 4these components are assigned to chalcogen dimers on the surface (Se1 and S1)and subsurface chalcogens (Se2 and S2) from the As-chalcogen exchangereaction in the topmost GaAs layers.Considering now the As 3d core level spectrum after a nominal In deposition of0.29 nm a new component (As4) appears in the photoemission spectrum fromthe Se modified GaAs(100) surface (figure 5.1 b.). It is shifted to lower bindingenergy by approximately 0.75 eV. In comparison with the As 3d core levelemission of the Se modified surface the Gaussian linewidth is further decreasedby approximately 60 meV as can already be seen in the spectrum from the morepronounced minimum between the spin-orbit split components. Therefore thenow apparent component has most probably only been obscured by the broaderemission of the Se treated surface. It can thus be supposed that it represents aninterface component of the topmost As layer near the Ga2Se3-like layer. This isalso supported by the fact that the intensity of the As4 component relative to theAs1 intensity stays constant with increasing In coverage. In the As 3dphotoemission spectrum of the S modified GaAs surface after evaporation of0.50 nm In a similar component is observed, which is shifted by 0.71 eVtowards lower binding energy (figure 5.1 e.). This coincides quite well with theshift of the additional As component observed on Se modified GaAs surface,which leads to the conclusion that it originates also from atoms at the interfaceof the Ga2S3 like layer on the top of the sample and the GaAs bulk.In the Ga 3d core level spectra on both chalcogen modified GaAs surfaces abroad feature appears at lower binding energy upon In deposition (figure 5.2 b.and e.). Due to the proximity of Ga 3d and In 4d binding energies it could

88

represent reacted Ga or In components. Clarification is obtained from anadditional experiment [Alo90, Cha93] allowing an unambiguous assignment.

For this purpose core level spectra were recorded at various photon energies inthe range between 60 and 190 eV. This causes a drastic change in the In 4dphotoionisation cross section due to the occurrence of a Cooper minimum[Yeh85]. As a result the In 4d contribution to the spectra decreases towardshigher photon energy to almost zero. Figure 5.4 a.) shows that the relativeintensity of the In component decreases, while the intensity of the Gacomponent slightly increases. The shape of the intensities agrees qualitativelywell to that what has been proposed by Yeh and Lindau [Yeh85] (figure 5.4 b.).Therefore it is concluded that the components named In1 and In2 in figure 5.2 b.

60 80 100 120 140 160 180 2000,01

0,1

a.)

Ga1 In3

Rel

. Int

ensi

ty /

arb.

uni

ts

Photon energy / eV

1

10

b.)C

ross

sec

tion

/ Mb

Ga In

Figure 5.4: a.) Relative intensities of the In3 and Ga1 component measured at differentphoton energies at a nominal In coverage of 0.5 nm on the S modified GaAs(100)surface. b.) Photoionisation cross section depending on the photon energy as calculatedby Yeh and Lindau [Yeh85].

89

and In1 in figure 5.2 e. stem from electron emissions of In. Altogether there arethree indium components separable in the spectra from the Se modifiedGaAs(100) surface. The In1 and In2 components are likely to be induced by Inat different adsorption sites on the Se modified GaAs surface. The bindingenergy of the In1 component amounts to 18.27 eV while the In2 component isshifted to higher binding energy by 0.43 eV. This coincides quite well with 18.8eV binding energy, which was found for the In 4d core level in the InSe bulk[Bro97]. On the S modified GaAs surface only one reacted component In1 isobserved at 18.8 eV binding energy. The missing In2 component is caused by aworse resolution in the photoemission measurement indicated by a Gaussianwidth, which is approximately 0.4 eV higher than for the In components in thespectra from the Se modified GaAs surface. Therefore there is no hint in thespectrum to fit the reacted In by more than one component. The assumption thatthe evaporated In does only react with the topmost Se layer on the sample issupported by the fact that a metallic Ga component is not observed in thespectra. This would be shifted between 0.7 and 0.9 eV towards higher bindingenergy relative to the component from the GaAs bulk depending on the size ofthe Ga clusters [Chen93]. The observation that the In does not dissolve thegallium chalcogenide layer is further supported by the comparison of the heatsof formation for the different compounds. The heat of formation of In2Se3

amounts to – 326.4 kJ/mol, which is much lower than – 408.8 kJ/mol for Ga2Se3

and the heat of formation of In2S3 amounts to – 355, 6 kJ/mol, which is muchlower that – 516.3 kJ/mol for Ga2S3.The In3 component (figure 5.2 b. and e.) is attributed to metallic In revealing anasymmetric lineshape and being much sharper than the reacted components. TheDoniac-Sunjic parameter which represents the metallic character amounts to0.138 which coincides quite well with the value of 0.1 found by Mao [Mao92].The metallic In3 component is shifted by approximately 1.5 eV and 1.6 eVtowards lower binding energy relative to the reacted In components for the Seand S modified GaAs surface, respectively. This shift of the metallic Incomponent towards lower binding energy has also been observed by others[Cha93, Chin87, Cim97]. With increasing In coverage the metallic componentbecomes more prominent and finally dominates the Ga 3d/In 4d spectra (figure5.2 c. and f.). The intensity ratios of the reacted Ga components to the Gacomponent from the GaAs bulk do not change with increasing In deposition.Moreover, the binding energy differences between all components in the spectraalso remain constant. Hence it is concluded that besides an initial reaction of Inwith the Se or S dimer atoms no further intermixing takes place at the interface.So the chalcogen modification inhibits a perturbation of the GaAs.The effect of the In deposition on the Se 3d and S 2p core level spectra isillustrated in figure 5.3. Due to the In deposition the intensity of the S1 and Se1component is decreased (figure 5.3 b. and d.). From this it can be concluded thatbesides an initial reaction of In probably with the S dimer atoms no further

90

intermixing takes place at the interface. The decrease of the S1 and Se1component is also compatible with a reaction of In with the S and Se dimeratoms, respectively.The growth mode of In on chalcogen modified GaAs(100) surfaces can beestimated from the intensities of the different components in the core levelspectra as a function of the In coverage. Figure 5.5 a. shows that all componentsof the As 3d, Ga 3d and In 4d core level spectra show the same decay with theexception of the In3 component. This supports the suggestion that In does notaffect the S modified GaAs surface, because otherwise the Ga3 and Ga4components, which originate from the gallium sulphide layer, should vanish dueto the dissolution of this layer. In contrast it is observed that these componentsdecrease exponentially with increasing In coverage similar to the Ga and Ascomponents from the GaAs bulk. The same is observed for the growth of In onthe Se modified GaAs surface, which leads to the same conclusion.In more detail it can be observed that at a nominal In coverage of approximately0.1 nm and higher the Ga3, Ga4 and As4 components which are dedicated to thegallium chalcogenide on the surface exhibit the same shape as the GaAs bulkcomponents Ga1 and As1 with increasing In coverage. This is also valid for theIn1 and In2 components. With the exception of In1 all components in figure 5.5a. and b. exhibit an exponential decrease, when the In coverage increases to 0.24and 0.29 nm and more on S and Se covered GaAs, respectively. In1 exhibits anexponential increase as a function of increasing In coverage. The calculation ofthe photoelectron attenuation length from the Ga1, Ga3, Ga4, As1 and In1components due to the growth of In on S modified GaAs from figure 5.5 a.between 0.24 and 1.00 nm nominal In coverage results to approximately 2.17nm. This is much higher than the escape depth of photoelectrons in GaAs in thesurface sensitive measurement mode at approximately 40 eV kinetic energy,which amounts to 0.5 nm (see chapter 3). Hence it is concluded that In formsislands after the initial adsorption of In atoms. The calculated attenuation lengthdecreases with increasing In coverage, which is interpreted by a strongerclustering of In atoms. This clustering has also been observed for the In growthon clean GaAs(100) surfaces [Spi90].The In growth on Se modified GaAs results in a photoelectron attenuation lengthof approximately 3.94. This value has been determined from the attenuation ofintensity of the Ga1, Ga3, As1, In1 and In2 components between a nominal Incoverage of 0.29 and 1.14 nm. With increasing In coverage the attenuationlength on Se modified GaAs shows the same behaviour as has been observed onS modified GaAs. But comparing these values absolutely for the S and Semodified GaAs surfaces it is striking that the attenuation length from the Semodified surface is much higher. This depends on the different mobilities of Inatoms on the S and Se modified GaAs surfaces. Hence the mobility of In atomson S modified GaAs is much higher than on Se modified GaAs, because the

91

0,0 0,5 1,0 1,5 2,0 4,0 4,5

In/S:GaAsa.)

As1 Ga1 In1 As4 Ga3 In3

Ga4Inte

grat

ed in

tens

ities

/ ar

b. u

nits

In coverage / nm

In/Se:GaAsb.)

As1 Ga1 In1 As4 Ga3 In2

In3

Figure 5.5: Integrated intensities of the different components in the As 3d and In 4d/Ga3d photoemission spectra as a function of increasing nominal In coverage on a.) sulphurand b.) selenium modified GaAs(100) surfaces. An In coverage of 0 nm corresponds tothe chalcogen modified GaAs(100) surface while a negative In coverage corresponds tothe clean GaAs surface.

92

clustering on this surface is stronger. That agrees with the observation made forthe growth of In on clean GaAs(100) [Spic90]. At last the thickness should bedetermined at which the evaporated In shows a metallic behaviour. This valuecan be evaluated from the occurrence of an asymmetric In component in thephotoemission spectrum, which occurs already at approximately 0.6 nm nominalIn coverage on S and Se modified GaAs surfaces. For In on clean GaAs surfacesa metallic behaviour is observed at 0.2 – 0.4 nm In coverage [Mao92]. Thismight be a consequence of a smaller mobility of the In atoms on the chalcogenmodified GaAs surface, which leads to a weaker clustering of the In atoms.

5.2 Mg on chalcogen modified GaAs(100) surfacesThe alkali metal Mg is known to be very reactive compared to In, which leads tothe expectation of stronger interface reactions on GaAs. Hence it is interesting toinvestigate the influence of the chalcogen modification of the GaAs surface onthe chemistry at the Mg/GaAs interface. Furthermore it is interesting to examinethe influence of chalcogens on the Schottky barrier formation for Mg on GaAs,which will be done in chapter 5.6. Mg was evaporated onto chalcogen modified GaAs(100) samples withincreasing thickness up to approximately 1 nm nominal thickness. After eachstep of modification photoelectron spectra were recorded of the different

participating elements. The photon energies for the photoemission measure-ments of the different core levels (Ga 3d (60 eV), As 3d (79 eV), Se 3d (88 eV),S 2p (195 eV) and Mg 2p (89 eV)) were always chosen to achieve maximumsurface sensitivity. To determine the binding energy of the Ga 3d and As 3d corelevel exactly after each spectrum the Fermi level position was measured on aclean piece of molybdenum. Thereafter the measured spectra were fitted usingspin-orbit split Voigt profiles to separate the contributions of chemical shiftedcomponents from the bulk.In the decomposition of the core level spectra the spin-orbit splitting, thebranching ratio and the Lorentzian width were always kept constant throughoutthe fits of the differently modified GaAs(100) surfaces. These values (compareto chapter 4) were first derived from the clean surface and thereafter optimised

Mg 2pLorentzian width / eV 0.1Branching ratio 2.0Spin-orbit splitting / eV 0.26Asymmetry factor 0.28

Table 5.2: Fit-parameters which are held constant during the different modification stepsof the GaAs(100) sample and for all Mg components occurring in the spectra.

93

to provide the best fit for the complete series. The constant parameters of the Mg2p core level in table 5.2 were derived from the spectra measured at the highestMg covarage. The other parameters such as binding energy, intensity,background (Shirley background) and Gaussian width were variable during thefits. The asymmetry parameters of the Doniach-Sunjic line shape of metalliccomponents was achieved from fitting the core level spectra in which themetallic component is most prominent.The chalcogen treatment of the clean GaAs(100) surface changes the As 3d andGa 3d core level spectra dramatically. Both As surface components disappear inthe spectra and only the bulk component (As1) and a small interface component(As4) remain, independent from the fact whether the samples are treated with Sor Se (figure 5.6 a. and d.).The Ga 3d core level spectra in figure 5.7 a. and d. exhibit two new intensesurface components (Ga4, Ga5) after the chalcogen treatment in comparison tothe spectra of the clean GaAs (100) surface (compare to chapter 4).From the fitting of the S 2p (figure 5.8 a.) and Se 3d (figure 5.9 a.) spectra twodistinct components can be recognised. The attribution of the components Se1and S1 to chalcogens on the surface and Se2 and S2 to chalcogens below thesurface have already been discussed in chapter 4.Looking at the As 3d spectra during increasing Mg coverage one observes thatthe As4 component remains on the lower binding energy side (figure 5.6 b. ande.). But the shift of this component decreases from approximately 0.75 eV on theS and Se modified surface to a saturation value of approximately 0.55 eV at 0.05nm of nominal Mg coverage. This is interpreted in terms of the formation of aMg3As2 like compound from the fact that this As4 component can not beattributed either to pure As-As bonds nor to As-chalcogen bonds because bothbonds should cause a shift to larger binding energy as can be concluded from thedifferences in electronegativity. The integrated intensity of the As4 componentshown in figure 5.11 a. and b. saturates also at a nominal Mg coverage ofapproximately 0.05 nm for S and 0.12 nm for Se. For S this coincides with thesaturation value of the binding energy shift of the As4 component as mentionedabove. This saturation indicates that the Mg-As compound segregates at thesurface, because its intensity does not decrease upon further Mg deposition. Incontrast the As1 component exhibits a strong dependence on the increasing Mgcoverage as can be seen in figure 5.11 a. and b.. From the exponential decreasein figure 5.11 a. and b. the escape depth of the photoelectrons from the As3dcore level is estimated to be 0.32 nm and 0.39 nm for Mg on Se and S modifiedGaAs surfaces, respectively. This coincides very well with the minimum in theescape depth of photoelectrons at 40 eV kinetic energy which is detected here.From this and the linear shape of the attenuation curve it is concluded that theinterface between the reacted surface and the GaAs bulk grows homogeneously(layer by layer) into the depth on Se and S modified GaAs(100) surfaces.

94

0.24 nm Mg

x 2

c.) 0.47 nm Mgf.)

0.12 nm Mg

Inte

nsity

/ ar

b. u

nits

b.) 0.12 nm Mge.)

44 43 42 41 40

As1 As4

Binding energy / eV

a.)

44 43 42 41 40

S:GaAsSe:GaAs

As4

As1

Binding energy / eV

d.)

Figure 5.6: As 3d core level spectra at different modification steps from GaAs(100)surfaces excited with 79 eV photon energy: a.) After Se modification at 435 °C, b.) afterevaporation of Mg with a nominal coverage of 0.12 nm and c.) 0.24 nm, d.) after Smodification at 480 °C, e.) after evaporation of Mg with a nominal coverage of 0.12 nmand f.) 0.47 nm.

95

0.95 nm Mgc.)

x 3

0.94 nm Mgf.)

x 3

0.12 nm Mg

Ga6

Ga7

Inte

nsity

/ ar

b. u

nits

b.)

x 3

0.12 nm Mg

Ga7

Ga6

e.)

x 2

22 21 20 19 18 17

Ga5

Ga4

Ga1

a.)

Binding energy / eV

22 21 20 19 18 17

S:GaAsSe:GaAs

Ga5

Ga4

Ga1

d.)

Binding energy / eV

Figure 5.7: Ga 3d core level spectra at different modification steps from GaAs(100)surfaces excited with 60 eV photon energy: a.) After Se modification at 435 °C, b.) afterevaporation of Mg with a nominal coverage of 0.12 nm and c.) 0.95 nm, d.) after Smodification at 480 °C, e.) after evaporation of Mg with a nominal coverage of 0.12 nmand f.) 0.94 nm.

96

163 162 161 160 159 158 157

S2S1a.)

Binding energy / eV

0.06 nm Mg

Inte

nsity

/ ar

b. u

nits b.)

S:GaAs

0.12 nm Mgc.)

Figure 5.8: S 2p core level spectra at different modification steps from GaAs(100) surfacesexcited with 193 eV photon energy, respectively: a.) After S modification at 480 °C, b.)after evaporation of Mg with a nominal coverage of 0.06 nm and c.) 0.12 nm.

97

58 57 56 55 54 53 52 51 50

Se2

Se1

a.)

Binding energy / eV

0.12 nm Mg

Mg2Mg1

b.)

Inte

nsity

/ ar

b. u

nits

0.95 nm Mg

Se:GaAs

Mg3c.)

Figure 5.9: Se 3d and Mg 2p core level spectra at different modification steps fromGaAs(100) surfaces excited with 89 eV photon energy: a.) After Se modification at 435°C, b.) after evaporation of Mg with a nominal coverage of 0.12 nm and c.) 0.95 nm.

98

54 53 52 51 50 49

S:GaAs

Mg2Mg1

a.)

Binding energy / eV

0.03 nm Mg

Mg3b.)

Inte

nsity

/ ar

b. u

nits

x 2

0.47 nm Mg

c.)

x 4

0.94 nm Mg

Figure 5.10: Mg 2p core level spectra at different modification steps from S modifiedGaAs(100) surfaces excited with 89 eV photon energy: a.) After evaporation of Mg with anominal coverage of 0.03 nm, b.) 0.47 nm and c.) 0.94 nm.

99

The Mg deposition also has a strong influence on the Ga 3d core level as can beseen in figure 5.7 b. and e. Two new components occur in the Ga 3d spectrumon the lower binding energy side. On S modified GaAs surfaces the shiftamounts to 0.66 eV (Ga7) and 1.12 eV (Ga6) towards lower binding energy,while on Se modified GaAs surfaces the shift amounts to 0.46 eV (Ga7) and1.04 eV (Ga6). The Ga6 and Ga7 component in the Ga 3d core level spectrumare assigned to metallic Ga in agreement with earlier investigations, where acomponent in the Ga 3d core level spectrum shifted by 0.9 eV towards lowerbinding energy relative to the bulk component was attributed to metallic Gaclusters [lay91].To separate between the bulk and surface components angle resolved spectrawere recorded. In figure 5.12 a. a Ga3d core level spectrum was measured at 60°off the normal direction of the sample, which leads to an increased surfacesensitivity. It is obvious that the components Ga4, Ga5, Ga6 and Ga7 areincreased in intensity relative to Ga1 in comparison to the spectrum measured innormal emission in figure 5.12 b. At higher photon energy the escape depth ofthe photoelectrons is increased which leads to a higher bulk sensitivity in thespectra. In figure 5.12 c. the Ga1 component is the most prominent feature in thespectrum, which is in agreement with the interpretation that it is assigned to thebulk.At higher Mg coverage the Ga6 component gets more pronounced and anasymmetric lineshape can clearly be observed, which confirms the metalliccharacter of this component (Doniach-Sunjic line shape). The Ga7 component,which can clearly be distinguished in figure 5.7 b. and e., is attributed to Gaatoms which are dissolved from the GaAs/GaxSy and GaAs/GaxSey interface,respectively. This component is no longer observed at a Mg coverage higherthan 0.23 nm, which leads to the conclusion that these Ga atoms now contributeto the Ga clusters represented by Ga6. The Ga4 and Ga5 components alsodisappear at this Mg coverage as can be seen in figure 5.7 c. and f. From thecomparison of their decay (figure 5.13 a. and b.) with the decay of the Ga bulkcomponent (Ga1), which decreases in intensity much stronger, it can beconcluded that the intensities of these components are not only attenuated by thelayer on the surface but also by the dissolving of the gallium chalcogenide.The dissolving of the gallium chalcogenide like compound is also indicated by achange in the binding energy shift of the two chalcogen components relative toeach other in figure 5.8 and 5.9. This is further supported by the change in therelative intensity ratio of Se1 to Se2 and S1 to S2, respectively. This is a sign fora change in the chemical environment of the chalcogen atoms, most likely beingdue to the formation of a magnesium chalcogenide like compound.A further proof for this conclusion can be observed in the Mg 2p core levels(figure 5.9 b. and 5.10 a.), where a broad feature occurs in the spectrum. Thebroad Mg peak covers the two proposed chemical environments of Mg, themagnesium chalcogenide like and the Mg3As2 like compounds, mentioned

100

0,0 0,1 0,2 0,3 0,4 0,5

Mg/Se:GaAs

a.)

Mg coverage / nm

As1 As4

Inte

grat

ed in

tens

ities

/ ar

b. u

nits

Mg/S:GaAs

b.)

As1 As4

Figure 5.11: Integrated intensities of the fitted components in the As 3d spectra at differentmodification steps of the GaAs(100) surface (the first points indicate the clean surface) a.)on Se modified GaAs and b.) on S modified GaAs.

101

37 38 39 40 41 42 43 44

a.)

Kinetic energy / eV

37 38 39 40 41 42 43 44

Ga5

Ga4Ga6

Ga7

Ga1b.)

x 4

x 8

125 126 127 128 129 130 131 132

c.)

Figure 5.12: Ga 3d core level spectra for 0.11 nm Mg coverage measured a.) at 60 eVphoton energy and 60° off the normal axis of the sample, b.) at 60 eV photon energy in thedirection of the surface normal and c.) at 150 eV photon energy normal to the samplesurface.

102

0,0 0,2 0,4 0,6 0,8 1,0

Mg/Se:GaAs a.)

Ga1 Ga4 Ga5 Ga6 Ga7

Inte

grat

ed in

tens

ities

/ ar

b. u

nits

Mg coverage / nm

Mg/S:GaAs b.)

Ga1 Ga4 Ga5 Ga6 Ga7

Figure 5.13: Integrated intensities of the fitted components in the Ga 3d spectra at differentmodification steps of the GaAs(100) surface (the first points indicate the clean surface) a.)on Se modified GaAs and b.) on S modified GaAs.

103

above. The two Mg components are separated by approximately 0.50 eV. Fromelectronegativity arguments it is concluded that the Mg1 component representsthe Mg-Chalcogen bonds and Mg2 the Mg-As bonds. This is in accordance withthe higher intensity of the Mg1 component, because the in the first steps of Mgevaporation mainly the gallium chalcogenide layer is dissolved and only athigher Mg coverage also the GaAs bulk is dissolved. This is further supportedby figure 5.10 b., where it is observed that the intensity of the Mg2 componentincreases relative to Mg1. At 0.23 nm nominal Mg coverage a slight asymmetryis observed in the Mg 2p spectrum, which increases with further Mg deposition.This asymmetry is assigned to a third Mg component (Mg3), which can beclearly separated with a binding energy shift of approximately 0.75 eV relativeto Mg2 on both Se and S modified GaAs surfaces. The Gaussian width of thisnew component is much smaller than for the other two Mg components and italso exhibits an asymmetry due to its metallic character (Doniac-Sunjic lineshape). This metallic component gets even more prominent after 0.95 nmdeposition of Mg (figure 5.9 c. and 5.10 c.). The occurrence of metallic Mg at0.23 nm Mg coverage points out that the chalcogens of the gallium chalcogenidelike layer on top of the GaAs(100) sample are consumed by the formation ofmagnesium chalcogenide, which coincides with the disappearance of the galliumchalcogenide related components (Ga4, Ga5) in the Ga 3d core level spectra.The observation of metallic Mg at 0.23 nm Mg coverage coincides also with theoccurrence of a Fermi edge distinguished in the valence band spectra of thesample, which is another indication for the existence of metallic Mg on top ofthe sample.

5.3 Na on Se modified GaAs(100) surfacesInterfaces between the alkali metal Na and GaAs(110) surfaces have been foundto be non-reactive for low temperatures while they are reactive at roomtemperature [Pri89]. So it is interesting to investigate whether Se is able toprevent a reaction between Na and GaAs. Furthermore it is interesting toexamine the influence of selenium on the Schottky barrier formation for Na onGaAs, which will be done in chapter 5.6.In this study Na is evaporated onto Se modified GaAs(100) samples withincreasing thickness. The thickness of the Na film could not be measured by aquartz crystal microbalance because of its low specific weight. Therefore thetime of evaporation under the same parameters of evaporation is taken as ameasure for the thickness of the evaporated Na film. Each modification step ismonitored by photoemission spectra of all participating elements. The photonenergies for the photoemission measurements of the different core levels (Ga 3d(60 eV), As 3d (79 eV), Se 3d (89 eV) and Na 2p (70 eV)) were always chosento achieve maximum surface sensitivity. To determine the binding energies ofthe Ga 3d and As 3d core level a Fermi edge on a clean peace of metal ismeasured at the same photon energy after each core level spectrum. Spin-orbit

104

split Voigt profiles are used to fit the measured core level spectra to separate thedifferent contributions in the spectra. During the fit the Lorentzian width, thebranching ratio and the spin-orbit splitting are kept constant. These values arederived from spectra of the clean and chalcogen modified GaAs surface andsubsequently optimised to fit the spectra during all modification steps. The otherparameters such as binding energy, intensity, Gaussian width and background(Shirley background) are variable during the fit.

After the Se modification of the GaAs surface the spectrum of the As 3d corelevel in figure 5.14 a. consists of only one component (As1). That means that theAs atoms now exist only in a GaAs bulk environment, because no surface orchemical shifted component exist anymore.The fit of the Ga 3d core level spectrum in figure 5.15 a. exhibits besides thecomponent Ga1 representing Ga atoms in the bulk like surrounding one newcomponent Ga3. This component is shifted by approximately 0.41 eV towardshigher binding energy.The Se 3d spectra in figure 5.16 a. consists of two different components fromwhich the Se1 component is attributed to Se dimers on top of the sample andSe2 to subsurface Se, which is bond to Ga atoms thus forming a gallium

60 s Na

As4

b.)In

tens

ity /

arb.

uni

ts2550 s Nad.)

x 2

44 43 42 41 40 39

As1

a.)

Binding energy / eV44 43 42 41 40 39

1350 s Nac.)

x 2

Binding energy / eV

Figure 5.14: As 3d core level spectra at different modification steps from GaAs(100)samples measured at 79 eV photon energy. a.) Se modified GaAs surface at 320 °C, b.)after 60 s, c.) after 1350 s and d.) after 2550 s evaporation of Na.

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selenide like compound. The energy separation between the two Se componentsamounts to approximately 0.92 eV.Na is evaporated onto this Se modified GaAs(100) surface. In figure 5.17 a. theNa 2p core level spectrum after 60 s of Na evaporation on the sample surface isshown. It is not possible to achieve a fit of this spectrum with uniqueparameters. But nevertheless it is observed that the spectrum must consist out oftwo different components, which is concluded from the asymmetry of thespectrum on the lower binding energy side. This may be explained by twodifferent chemical environments of the Na atoms. Since Na is known to be quitereactive at room temperature [Pri89] it might have reacted with the Se dimers ontop of the sample and also remain in Na clusters.To support this hypothesis the effect of the Na deposition on the core levelspectra of the other participating elements has to be taken into account. The

additional As4 component, which occurs after 60 s of Na evaporation (figure5.14 b.) originates from As atoms at the interface between the gallium selenidelike compound and the GaAs bulk. The component is shifted by approximately0.81 eV towards lower binding energy relative to the As bulk component. Thisvalue is comparable to the 0.75 eV shift, which has been observed for the Asinterface component at S modified GaAs(100) surfaces (compare to chapter 4).

60 s Na

Inte

nsity

/ ar

b. u

nits

b.) 2550 s Na

x 4

d.)

22 21 20 19 18 17

Ga1

Ga3

a.)

Binding energy / eV22 21 20 19 18 17

870 s Na

x 2

Ga4

c.)

Binding energy / eV

Figure 5.15: Ga 3d core level spectra at different modification steps from GaAs(100)samples measured at 60 eV photon energy. a.) Se modified GaAs surface at 320 °C, b.)after 60 s, c.) after 870 s and d.) after 2550 s evaporation of Na.

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The reason why it is now visible after the deposition of Na onto the surface maylay in a sharpening of the As 3d spectrum due to photoelectron diffraction. Thisis further supported by the fact that the Gaussian width of the components in theAs 3d core level spectrum is diminished by 20 meV. With increasingevaporation time of Na the energy shift of this component decreases and after120 s of Na deposition a saturation value of approximately 0.54 eV is achieved.Therefore the As4 component does not represent As at the gallium selenide-GaAs interface anymore. Moreover As4 is now interpreted with the formation ofAs-Na bonds which is possible because of the dissolution of the gallium selenideon the surface of the sample. After longer Na evaporation the intensity of theAs4 component increases even more (figure 5.14 c. and d.), which furthersupports the theory of the formation of As-Na bonds. The formation of As-alkalibonds has also been observed for the evaporation of Mg on chalcogen modifiedGaAs(100) surfaces, where nearly the same shift occurs for the As-Mg bonds.

The Na 2p spectra gets more symmetric with increasing Na coverage (figure5.17 b. and c.). Nevertheless it must contain two different components, becausethe Na exists in two different chemical environment. On the one side it is

60 s Nab.) 2550 s Nad.)

58 57 56 55 54 53

a.)

Se2Se1

Binding energy / eV

Inte

nsity

/ ar

b. u

nits

58 57 56 55 54 53

870 s Nac.)

Binding energy / eV

Figure 5.16: Se 3d core level spectra at different modification steps from GaAs(100)samples measured at 89 eV photon energy. a.) Se modified GaAs surface at 320 °C, b.)after 60 s, c.) after 870 s and d.) after 2550 s evaporation of Na.

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bonded to As as has been discussed above and on the other side it has reactedwith Se as will be discussed below.

Figure 5.18 b. exhibits the development of the intensities of the differentcomponents in the As 3d core level spectra as a function of the Na evaporationtime. The intensity of the As1 component decreases with increasing Naevaporation time, which results from the growth of Na compounds on top of thesample, so that the intensity of the bulk components is attenuated. On the otherside the intensity of the As4 component increases slightly with increasing Na

35 34 33 32 31 30 29

60 s Naa.)

Binding energy / eV

870 s Nab.)

Inte

nsity

/ ar

b. u

nits

2550 s Nac.)

Figure 5.17: Na 2p core level spectra at different modification steps from GaAs(100)samples measured at 70 eV photon energy. Se modified GaAs surface at 320 °C, a.) after60 s, b.) after 870 s and c.) after 2550 s evaporation of Na.

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evaporation time. Therefore it is concluded that this component originates fromAs atoms which are located on the surface. The slight increase can be attributedto further reaction of As with Na under the increasing exposure of Na atoms.

The Ga 3d core level does not change very much after 60 s of Na evaporation(figure 5.15 a.). The energy shift between the Ga1 and Ga3 component remainsmainly the same while only intensity of Ga3 is increased in relation to Ga1. This

0 500 1000 1500 2000 2500

a.)

Ga1 Ga3 Ga4

Time of Na deposition / s

b.)

As1 As4

Inte

nsity

/ ar

b. u

nits

Se1 Se2

c.)

Figure 5.18: Integrated intensities of the deconvoluted components in thedifferent core level spectra of the GaAs(100) surface at increasing Naevaporation time (the first points in a.) and b.) indicate the clean surface). a.)Ga 3d, b.) As 3d and c.) Se 3d.

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might be an effect of photoelectron diffraction. So the small amount of Na atomshas only reacted with the topmost Se layer of the Se modified GaAs sample.After a longer evaporation time an additional component (Ga4) on the lowerbinding energy side occurs in the Ga 3d spectrum (figure 5.15 c.). The bindingenergy shift between this component and the bulk component (Ga1) amounts toapproximately 0.68 eV. Therefore it can’t be assigned to metallic Ga since thebinding energy shift between metallic Ga and Ga in GaAs should amount to 0.9eV [Joy87]. The reduced chemical shift can be interpreted in terms of a reducedcore hole screening due to the limited size of Ga clusters and thus anincompletely developed metallicity [Che93]. A similar component also occurredin the spectra of the Ga 3d core level after the deposition of Mg (chapter 5.2).These small Ga clusters are formed from Ga atoms which are liberated from thegallium selenide like layer on top of the Se modified GaAs(100) sample,because of the reaction of Na with Se. This reaction is favoured due to the lowerelectronegativity of Na in comparison to Ga. After 2550 s evaporation of Nanearly the whole gallium selenide like layer is dissolved as can be concludedfrom figure 5.15 d., where the Ga4 component increases nearly to the sameheight as the Ga1 component while the Ga3 component nearly vanishes.This is even more obvious from figure 5.18 a., where the integrated intensities ofthe different Ga components in dependence of the Na evaporation time areshown. The intensity of the Ga1 component has approximately the same shapeas the intensity of the As1 component in figure 5.18 b. Therefore it is concludedthat the attenuation in intensity is due to the growth of a sodium selenide likecompound on top of the sample and that the Ga atoms in the bulk likesurrounding remain unaffected by the Na treatment of the sample. The change inintensity of the Ga1 and Ga3 component at short Na evaporation times may bedue to photoelectron diffraction. Comparing the intensity shape of the Ga3component with Ga1 it is obvious that the decrease in intensity of Ga3 is muchstronger than that of Ga1.That indicates that the chemical bonds, which are responsible for the occurrenceof Ga3, are changed due to the Na deposition, which is interpreted by thedissolution of the gallium selenide like compound on top of the Se modifiedGaAs(100) sample. These Ga atoms, which are set free from the galliumselenide, then show up as Ga4 component, which complementary increases inintensity with increasing Na evaporation time.The formation of a sodium selenide on top of the sample also influences the Se3d core level spectra. After evaporation of 60 s Na the intensity of the Se1component decreases (figure 5.16 b.). That leads to the conclusion that the Naatoms first react with the Se dimers on top of the Se modified GaAs(100)sample. This reaction is also supported by the decreasing of the binding energyshift between the Se1 and Se2 component from approximately 0.95 eV to 0.77eV, which results from the change in the chemical environment of the Se atoms.Although the Se atoms exist in three different chemical environments, the Se 3d

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core level spectra can be satisfactory fitted by only two different components.The reason for this may be that the difference in binding energy for the threedifferent components is to small to separate them in the spectrum. After 390 sNa evaporation the binding energy shift has achieved a value of 0.56 eV, whichdoes not change under further Na evaporation (figure 5.16 c.). That indicatesthat an equilibrium has been reached. The Se dimers on top of the sample nolonger exist and now the Se1 component in the Se 3d spectrum represents thesodium selenide compound. It is in agreement with electronegativity argumentsthat the sodium selenide compound occurs on the higher binding energy side ofthe spectrum, because the electronegativity of Na is lower than that of Ga. After2550 s of Na evaporation finally only the Se1 component remains in the Se 3dcore level spectrum. Therefore the gallium selenide like layer is dissolved anddoes not exist anymore. That can also be concluded from the shape of theintegrated intensities of the Se components in figure 5.18 c., where the Se1component mainly increases while the Se2 component mainly decreases due tothe increasing Na evaporation time. At the end of the experiment the surface ofthe sample is covered by a sodium selenide like layer and pure Na. Byevaporating more Na to the surface no change in the spectra can be observed sothat is must be concluded that a saturation is achieved and no more Na can stickon the surface at room temperature. At least a Fermi edge has been observedafter 870 s Na evaporation.

5.4 Sb on chalcogen modified GaAs(100) surfacesThe growth of Sb on GaAs(110) has attracted a lot of interest until now, becausethe Sb is known to be unreactive and forms an ordered first monolayer onGaAs(110) [Cao88, Ess90, Ess92, Kim93, Mag96, Mat85, Now95, Schä87,Tul86]. But there is hardly any study of Sb on GaAs(100) [Spi90] and so it isvery interesting to investigate the influence of S and Se on the chemistry and thegrowth of Sb on GaAs(100). Furthermore it is interesting to examine theinfluence of chalcogens on the Schottky barrier formation for Sb on GaAs,which will be done in chapter 5.6.

The chalcogen modified GaAs(100) surfaces serve as a starting point for thedeposition of Sb, which is evaporated in several steps. The thickness of the Sb

Sb 4dLorentzian width / eV 0.1Branching ratio 1.34Spin-orbit splitting / eV 1.25

Table 5.3: Fit-parameters which are held constant during the different modification stepsof the GaAs(100) sample and for all Sb components occurring in the spectra.

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film is monitored by a quartz crystal microbalance. After each evaporation stepphotoemission spectra are recorded from the As 3d (hν = 79 eV), Ga 3d (hν =60 eV), S 2p (hν = 195 eV), Se 3d (hν = 88 eV) and Sb 4d (hν = 70 eV) corelevels. The photon energies are always selected for maximum surfacesensitivity. The measured spectra were fitted using spin-orbit split Voigt profilesafter the subtraction of a Shirley background. In the fitting procedure theLorentzian width, the branching ratio and the spin-orbit splitting are heldconstant (table 5.3), while the intensity, the binding energy, the Gaussian widthand the parameters for the Shirley background are variable.The deposition of S and Se onto the hot sample, respectively, results in thedisappearance of the As surface components of the clean GaAs(100) surface(compare to chapter 4). So mainly the bulk component of GaAs remains (figure5.19 a. and d.). The small As4 component which occurs in the spectrum of the Smodified surface is shifted by approximately 0.70 eV.In the Ga 3d core level spectrum of the S modified GaAs(100) surface (figure5.20 a.) two chemically shifted components Ga3 and Ga4 occur, which areshifted by 0.46 and 0.90 eV relative to the bulk component (Ga1). The Semodification induces only one new chemically shifted component (Ga3) atapproximately 0.36 eV relative to the bulk (figure 5.20 d.).The spectra of the S 2p core level have a very poor resolution (ca. 1 eV), whichis related to the low resolution of the monochromator at higher energies.However, from these spectra it can be deduced that two different S components(S1 and S2) exist with 0.94 eV difference in binding energy. This agrees wellwith the Se 3d core level spectrum from the Se modified GaAs(100) surface(figure 5.21 d.), where also two components occur, which are separated by abinding energy shift of 0.91 eV. The origin of all observed components hasalready been discussed in chapter 4.The subsequent Sb deposition does not change the S 2p spectra very much(figure 5.21 b. and c.), only the intensity of the S1 component is lowered slightlycompared to S2, which can be explained by the formation of a bond between Sband S. The same is observed for the Se 3d spectra in figure 5.21 e. and f.The Sb 4d spectra from the S and Se modified GaAs(100) surface afterdeposition of Sb exhibit two interface components (Sb2 and Sb3) and onecomponent which can be attributed to amorphous antimony (Sb1) (figure 5.22 a.and d.). The binding energy shifts of the Sb2 and Sb3 component relative to Sb1amount on the S modified GaAs surface 0.85 and 1.70 eV and on the Semodified GaAs surface 0.57 eV and 1.16 eV. This difference in the bindingenergy shifts between the S and Se modified GaAs(100) surface is in accordancewith the higher electronegativity of S.Even at low nominal Sb coverage the Sb1 component occurs in the spectra. Thisleads to the conclusion that there is only one interface layer of antimonyrepresented by Sb2 and Sb3.

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c.) f.)

Inte

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/ ar

b. u

nits

b.) e.)

x 2

44 43 42 41 40 39

As1

As4

Binding energy / eV

a.)

44 43 42 41 40 39

As1

Binding energy / eV

d.)

Figure 5.19: As 3d core level spectra at different modification steps from GaAs(100)surfaces excited with 79 eV photon energy. a.) After S modification at 440 °C, b.) after0.06 nm, c.) after 0.88 nm nominal Sb coverage, d.) Se modified GaAs surface at 370 °C,e.) after 0.16 nm and f.) after 0.55 nm nominal Sb coverage.

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x 2

c.)

x 4

f.)

Inte

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b. u

nits

b.) e.)

x 2

22 21 20 19 18

Ga4

Ga3Ga1

Binding energy / eV

a.)

22 21 20 19 18

Ga3

Ga1

Binding energy / eV

d.)

Figure 5.20: Ga 3d core level spectra at different modification steps from GaAs(100)samples measured at 60 eV photon energy: a.) S modified GaAs surface at 440 °C, b.) after0.22 nm, c.) after 1.77 nm nominal Sb coverage, d.) Se modified GaAs surface at 370 °C,e.) after 0.16 nm and f.) after 1.07 nm nominal Sb coverage.

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c.) f.)

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b.) e.)

x 2

163 162 161 160 159 158 157

S1

S2

Binding energy / eV

a.)

57 56 55 54 53 52

Se2Se1

Binding energy / eV

d.)

Figure 5.21: S 2p and Se 3d core level spectra at different modification steps fromGaAs(100) surfaces excited with 195 eV and 89 eV photon energy, respectively. S 2pspectra after a.) S modification at 440 °C, b.) 0.03 nm and c.) 0.11 nm nominal Sbcoverage. Se 3d spectra after d.) Se modification at 370 °C, e.) 0.09 nm and f.) 0.55 nmnominal Sb coverage.

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x 10

c.)

x 4

f.)

x 2

Inte

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/ ar

b. u

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b.) e.)

x 2

36 35 34 33 32 31 30

Sb2

Sb3 Sb1

Binding energy / eV

a.)

36 35 34 33 32 31 30

Sb1

Sb2

Sb3

Binding energy / eV

d.)

Figure 5.22: Sb 4d core level spectra at different modification steps from GaAs(100)surfaces excited with 70 eV photon energy. a.) After 0.03 nm, b.) after 0.11 nm and c.)after 1.77 nm nominal Sb coverage on the Se modified GaAs(100) surface. d.) After 0.01nm e.) after 0.17 nm and f.) after 2.68 nm nominal Sb coverage on the S modifiedGaAs(100) surface.

116

0,0 0,2 0,4 0,6 0,8 1,0 1,2 1,4 1,6 1,8 3,4

BA

a.) Ga1 Ga3 Ga4 As1 Sb1

Inte

grat

ed in

tens

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/ a

rb.u

nits

Sb coverage / nm

BA

b.)

Ga1 Ga3 As1 Sb1

Figure 5.23: Integrated intensities of the fitted components in the different core levelspectra of the chalcogen modified GaAs(100) surface at increasing nominal Sb coverage(the first points in a.) and b.) indicate the clean surface) from a.) the S modified and b.) theSe modified GaAs(100) surface.

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These two components represent two different adsorption sites of the antimonyon top of the chalcogen modified GaAs(001) surfaces. Increasing Sb depositionleads to an increase of the intensity of the Sb1 component while the other twocomponents decrease, which supports the assignment of Sb2 and Sb3 to theinterface (figure 5.22 b. and e.). At high Sb coverage (figure 5.22 c. and f.) onlythe Sb1 component remains in the Sb 4d spectrum.The shape of the antimony core level does not show any asymmetry. Thereforeit can be concluded that the achieved film thickness is not sufficient to exhibitmetallic behaviour which is supported by the lack of a Fermi edge in the valenceband spectra.Although the antimony adsorbs on the surface of the sample there is no sign fora strong reaction neither in the S 2p or Se 3d nor in the Ga 3d spectra, whichwould yield in the formation of antimony chalcogenides and segregated Gaatoms or Ga clusters. The binding energy shifts of the different chemicallyshifted components in the As 3d and Ga 3d spectra do not change due to theincreasing Sb deposition (figures 5.19 and 5.20). Even the relative intensities ofthe shifted components relative to the bulk component do not change withincreasing Sb deposition (figure 5.23.). The only observation which can be madeis a decrease in intensity of the core level spectra of the other elements due tothe increasing Sb film thickness. This means that no alloying of antimony andgallium liberated from the gallium chalcogenide like compound takes place.The growth mode of the sample can be deduced from the GaAs(100) bulkcomponents (As1 and Ga1) and the amorphous antimony (Sb1). As can be seenin figure 5.23 a. and b. the intensity of the GaAs components decreases from thefirst to the second point drastically due to the growth of the galliumchalcogenide like layer.Thereafter in region A the growth of one complete monolayer of antimony canbe seen up to a nominal coverage of approximately 0.12 nm on S modifiedGaAs and 0.17 nm on Se modified GaAs. This can be deduced from thecoincidence of the escape depth of photoelectrons in GaAs which amounts to 0.5nm and the decay constant of the exponential decay of the Ga1 and As1intensities. Thereafter in area B islanding is observed which can be concludedfrom the decreased attenuation of As1 and Ga1 and the rather low increase ofSb1 with further evaporation of antimony. Therefore the growth mode of Sb onchalcogen passivated GaAs can be described by the Stranski-Krastanov growthin contrast to Frank-van-der-Merwe growth observed for Sb on clean GaAs(100)[Spin90].

5.5 Ag on chalcogen modified GaAs(100) surfacesThe study of the Schottky barrier formation is complicated by the fact that themetal/semiconductor interface is very complex with many processes occurringsimultaneously. Among these complicating processes are diffusion and chemicalreaction. To minimise these effects, one would like to choose a metal which

118

reacts weakly with GaAs. Therefore Ag is an excellent choice based on thisconsideration. Evidence of cluster formation even at room temperature isindicative of its low reactivity with GaAs [Lud83, Chi85], because it reveals thestrong preference of the Ag atom to bond to other Ag atoms rather than to a Gaor As atom. Therefore the growth of Ag on clean GaAs surfaces has attracted alot of interest in the past [Cao87, Chi88, Chi89, Chi85, Lud82, Spi90, Sti88,Wad 89, Wal93], but there is hardly any study dealing with the influence of thechalcogen modification on the Ag/GaAs(100) interface. Furthermore it isinteresting to examine the influence of chalcogens on the Schottky barrierformation for Ag on GaAs, which will be done in chapter 5.6.To investigate the interface formation the chalcogen modified GaAs(100)samples are subsequently covered by silver with increasing thickness monitoredby a quartz microbalance. At each modification step photoemission spectra arerecorded from the As 3d (hν=79 eV), Ga 3d (hν=60 eV), Se 3d (hν=89 eV) and S2p (hν=195 eV) core levels and valence band spectra (hν=24 eV) with photonenergies providing maximum surface sensitivity. The measured spectra are fittedusing spin-orbit split Voigt profiles and a Shirley background. In the fit routinethe values for the spin orbit splitting, the branching ratio and the Lorentzianwidths are kept constant for all spectra, whereas the intensities, the bindingenergy shifts and the Gaussian widths are variable. The same Gaussian width isused for all components in one certain spectrum.The chalcogen modification of the clean GaAs(100) surface causes a dramaticchange in the spectra of the As 3d core level spectrum. As observed in figure5.24 a. only a bulk component (As1) and a small component at lower bindingenergy (As4) occur in the As 3d spectrum on the S modified GaAs surface. Thiscomponent is shifted by approximately 0.70 eV relative to the bulk component.On the Se modified GaAs surface only a bulk component (As1) is observed(figure 5.24 d.). In this As 3d core level spectrum there is no indication for afurther component, which might be reasoned by the low intensity of thiscomponent and an inhomogenity of the gallium selenide surface layer.The S modification at 490 °C causes a change in the lineshape of the Ga 3d corelevel spectrum (figure 5.25 a.). The two new surface components are shiftedtowards higher binding energy by 0.50 eV (Ga3) and 0.96 eV (Ga4),respectively. The Se modification of the GaAs(100) surface at 380 °C causesonly one new chemically shifted component Ga3 (figure 5.25 d.). The Ga3component is shifted by approximately 0.43 eV towards higher binding energy.The lineshape of the S 2p core level spectrum (figure 5.26 a.) measured on theGaAs(100) surface which is exposed to an atomic S beam at 490 °C indicates atleast two different components. The energy shift between these two Scomponents amounts to approximately 0.96 eV.This is similar to the GaAs surface, which is modified by Se at 380 °C. In the Se3d core level spectrum occur also two different components, which are shifted toeach other by approximately 0.88 eV.

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1.96 nm Ag

c.)

0.91 nm Agf.)

0.06 nm Ag

Inte

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b.) 0.06 nm Ag

Se:GaAsS:GaAs

As4

e.)

44 43 42 41 40 39

As1

As4

Binding energy / eV

a.)

44 43 42 41 40 39

As1

Binding energy / eV

d.)

Figure 5.24: As 3d core level spectra at different modification steps from GaAs(100)surfaces excited with 79 eV photon energy: a.) After S modification at 490 °C, b.) afterevaporation of Ag with a nominal coverage of 0.06 nm and c.) 1.96 nm. d.) After Semodification at 380 °C, e.) after evaporation of Ag with a nominal coverage of 0.06 nmand f.) 0.91 nm.

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1.96 nm Ag

x 2

c.) 0.91 nm Agf.)

0.06 nm Ag

Inte

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b. u

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b.) 0.06 nm Age.)

22 21 20 19 18

Ga4

Ga3

Ga1

Binding energy / eV

a.)

22 21 20 19 18

Se:GaAsS:GaAs

Ga1

Ga3

Binding energy / eV

d.)

Figure 5.25: Ga 3d core level spectra at different modification steps from GaAs(100)surfaces excited with 60 eV photon energy: a.) After S modification at 490 °C, b.) afterevaporation of Ag with a nominal coverage of 0.06 nm and c.) 1.96 nm. d.) After Semodification at 380 °C, e.) after evaporation of Ag with a nominal coverage of 0.06 nmand f.) 0.91 nm.

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The Se1 and S1 components are assigned to chalcogen dimers on top of thesample, while the Se2 and S2 components originate from incorporatedchalcogen atoms (compare to chapter 4).Subsequently Ag is deposited onto the chalcogen modified GaAs(100) surfaceswith increasing thickness. The evaporation of a small amount of Ag (0.06 nm)onto these surface causes nearly no changes in the As 3d core level spectra. Theintensity of the As 3d spectrum from the S modified GaAs surface (figure 5.24b.) decreases slightly, while the relation of the intensities of the As1 and As4component remains mainly constant. In the As 3d spectrum of the Se modifiedGaAs(100) surface an additional component (As4) can be separated. This mightbe explained by the fact that an inhomogeneous Fermi level pinning is cancelleddue to the Ag deposited on top of the sample. Even a high Ag coverage (> 0.9nm) does not change the chemical environment of the As atoms, which isindicated by the constant relation between the intensity of the As1 and As4component and the fact that the binding energy shift between these twocomponents does not change (figure 5.24 c. and f.).

0.12 nm Ag

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b.) 0.12 nm Agd.)

163 162 161 160 159 158 157

S2 S1

Binding energy / eV

a.)

57 56 55 54 53 52

Se:GaAsS:GaAs

Se2Se1

Binding energy / eV

c.)

Figure 5.26: S 2p and Se 3d core level spectra at different modification steps fromGaAs(100) surfaces excited with 194 eV and 89 eV photon energy, respectively: a.)After S modification at 490 °C and b.) after evaporation of Ag with a nominal coverageof 0.12 nm. d.) After Se modification at 380 °C and e.) after evaporation of Ag with anominal coverage of 0.12 nm.

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-10 -8 -6 -4 -2 0

Binding energy / eV

a.)

b.)

/ 0,8

c.)

d.)

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tens

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/ ar

b. u

nits

/ 1,4

/ 2

/ 2,6e.)

f.) / 2,8

/ 3,4

h.)

g.)

Figure 5.27: Valence band spectra at different modification steps from GaAs(100)surfaces excited with 24 eV photon energy: a.) After S modification at 490 °C, b.) afterevaporation of Ag with a nominal coverage of 0.06 nm, b.) 0.12 nm, c.) 0.24 nm, d.)0.48 nm, e.) 0.98 nm, f.) 1.47 nm and g.) 1.96 nm. The valence band spectra from Ag onthe Se modified GaAs surface are very similar.

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The attenuation of the intensity of the As 3d core level spectrum is the onlyeffect of the increasing Ag deposition.The deposition of a small amount of Ag (0.06 nm) onto the S modified GaAssurface has the effect that the intensities of the Ga3 and Ga4 componentsincrease in the Ga 3d core level spectrum (figure 5.25 b.).In the next steps of Ag evaporation (0.12 nm and more) it is observed that theintensity of the Ga3 component decreases with a higher attenuation constantthan the Ga1 component while the intensity of the Ga4 component increaseseven more (figure 5.29 a.).Since the Ga3 and Ga4 components derive from Ga atoms nearer to the surfaceof the sample than the Ga1 component this might be an effect of photoelectrondiffraction due to the growth of an ordered Ag layer on the surface. This isfurther supported by the fact that the binding energy shifts of the Ga3 and Ga4components relative to the Ga1 component remain constant due to the Ag

deposition. So there is no sign of a chemical reaction. At higher Ag coverage theintensity ratio between the surface and the bulk component is constant and onlythe total intensity is decreased (figure 5.25 c.).For small amounts of Ag (0.06 nm) on the Se modified GaAs surfaces, it isobserved that the intensity of the Ga3 component decreases with a smaller

0,1 1

1,4

1,5

1,6

1,7

1,8

1,9

2,0

2,1

2,2

2,3

2,4

FWH

M /

eV

Ag coverage / nm

Figure 5.28: Full width at half maximum (FWHM) of the Ag 4d band, which occurs inthe valence band spectrum at increasing Ag coverage of the S modified GaAs(100)surface.

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attenuation length than the bulk component Ga1 (figure 5.25 e.). This might alsoexplained by photoelectron diffraction as has been described above for Ag onthe S modified GaAs surface. The binding energy shift between Ga1 and Ga3remains constant up to a high Ag coverage so that it can be concluded that theadsorption of Ag does not disrupt the gallium selenide layer on top of thesample. That coincides with the observation for Ag on the S modified GaAs.The only effect of the large Ag deposition (> 0.9 nm) is the strong attenuation ofthe total intensity of the Ga 3d core level spectrum.Considering now the effect of the Ag deposition on the chalcogen spectrareveals a diminished binding energy shift between the two different componentsof the chalcogen core level spectra. In the S 2p as well as in the Se 3d core levelspectrum the binding energy shift is lowered by approximately 0.2 eV (figure5.26 b. and d.).This fact is explained with the formation of a bond between Ag and thechalcogen dimers which changes the chemical environment of the chalcogendimers on top of the sample. The intensity ratio of the two chalcogencomponents remains constant due to the increasing Ag coverage which supportsthe assumption that the gallium chalcogenide layer on top of the sample is notdisrupted.From the Ag atoms only the Ag 4d band is accessible at the monochromatorwhich is used for the experiments. The binding energy of the Ag 4d levelamounts to approximately 6 eV and therefore it occurs in the valence bandspectrum of the sample. Figure 5.27 shows that the spectrum from the cleansurface (a.) exhibits a minimum in intensity at approximately 6 eV bindingenergy, which is filled by the emission from the Ag 4d level after the depositionof 0.06 nm Ag (b.). With increasing Ag coverage the Ag 4d doublet increases inintensity and its lineshape develops to the normal shape (figure 5.27 c. – h.)which has also been observed by others [Chi85, Chi89, Lud82, Sti88, Wer86].The full width at half maximum (FWHM) of the Ag 4d band can be used torecognise whether the evaporated Ag film on top of the sample exhibits ametallic behaviour. A narrow Ag 4d band represents an atomic character of theAg atoms, while with the widening of the band a transition to metallic clusterstakes place [Wer86]. In figure 5.28 it is observed that the FWHM increases froma value of 1.45 eV at 0.06 nm nominal Ag coverage up to 2.25 eV at 0.48 nm,thereafter there is nearly no further increasing observed. Therefore it isconcluded that with a nominal Ag coverage of 0.48 nm the Ag clusters on top ofthe chalcogen modified GaAs surface exhibit a metallic character. At 0.48 nmnominal Ag coverage a Fermi edge occurs in the valence band spectrum, whichalso indicates the metallic character of the deposited Ag film.The deposition of Ag causes an attenuation of the integrated intensities of thecomponents in the core level spectra measured on GaAs(100) at differentmodification steps. The attenuation of distinct components in the spectra from

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the S modified GaAs surface can be separated into two distinct regions A and B(Figure 5.29).

0,0 0,2 0,4 0,6 0,8 1,0 1,2 1,4 1,6 1,8 2,0

Ag/S:GaAs a.)BA

Ga1 Ga4 Ga3 As1

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/ ar

b.un

its

Ag coverage / nm

Ag/Se:GaAs b.)

BA

Ga1 Ga3 As1

Figure 5.29: Integrated intensities of the fitted components in the As 3d and Ga 3d corelevel spectra at different modification steps of the GaAs(100) surface (the first pointsindicate the clean surface, while the second points indicate the chalcogen modifiedGaAs(100) surface) a.) on S modified GaAs and b.) on Se modified GaAs.

126

Region A starts at 0.06 nm and region B at approximately 0.5 nm Ag coverage(figure 5.29 a.). In the area in front of region A it can be observed that the Smodification causes a decrease in the intensities of the As and Ga bulkcomponents (As1, Ga1) due to the formation of the gallium sulphide like layeron top of the sample. The Ga1, Ga3 and As1 components behave similar inregions A and B. Calculating the nominal escape depth of the photoelectronsfrom the exponential decay of the intensity curves leads to a value of 1.7 nm and4.5 nm in region A and B, respectively. Comparing these values to the escapedepth of photoelectrons at approximately 40 eV kinetic energy, which amountsto about 0.5 nm, leads to the conclusion that the deposited Ag forms islands ontop of the sample. The two distinct regions indicate that this islanding getsstronger after the deposition of 0.5 nm. The behaviour of the Ga4 component inregion A revealing a stronger attenuation after an initial increase is not wellunderstood yet since. Its nominal escape depth of 1 nm is lower than observedfor the other components in this region, while it coincides with the others inregion B.The deposition of Ag on the Se modified GaAs surface also shows mainly tworegions which differ in their attenuation length (figure 5.29 b.). Region Aproceeds into region B at 0.12 nm of nominal Ag coverage, which is slightlylower than the border found on S modified GaAs surfaces. The attenuation in thebulk components Ga1 and As1 from the first to the second point is caused by thedeposition of Se onto the GaAs surface at elevated temperature which results inthe growth of a gallium selenide like layer on top of the sample. The subsequentAg deposition also causes a strong decrease in the intensities of the two bulkcomponent with an attenuation length of approximately 0.7 nm. This value iscomparable to the escape depth of photoelectrons at 40 eV kinetic energy, whichleads to the conclusion that the Ag grows nearly two dimensional up to 0.12 nmcoverage. At higher Ag coverage the attenuation length increases toapproximately 2.3 nm, which means that the Ag atoms form islands. Thisattenuation length is found for the bulk components Ga1 and As1 as well as forthe surface component Ga3, which supports the suggestion that the galliumselenide like layer on top of the sample is not destroyed by the deposition of Ag.The comparison of the growth mode of Ag on chalcogen modified GaAs(100)surfaces thus reveals that in the case of Se a growth mode similar to theStranski-Krastanov mode is found, while in the case of S no complete Ag layeris formed and the formation of Ag islands starts even at a low Ag coverage. Thatmight be explained by a higher mobility of the Ag atoms on the S surface thanon the Se modified GaAs surface. In conclusion the growth of Ag on chalcogenmodified GaAs(100) surfaces differs strongly from the growth of Ag on theclean GaAs(100) surface, which reveals a Volmer-Weber growth [Spic90].

127

5.6 Electronic properties of chalcogen modified metal/GaAs(100)contacts

In the preceding chapters the chemical properties of the chalcogen modifiedmetal/GaAs(100) contacts have been considered. Now electronic properties ofthese interfaces are the main points of interest such as the position of the Fermilevel at different modification steps and the formation of the Schottky barrier.The measurement of the Fermi level by photoemission during increasingadsorbate deposition is quite difficult, because no sharp valence band edge ofthe underlying semiconductor can be observed at higher adsorbate coverage. Butit is a suitable method to determine the Fermi level of the clean surface. Onclean n- and p-doped GaAs surfaces the Fermi level relative to the valence bandmaximum (VBM) was carefully determined as has been described in chapter 3.Additionally the binding energies of the Ga 3d5/2 and As 3d5/2 components weremeasured relative to the Fermi level (see table 4.2). Thereafter the Fermi levelrelative to the VBM (see table 4.3) were subtracted from these values leading to18.77 eV for Ga 3d5/2 and 40.65 eV for As 3d5/2 coincident for n- and p-dopedGaAs(100), which agrees well with the values determined by Eastman et al.[Eas80]. The subtraction of these binding energies relative to the VBM from thebinding energies of Ga 3d and As 3d, which were measured relative to the Fermilevel on adsorbate covered surfaces, leads to the position of the Fermi levelrelative to the valence band maximum [Kra80]. For this method it is importantto measure the photoemission spectra in a surface sensitive mode so that onlythe first layers are probed. Otherwise the determined Fermi level would notprovide the surface value.In the following the values derived from the shift of the Ga 3d bulk component(Ga1) are considered, because they agree with the values derived from the As 3dbulk component (As1).Figures 5.30 to 5.34 show the position of the Fermi level relative to the valenceband maximum (VBM) for different metal adsorbates on S and Se modifiedGaAs surfaces, respectively.It is generally observed, that the chalcogen treatment moves the Fermi levelcloser to the conduction band minimum (CBM), thus indicating a decrease andincrease of band bending on samples n- and p-doped, respectively. Therefore ishas been concluded in chapter 4 that the chalcogen treatment leads to a newdistribution of surface states within the gap rather than a removal of the surfacestates from the gap (chapter 4).In the following it will be concentrated on the influence, which differentmetallic adsorbates have on the position of the Fermi level. The adsorbatedeposition on the chalcogen modified n-GaAs surfaces causes in nearly allinvestigated cases a more or less strong shift of the Fermi level towards theCBM. After the first evaporation step the Fermi level lies between 0.9 eV and1.1 eV.

128

0 1 2 3 4 5

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1,1

1,2In/Se:GaAs a.)

n-GaAs(100) p-GaAs(100)

Ef

- V

BM

/ eV

In coverage / nm

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1,1

1,2In/S:GaAs b.)

n-GaAs(100) p-GaAs(100)

Figure 5.30: Position of the Fermi level on chalcogen modified n- and p-dopedGaAs(100) surfaces as a function of In coverage: a.) Se- and b.) S-modification. Thefirst points indicate the Fermi level on the clean surface, while the second pointsrepresent the chalcogen modified surface.

129

0,0 0,4 0,8 1,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1,1

1,2

Mg/Se:GaAs a.)

n-GaAs(100) p-GaAs(100)

Ef

- VB

M /

eV

Mg coverage / nm

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1,1

1,2

Mg/S:GaAs b.)

n-GaAs(100) p-GaAs(100)

Figure 5.31: Position of the Fermi level on chalcogen modified n- and p-dopedGaAs(100) surfaces as a function of Mg coverage: a.) Se- and b.) S-modification. Thefirst points indicate the Fermi level on the clean surface, while the second pointsrepresent the chalcogen modified surface.

130

This overshoot occurs only at a very low adsorbate coverage of approximately0.05 nm. This observation is made for In (figure 5.30), Mg (figure 5.31), Na(figure 5.32) and Ag (figure 5.34). The strong decrease in band bending isattributed to the initial stage of the adsorbate deposition when the adsorbatesdisplay not yet metallic properties.It is concluded that adsorbate-induced states of acceptor type are build on thechalcogen modified GaAs surface. This is the reason why the strong decrease inband bending is only observed on n-doped GaAs. At approximately 0.2 nmadsorbate coverage the overshoot decreases again and the Fermi levelapproximates to its saturation value, which remains nearly constant up to themaximum coverage. This in interpreted by a continuous transition from bandbending due to adsorbate-induced states to a situation dominated by metalinduced gap states (MIGS). The adsorbate coverage at which the metallizationof the different overlayers is observed (compare to chapter 5.1 – 5.5) agrees wellwith this 0.2 nm. In almost all cases of investigated adsorbates the Fermi levelreaches a saturation value with reference to the adsorbate thickness. These Fermilevel values are shown in table 5.4.Mg represents an exception, because no saturation is achieved for the Fermilevel position on Se as well as on S modified GaAs surfaces.

0 500 1000 1500 2000 2500

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1,1

1,2

Na/Se:GaAs

n-GaAs(100) p-GaAs(100)

Ef

- VB

M /

eV

Na coverage / s of evaporation

Figure 5.32: Position of the Fermi level on Se modified n- and p-doped GaAs(100)surfaces as a function of Na coverage. The first points indicate the Fermi level on theclean surface, while the second points represent the chalcogen modified surface.

131

0,0 0,4 0,8 1,2 1,6 2,0 2,4 2,8

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1,1

1,2Sb/Se:GaAs a.)

n-GaAs(100) p-GaAs(100)

Ef

- VB

M /

eV

Sb coverage / nm

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1,1

1,2Sb/S:GaAs b.)

n-GaAs(100) p-GaAs(100)

Figure 5.33: Position of the Fermi level on chalcogen modified n- and p-dopedGaAs(100) surfaces as a function of Sb coverage: a.) Se- and b.) S-modification. Thefirst points indicate the Fermi level on the clean surface, while the second pointsrepresent the chalcogen modified surface.

132

0,0 0,4 0,8 1,2 1,6 2,0

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1,1Ag/Se:GaAs a.)

n-GaAs(100) p-GaAs(100)

Ef

- V

BM

/ eV

Ag coverage / nm

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1,1

Ag/S:GaAs b.)

n-GaAs(100) p-GaAs(100)

Figure 5.34: Position of the Fermi level on chalcogen modified n- and p-dopedGaAs(100) surfaces as a function of Ag coverage: a.) Se- and b.) S-modification. Thefirst points indicate the Fermi level on the clean surface, while the second pointsrepresent the chalcogen modified surface.

133

This results from the strong reactivity of Mg, which leads to the dissolution ofthe gallium chalcogenide layer on top of the GaAs samples and the segregationof metallic Ga clusters. Hence it has been very difficult to determine the bindingenergies of the Ga and As bulk components above 0.24 nm nominal Mg

coverage. So it was not possible to evaluate the exact Fermi level position forhigher Mg coverage from the photoemission experiments. In this case the switch

Interface EF / eV ΦBPES /

eVΦB

IV / eV n

Ag/Se:n-GaAs

0.73 0.70 0.73 1.24

Ag/Se:p-GaAs

0.59 0.59 0.55 1.20

Ag/S:n-GaAs 0.80 0.63 0.73 1.25Ag/S:p-GaAs 0.74 0.74 0.51 1.25Sb/Se:n-GaAs 0.81 0.62 - -Sb/Se:p-GaAs 0.60 0.60 - -Sb/S:n-GaAs 1.05 038 - -Sb/S:p-GaAs 0.66 0.66 - -In/Se:n-GaAs 1.01 0.42 - -In/Se:p-GaAs 0.97 0.97 - -In/S:n-GaAs 0.85 0.58 - -In/S:p-GaAs 0.73 0.73 - -Mg/Se:n-GaAs

0.79 0.64 0.55 1.32

Mg/Se:p-GaAs

0.79 0.79 0.53 1.35

Mg/S:n-GaAs 0.88 0.55 0.59 1.10Mg/S:p-GaAs 0.77 0.77 0.55 1.62Na/Se:n-GaAs

0.54 0.89 - -

Na/Se:p-GaAs

0.60 0.60 - -

Table 5.4: Position of the Fermi level EF relative to the valence band maximum (VBM)measured by Photoemission spectroscopy (PES), Schottky barrier heights (ΦB

PES)measured by PES, Schottky barrier heights (ΦB

IV) determined from current-voltage (IV)measurements and ideality factors (n) from the determination of the Schottky barrierheight from IV measurements of different chalcogen modified metal/GaAs(100)Schottky contacts.

134

to current-voltage measurements is advisable for the determination of theSchottky barrier height, because these experiments can be applied for a thickmetal coverage. The results will be described below.Comparing the Fermi level positions for the different examined adsorbates on n-and p-doped GaAs samples at their saturation value, it is found out that theycoincide quite well. The maximum deviation amounts to approximately 0.14 eV.This is in agreement with the theory, that the Schottky barrier heights on n- andp-doped semiconductor samples should sum up to the band gap for a metallicadsorbate (compare to chapter 2).The Schottky barrier height on n-doped samples is determined by the energyseparation between the Fermi level and the CBM, while on p-doped samples theSchottky barrier height equals the difference between the Fermi level and theVBM.An exception is made by Sb on Se (figure 5.33 a.) or S (figure 5.33 b.) modifiedGaAs, which exhibits no coincidence in the Fermi level position after thesaturation value is reached. Also the Fermi level does not show an overshoot onn-doped GaAs. That leads to the conclusion that Sb does not get metallic up tonominal coverage of 2.8 nm, which agrees well with the fact that no Fermi levelis observed on the sample up to this coverage (compare to chapter 5.4). OnGaAs(110) it has been observed that the Schottky barrier heights for Sb on n-and p-doped samples coincide for a coverage of more than 100 monolayers,which is much more than the amount of Sb which was deposited in thisinvestigation [Ess90].In the case of n-doped GaAs the saturation value for the Fermi level of Sb on Seand S modified GaAs(100) amounts to 0.81 eV and 1.05 eV, respectively (table5.4). This is considerable higher in comparison to the Fermi level of Sb onunmodified n-doped GaAs(110), which amounts to 0.75 eV [Cao88]. Thereforea reduction in band bending has been achieved due to the chalcogenmodification of the interface. The opposite is found for p-doped GaAs, wherethe Fermi level amounts to 0.60 eV and 0.66 eV for Sb on Se and S modifiedGaAs(100), respectively. This leads to the conclusion that the band bending onp-doped GaAs is increased, because the Fermi level for Sb on clean p-dopedGaAs(110) amounts to 0.50 eV.From the decrease of band bending on n-doped and increase on p-doped samplesit is concluded that the chalcogen modification has not lead to a removal of allstates in the band gap. Furthermore new states of acceptor and donor type arecreated in the band gap, which pin the Fermi level.In the case of Mg and Ag additionally to the determination of the Schottkybarrier height by photoemission spectroscopy current-voltage measurementswere applied. In contrast to the photoemission spectroscopy which is onlyapplicable for thin overlayers the current-voltage measurement can be appliedfor thick metallic overlayers.

135

-0,2 -0,1 0,0 0,1 0,2 0,3 0,410-11

10-10

10-9

10-8

10-7

10-6

10-5

ΦB = 0.51 eV

n = 1.25RS = 38 Ω

a.)

Cu

rren

t /

A

Voltage / V

10-7

10-6

10-5

10-4

10-3

10-2-0,2 -0,1 0,0 0,1 0,2 0,3 0,4

ΦB = 0.59 eV

n = 1.10RS = 5 Ω

b.)

Ag/S:p-GaAs

Mg/S:n-GaAs

Figure 5.35: Current-Voltage characteristics, which are measured on S-modified Ag/n-GaAs(100) and S-modified Mg/p-GaAs(100) Schottky diodes.

136

The samples for these measurements were prepared in the same way as thesamples for the photoemission measurements. The only difference lies in thefact that the metal was evaporated through a grid thus forming metal dots on topof the sample in contrast to a continuous film.After the formation of the metal contacts the samples were taken out of thevacuum and the I-V measurements were carried out ex situ. Two examples forthe measured I-V characteristics are shown in figure 5.35 a. and b. According tochapter 3.4 the measured I-V curves are fitted using formula 3.18. The evaluatedvalues for the Schottky barrier heights ΦB and the ideality factors n are shown intable 5.4.From the photoemission measurements of Mg on chalcogen modified GaAs ithas been observed that the slope of the Fermi level (figure 5.31) increases withincreasing Mg coverage. Therefore it is expected that the barrier heightsdetermined by current-voltage measurements will exhibit a decreased andincreased barrier height on n- and p-doped samples in comparison to the valuesdetermined from photoemission, respectively. This expectation is true in thecase of n-doped GaAs, while in the case of p-doped GaAs the barrier height isdecreased. The ideality factors for the Mg/GaAs Schottky contacts exceed theideal value (n = 1) to a high amount with exception of Mg/S:n-GaAs. This mightbe explained by laterally inhomogenous Schottky contacts. Schmitsdorf et al.found that the barrier heights for these Schottky contacts decreases as a functionof increasing ideality factor [Schm95a]. Hence the barrier height for the idealSchottky contact might be higher.In the case of Ag the barrier heights determined from photoemission agree wellwith the values determined by photoemission with the exception of Ag/S:p-GaAs. This might be also explained with an inhomogenity in the Schottkycontact.In chapter 2 it has been shown that the Schottky barrier height of ametal/semiconductor contact can be calculated according to the model of metalinduced gap states (MIGS). Therefore the slope parameter SX has to be estimatedby using equation 5.1, which was derived from equation 2.7.

( )1

129,0 2

+−⋅=

∞

i

XX

AS

εε (5.1)

Inserting 0.86 for the parameter AX [Mön95], which is valid for the use ofMiedema’s electronegativities, 10.9 for the dielectric constant ε∝ of GaAs atroom temperature (RT) [Mad96] and 4 for the dielectric constant εi at theinterface of the semiconductor and the metal [Lud88] in formula 5.1 yields0.106 for the Miedema slope parameter SX .

137

( )SMxCNLBp S Χ−Χ⋅+Φ=Φ (5.2)

For the calculation of the Schottky barrier heights it is essential to know thebranch point energy of GaAs. This value was first estimated from measurementof the indirect band gap of GaAs [Ter85], but it was also calculated by theempirical tight-binding approximation (ETB), the Green’s function method(GFM) and the dielectric midgap energies [Mön96b and references therein]. Inthese references the values for the branch point energy of GaAs range between0.50 eV and 0.55 eV. Here the experimentally estimated value of 0.52 eV[Ter85] is taken for the branch point energy. The Schottky barrier heights fordifferent metals on GaAs can be calculated by formula 5.2 by using thecalculated Miedema slope parameter SX , the branch point energy of GaAs ΦCNL,the electronegativity of the different metals ΧM and the electronegativity ofGaAs ΧS, which amounts to 4.44 [Mön95] (table 5.5).These values are only valid for the abrupt metal/semiconductor interface.

Additionally secondary effects have to be taken into account, because theexamined samples do not exhibit an abrupt metal/semiconductor interface, e.g.the formation of an interface dipole or defect states at the interface.The comparison of the experimentally determined Schottky barrier heights ofthe chalcogen modified (table 5.4) with the unmodified metal/GaAs contacts(table 5.5) exhibits a decrease and an increase of the barrier height on n-dopedand p-doped GaAs, respectively.A similar behaviour was observed for lead contacts on hydrogen-passivatedSi(111) surfaces where the hydrogen-passivation leads to an increase of theposition of the Fermi level with respect to VBM [Kam95a]. The change in

Element ENMiedema ΦBntheo ΦBp

theo ΦBnexp ΦBp

exp

Ag 4.45 0.91 0.52 0.90 0.50Sb 4.40 0.91 0.52 0.68 0.50In 3.90 0.96 0.47 0.83 0.40Mg 3.45 1.00 0.43 0.62 0.55Na 2.70 1.06 0.37 0.90 0.53

Table 5.5: Theoretically (ΦBtheo) and experimentally (ΦB

exp) determined Schottky barrierheights of an unmodified metal/GaAs contact. The theoretically barrier heights arecalculated using the MIGS model [Mön95]. The experimentally barrier heights aredetermined using photoemission spectroscopy in the cases of Ag [Vit89a, Bri88], Sb[Cao88], In [Mao92] and Na [Pri89] and I-V measurements in the cases of Ag [Wal84]and Mg [Wal84].

138

barrier height is attributed to an additional hydrogen-induced interface dipoleand is explained in the MIGS-and-electronegativity model. In this case theelectronegativity of the passivating atoms is larger than the electronegativity ofthe substrate atom. This results in charge of positive sign on the semiconductorside of the interface which is compensated by equal charge of opposite sign inthe metal induced gap states [Hei65]. Therefore the Fermi level moves closer tothe conduction band minimum, i.e. in other words, the barrier heights on n- andp-type substrates are decreased and increased, respectively.In chapter 4.2 and 4.3 it has already been described that the chalcogenpassivation of the GaAs surface forms Ga-chalcogen dipoles on top of the GaAssamples. According to chapter 5 it has been observed that the metal depositiononto the chalcogen modified surfaces leaves the Ga-chalcogen dipole layerunaffected with exception of Mg and Na. Nevertheless the change in barrierheight is calculated according to the voltage drop across this dipole double layer.The charge transfer ∆qi between the Ga and the chalcogen atoms can becalculated according to formula 5.3.

2035.016.0 ChalGaChalGai XXXXq −⋅+−⋅=∆ (5.3)

With the electronegativities of Ga (XGa = 1.81 [Pau60]), Se (XSe = 2.55 [Pau60])and S (XSe = 2.58 [Pau60]) the charge transfer ∆qi amounts to 0.138 and 0.144between Ga and Se and between Ga and S, respectively.

⊥⊥ ⋅∆⋅= iii dqep 0 (5.4)

The dipole moment normal to the surface can be calculated according to formula5.4. The distance of the Ga and chalcogen atom normal to the surface isregarded to be quite similar for the Ga-Se and Ga-S bond, respectively. Thisdistance amounts to 0.11 nm [Hir98, Sug97]. Using this value and the chargetransfer ∆qi from Ga to the respective chalcogen leads to a value of 2.43⋅10-30

Cm and 2.54⋅10-30 Cm for the dipole moment normal to the surface for the Ga-Se and Ga-S dipole, respectively.The change in the surface barrier height due to these dipoles can be calculatedaccording to the formula given by Topping [Top27].

ad

ad

i

iB N

N

pe⋅

+⋅±=∆Φ ⊥

2/30

0

91 αεε (5.5)

139

To calculate the change in the surface barrier height the dielectric constant forthe semiconductor interface is needed. For GaAs this has been estimated to εi ≈ 4[Lud88]. With the elementary charge e0, the dielectric constant in vacuum ε0,3.77⋅10-24 cm3 and 2.90⋅10-24 cm3 for the polarizability α of Se and S, and6.26⋅1014 cm-2 for the density of the dipoles at the surface, which coincides withthe density of atoms on the GaAs(100) surface [Mön95], the change in thesurface barrier height ∆ΦB amounts to 0.28 eV and 0.32 eV for Se and Smodified metal/GaAs interfaces, respectively.The barrier height of metal/n-GaAs contacts should be reduced by ∆ΦB, becausethe chalcogen atom of the interface dipole is negatively charged due to its higherelectronegativity while the Ga atom is positively charged. Applying the samearguments the Schottky barrier height of metal/p-GaAs should be increased bythis value. To calculate the Schottky barrier heights of the chalcogen modifiedmetal/GaAs contacts the estimated change in barrier height is subtracted from oradded to the barrier heights calculated by the MIGS model (table 5.5). Theseresults are shown in figure 5.36 by the black lines.

0,1

0,3

0,5

0,7

0,9

1,1

1,3 metal/S:n-GaAsb.)

ΦB

n / e

V

0,1

0,3

0,5

0,7

0,9

1,1

1,3metal/S:p-GaAsd.)

ΦB

p / eV

2,7 2,9 3,1 3,3 3,5 3,7 3,9 4,1 4,3 4,5

0,1

0,3

0,5

0,7

0,9

1,1

1,3 metal/Se:n-GaAsa.)

Electronegativity (Miedema)

2,6 2,8 3,0 3,2 3,4 3,6 3,8 4,0 4,2 4,4

0,1

0,3

0,5

0,7

0,9

1,1

1,3metal/Se:p-GaAsc.)

Electronegativity (Miedema)

Figure 5.36: Schottky barrier heights of a.) Se modified metal/n-GaAs(100) contacts, b.)S modified metal/n-GaAs(100) contacts, c.) Se modified metal/p-GaAs(100) contactsand d.) S modified metal/p-GaAs(100) contacts. The lines represent the Schottky barrierheight calculated according to the MIGS model, which has been diminished (n-doped)or increased (p-doped) by the amount calculated for the Ga-chalcogen dipole.

140

Comparing these theoretically calculated values with the experimentallyachieved data, it is observed that the MIGS model is a good approximation forthe Schottky barrier height of chalcogen modified metal/GaAs(100) contacts.But there is still some discrepancy, which might be explained by additionaladsorbate-induced interface donors or acceptors.

141

6 SummaryIn this study the passivation of clean GaAs(100) surfaces by chalcogens and theeffect of chalcogens on different metal/GaAs interfaces has been investigated.The main results can be summarised as followed:• The clean GaAs surface was prepared by gentle annealing (380 °C) of an As

capped GaAs(100) sample in ultrahigh vacuum leading to an As-rich c(4x4)reconstruction. The reconstruction could be established by LEED andphotoemission spectroscopy of the As 3d and Ga 3d core level spectra. TheFermi level relative to the valence band maximum (VBM) amounts to 0.49eV and 0.34 eV for n- and p-doped GaAs(100), respectively.

• The evaporation of Se or S onto the clean As rich GaAs surfaces at roomtemperature leads to an heterovalent exchange between As and the respectivechalcogen. Also arsenic chalcogenides are formed. At temperatures above300 °C the arsenic chalcogenides evaporate from the surface. The chalcogenmodification of GaAs at temperatures above 340 °C leads to a sharp (2x1)LEED pattern. The surface is covered with chalcogen dimers and also the Asatoms in the third layer have been substituted by chalcogen atoms, accordingto a model proposed by Pashley [Pas84]. This model fits with the resultsfrom the fitting of As 3d, Ga 3d, Se 3d and S 2p core level spectra. Thehigher binding energy component in the chalcogen spectrum is clearlyassigned to surface dimers, while the lower binding energy componentrepresents incorporated chalcogen atoms in the Ga2Se3- and Ga2S3-like layer.

• The surface Fermi level relative to VBM on n-doped GaAs after Se and Smodification amounts to 0.80 eV and 0.90 eV, respectively. On the Se or Smodified p-doped GaAs it amounts to 0.49 eV and 0.51 eV, respectively.Hence the chalcogen modification shows a decreased band bending on n-doped GaAs, while it is increased on p-doped GaAs. That means that newdonor and acceptor type states are created within the band gap, which differfrom the “pinning” states on the clean GaAs(100) surface. That is in contrastto the prediction that the chalcogen modification only leads to a removal ofstates in the bandgap and yields to flat bands on the surface.

• The change in ionisation energy due to the chalcogen modification could bepredicted by assuming a double layer consisting of Ga-chalcogen dipoles onthe surface, which agrees well with the experimentally determined values forS and Se on GaAs.

• On Te modified GaAs surfaces the As 3d and Te 4d core level spectra areseparable by using the Cooper minimum of the Te 4d photoemission crosssection, nevertheless the binding energies of both core levels coincide. TheTe 4d spectrum exhibits similar to the observation of the chalcogen spectrafrom the Se or S modified GaAs surfaces a surface component at higherbinding energy and a subsurface component at lower binding energy. But the

142

Te-As exchange reaction does not take place to the same extent as for Se or Sand hence no effect on the band bending has been observed due to the Temodification.

• The deposition of Mg on Se or S modified GaAs surfaces exhibits a strongreaction. An exchange between Mg and Ga takes place leading to theformation of magnesium chalcogenides. Pure Ga clusters segregate on thesurface due to the dissolution of the gallium chalcogenide layer. Additionallya Mg-As compound segregates on the surface, which indicates a disruption ofthe GaAs bulk. The interface between the reacted layers on the surface andthe GaAs bulk “grows” homogeneously in the direction of the bulk as afunction of Mg coverage. Similar observations have been made for the Nadeposition. After 0.23 nm nominal Mg deposition and after 870 s of Nadeposition a Fermi edge is observed on the samples, respectively, indicatinga metallization of the surface.

• The deposition of In on Se or S modified GaAs leaves the surface nearlyunaffected. The In bonds to the chalcogen dimers on the surface, but does notdisrupt the gallium chalcogenide like layer. This is an improvement incomparison to the In deposition on clean GaAs, where it disrupts the firstlayers of the GaAs bulk. Above 0.24 nm of nominal In coverage the In atomsstart to form clusters and the formation of islands is observed. The clusteringof In atoms is stronger on S than on Se modified GaAs surfaces, whichimplies a higher mobility of In atoms on the S modified GaAs surface.Similar results are obtained for Ag and Sb. Both bond to the chalcogendimers on top of the sample while they do not disrupt the galliumchalcogenide layer. Above 0.12 nm coverage Ag islands are formed on thesurface. Similar to the observation from In the mobility of Ag on S modifiedGaAs is higher than on Se modified GaAs, which is concluded from thestronger clustering of the Ag atoms on the S modified GaAs. For Sb,islanding is observed for a nominal coverage above 0.17 nm. This growthmode is in contrast to the layer by layer growth mode, which was observedfor the Sb deposition on clean GaAs(100). The growth modes of Ag and Inmainly agree with the modes shown on the unmodified surface. The Ag andIn layers exhibit metallic properties at 0.48 nm and 0.60 nm nominalthickness, which is concluded from the occurrence of a Fermi edge. Incontrast the Sb layer does not show any metallic behaviour up to themaximum thickness of 3.3 nm, which means that it remains amorphous.

• The position of the Fermi level on chalcogen modified n-doped GaAsexhibits an overshoot up to 0.2 nm Na, Mg, In or Ag coverage. This resultsfrom adsorbate-induced states of acceptor character. For higher adsorbatecoverage (> 1 nm) the barrier heights of the investigated metal contacts on n-GaAs (p-GaAs) decreased (increased) due to the chalcogen modification incomparison to metal contacts on the clean surfaces. The general trend of thebarrier height is approximated by the application of the metal-induced gap

143

states (MIGS) model and estimating a Ga-chalcogen dipole layer. This dipolelayer leads to a voltage drop, which decreases the Schottky barrier height onn-doped GaAs and increases it on p-doped GaAs.

144

145

AcknowledgementAt this point I want to thank everybody who participated in the finishing of thiswork. My special thanks are for

• Prof. Dr. D.R.T. Zahn for the interesting topic of my thesis and the manifoldhelpful discussions,

• the „Bundesministerium für Bildung und Forschung“ under grant No 05622OCA 3 for the financial support,

• the BESSY staff for their help during the measurements, especially Dr. W.Braun,

• M. Lübbe, A. Schneider, Dr. D. Drews, Dr. T. Werninghaus, Dr. A. Patchettand Dr. C. Schultz for their help during the measurements and helpfuldiscussions,

• Dr. T.U. Kampen for the revision of my thesis,

• A. Fechner and the workshop crew for their advise at the construction ofuseful vacuum parts,

• S. Raschke and the rest of the semiconductor group in Chemnitz for the goodatmosphere and the “manifold” culinary pleasure in the evening,

• Dr. D. Westwood from the Department of Physics and Astronomy of theUniversity of Wales, College of Cardiff for preparing the substrates,

• my wife Christiane without whose help and support this work would not befinished.

146

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Ich erkläre, dass ich die vorliegende Arbeit selbständig und nur unter Verwendung derangegebenen Literatur und Hilfsmittel angefertigt habe.

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Thesen• Es wurde die Passivierung von sauberen GaAs(100) Oberflächen durch

Chalkogene und der Effekt von Chalkogenen auf verschiedene Metall/GaAsGrenzflächen untersucht.

• Die reine GaAs Oberfläche wurde durch vorsichtiges Heizen bis 380 °C vonAs-gecappten GaAs(100) Oberflächen im Ultrahochvakuum präpariert. DiesePräparation führte zu der As-reichen c(4x4) Oberflächenrekonstruktion. DieRekonstruktion wurde durch LEED (Beugung niederenergetischerElektronen) und Photoemissionsspektroskopie der As 3d und Ga 3dRumpfniveaus bestätigt. Das Feminiveau relativ zum Valenzbandmaximum(VBM) lag bei jeweils 0,49 eV und 0,34 eV für n- und p-dotiertesGaAs(100).

• Die Bedampfung der reinen, As reichen GaAs Oberfläche mit Se oder S beiRaumtemperatur führt zu einem heterovalenten Austausch zwischen As unddem jeweiligen Chalkogen. Bei Temperaturen über 300 °C verdampfen dieArsenchalkogenide von der Oberfläche. Die Chalkogenmodifikation vonGaAs bei Temperaturen oberhalb von 340 °C führt zu einem scharfen (2x1)LEED Bild. Die Oberfläche ist mit Chalkogen Dimeren bedeckt und auch dieAs Atome in der dritten Lage wurden durch Chalkogen Atome ersetzt, inÜbereinstimmung mit einem von Pashley [Pas94] vorhergesagten Modell.Dieses Modell stimmt mit den Resultaten aus den gefitteten As 3d, Ga 3d, Se3d und S 2p Rumpfniveauspektren überein. Die Komponente im ChalkogenSpektrum mit höherer Bindungsenergie ist klar dem Oberflächendimerzugeordnet, während die Komponente bei niedrigerer Bindungsenergie dieChalkogen Atome in der Ga2Se3- und Ga2S3-ähnlichen Schicht repräsentiert.

• Das Oberflächen-Ferminiveau auf n-dotiertem GaAs nach Se oder SModifikation beträgt 0,80 eV bzw. 0,90 eV relativ zu VBM. Auf Se- und S-modifiziertem, p-dotiertem GaAs beträgt das Oberflächen Ferminievau 0,49eV bzw. 0,51 eV. Das heißt, dass die Chalkogen Modifikation auf n-dotiertem GaAs eine verringerte Oberflächenbandverbiegung gegenüber derreinen Oberfläche zeigt und auf p-dotiertem GaAs eine erhöhteOberflächenbandverbiegung. Das bedeutet, dass neue Donator- undAkzeptor-artige Zustände innerhalb der Bandlücke erzeugt werden, die sichvon den „Pinning“-Zuständen auf der reinen GaAs(100) Oberflächeunterscheiden. Das ist im Gegensatz zu der Vorhersage, dass die ChalkogenModifikation zu einer zustandsfreien Bandlücke führt und flache Bänder ander Oberfläche erzeugt.

• Die Änderung in der Ionisierungsenergie aufgrund der ChalkogenModifikation konnte durch die Annahme einer Doppelschicht an derOberfläche bestehend aus Ga-Chalkogen Dipolen vorhergesagt werden. Die

170

experimentell bestimmten Werte stimmen mit den für Se bzw. S berechnetenWerten gut überein.

• Die As 3d und Te 4d Rumpfniveauspektren der Te modifizierten GaAsOberfläche konnten aufgrund des Cooper Minimums im Te 4dPhotoemissionswirkungsquerschnitts gut getrennt werden, obwohl dieBindungsenergien der beiden Rumpfniveaus übereinstimmen. Das Te 4dSpektrum zeigt ähnlich zu der Beobachtung der Chalkogen Spektren von derSe oder S modifizierten Oberfläche eine Oberfächenkomponente bei höhererBindungsenergie und eine Komponente von Te Atomen unterhalb derOberfläche bei niedrigerer Bindungsenergie. Der Te-As Austausch findetaber nicht in dem gleichen Ausmaß statt wie bei Se oder S und ebenso wurdedurch die Te Modifikation keine Änderung der Oberflächenbandverbiegungbeobachtet.

• Die Bedampfung der Se oder S modifizierten GaAs Oberfläche mit Mgführte zu einer starken Reaktion. Ein Austausch zwischen Mg und Ga findetstatt, der zur Bildung von Magnesiumchalkogeniden führt. Reine Ga Clusterbilden sich an der Oberfläche durch die Auflösung der GalliumchalkogenidSchicht. Zusätzlich bilden sich Mg-As Verbindungen an der Oberfläche, dieeine Reaktion des Mg mit dem GaAs Rumpf anzeigen. Die Grenzflächezwischen der reagierten Schicht an der Oberfläche und dem GaAs Rumpfwächst homogen in die Tiefe als Funktion der Magnesium Bedeckung.Ähnliche Beobachtungen wurden für die Bedampfung mit Na gemacht. Beieiner Bedeckung von 0,23 nm Mg bzw. nach 870 s Na Bedampfung zeigtsich jeweils eine Fermikante, die die Metallisierung der Oberfläche anzeigt.

• Die Bedampfung von Se oder S modifiziertem GaAs mit In läßt dieOberfläche nahezu unbeeinflußt. Das In bindet sich an die Chalkogen Dimerean der Oberfläche, aber es zerstört nicht die Galliumchalkogenid-artigeSchicht. Dies ist eine Verbesserung im Vergleich zu der In Bedampfung vonreinem GaAs, wobei eine Störung in den ersten Lagen von GaAs beobachtetwurde. Oberhalb von 0,24 nm nomineller In Bedeckung beginnen die InAtom sich zu Clustern zusammenzuschließen und die Bildung von In Inselnwird beobachtet. Das Clustern der In Atome ist auf S modifiziertenOberflächen stärker als auf Se modifizierten, was zu dem Schluß führt, dassdie Mobilität der In Atom auf S modifizierten Oberflächen größer ist.Ähnliche Ergebnisse wurden für Sb und Ag erhalten. Beide binden sich andie Chalkogen Dimere an der Oberfläche, sie führen aber nicht zu einerZerstörung der Galliumchalkogenid Schicht. Oberhalb von 0,12 nmBedeckung bilden sich Ag Inseln auf der Oberfläche. Ähnlich wie bei In istdie Mobilität von Ag auf S modifizierten GaAs größer als auf Semodifiziertem, was aus dem stärkeren Clustern der Ag Atome auf Smodifiziertem GaAs geschlossen wird. Für Sb wird die Bildung von Inselnab einer Bedeckung von 0,17 nm beobachtet. Dieser Wachtumsmodus stehtim Gegensatz zu dem lagenweisen Wachstum, das bei der Sb Bedampfung

171

von reinen GaAs(100) Oberflächen beobachtet wurde. Die Wachtumsmodivon Ag und In stimmen überein mit denen, die für das Wachstum von Agund In auf reinen GaAs Oberflächen beobachtet wurden. Die Ag und InSchichten zeigen metallische Eigenschaften ab einer nominellenSchichtdicke von 0,48 nm bzw. 0,60 nm, wie aus der Beobachtung einerFermikante auf der Probe geschlossen werden kann. Im Gegensatz zeigt dieSb Schicht kein metallisches Verhalten bis zu der maximalen Schichtdickevon 3,3 nm, was bedeutet, dass sie amorph bleibt.

• Die Position des Ferminiveaus auf Chalkogen modifiziertem, n-dotiertemGaAs zeigt eine starke Auslenkung in Richtung des Leitungsbandes bis zueiner Dicke von 0,2 nm der Bedeckung durch Na, Mg, In oder Ag. Dasresultiert aus der Bildung von Adsorbat-induzierten Zuständen mit AkzeptorCharakter. Für höhere Adsorbatbedeckungen (> 1nm) vermindert (erhöht)sich die Barrierenhöhe der untersuchten Metallkontakte auf n-GaAs (p-GaAs) aufgrund der Chalkogen Modifikation im Vergleich zuMetallkontakten auf reinen Oberflächen. Der allgemeine Trend derBarrierenhöhe kann näherungsweise durch die Anwendung des ModellsMetall-induzierter Bandlückenzustände (Metal-induced gap states) und dieAnnahme einer Ga-Chalkogen Dipolschicht erklärt werden. DieseDipolschicht führt zu einem Spannungsabfall, der die Barrierenhöhe auf n-dotiertem GaAs vermindert und auf p-dotiertem GaAs erhöht.

172

173

Curricula vitaeName: Stefan Hohenecker

Geburtstag: 16.05.1968

Geburtsort: Hückeswagen

Staatsangehörigkeit: Deutsch

Familienstand: verheiratet

Schulbildung: 1974 – 1978 Grundschule Görresschule in Neuss1978 – 1987 Quirinus-Gymnasium in Neuss

Wehrdienst: 1987 – 1988 Wehrdienst in Munster und Lüneburg

Studium: 1988 – 1991 Physik-Grundstudium an der Heinrich- Heine-Universität in Düsseldorf, Abschluß: Vordiplom1991 – 1995 Hauptstudium an der Universität Dortmund, Abschluß: Diplom

Berufstätigkeit: 1995 – 1998 Wissenschaftlicher Mitarbeiter im Foschungsvorhaben: Chalkogen- Modifikation von Halbleitergrenz- flächen: die Wechselwirkung von Metallen mit Chalkogen/III-V- und III2VI3/III-V-Grenzflächen am Lehrstuhl für Halbleiterphysik der TU Chemnitzseit 1998 Mitarbeiter der Firma “software design & management sd&m AG”