Atomic-scale studies of hydrogenated semiconductor surfaces

Progress in Surface Science 81 (2006) 1–51 www.elsevier.com/locate/progsurf

Review

Atomic-scale studies of hydrogenated semiconductor surfaces A.J. Mayne *, D. Riedel, G. Comtet, G. Dujardin CNRS, Laboratoire de Photophysique Mole´culaire, Baˆt. 210, Universite´ Paris-Sud, 91405 Orsay, France

Abstract The adsorption of hydrogen on semiconductors strongly modifies the electronic and chemical properties of the surfaces, whether on the surface or in the sub-surface region. This has been the starting point, in recent years, of many new areas of research and technology. This paper will discuss the properties, at the atomic scale, of hydrogenated semiconductor surfaces studied with scanning tunnelling microscopy (STM) and synchrotron radiation. Four semiconductor surfaces will be described – germanium(1 1 1), silicon(1 0 0), silicon carbide(1 0 0) and diamond(1 0 0). Each surface has its particularities in terms of the physical and electronic structure and in regard to the adsorption of hydrogen. The manipulation of hydrogen on these surfaces by electronic excitation using electrons from the STM tip will be discussed in detail highlighting the excitation mechanisms. The reactivity of these surfaces towards various molecules and semiconductor nanocrystals will be illustrated.
2006 Elsevier Ltd. All rights reserved. Keywords: Scanning tunnelling microscopy (STM); Synchrotron radiation; Semiconductor surfaces; Hydrogen; Molecules; Manipulation; Desorption; Electronic excitation; Photon excitation

Contents 1. 2.

*

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Preparation of hydrogenated surfaces on semiconductors . . . . . . . . . . . . . . . . . . . . . .

Corresponding author. Tel.: +33 1 69 15 75 02; fax: +33 1 69 15 67 77. E-mail address: [email protected] (A.J. Mayne).

0079-6816/$ - see front matter
2006 Elsevier Ltd. All rights reserved. doi:10.1016/j.progsurf.2006.01.001

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2.1.

3.

4.

5.

6.

Hydrogenated surface preparation in UHV . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1.1. (1 0 0) Surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1.2. (1 1 1) Surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2. Hydrogenated surface preparation by chemical methods . . . . . . . . . . . . . . . . . . . 2.3. Hydrogen plasma . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Hydrogen desorption methods . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1. Electronic excitation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2. Photonic excitation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3. Thermal desorption . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Experimental and theoretical studies. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1. Hydrogen on germanium(1 1 1) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1.1. Introduction: the clean Ge(1 1 1)-c(2 · 8) surface . . . . . . . . . . . . . . . . . . . . 4.1.2. Adsorption sites of hydrogen on Ge(1 1 1)-c(2 · 8) . . . . . . . . . . . . . . . . . . . 4.1.3. STM manipulation of individual hydrogen atoms . . . . . . . . . . . . . . . . . . . 4.1.4. Reactivity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2. Hydrogen on silicon(1 0 0) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.1. Introduction: the clean Si(1 0 0) surface . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.2. Electronic structure of hydrogenated and partially hydrogenated Si(1 0 0) surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.3. STM desorption of hydrogen atoms . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.4. Desorption of hydrogen on Si(1 0 0):H with lasers . . . . . . . . . . . . . . . . . . . 4.2.5. Reactivity: the adsorption of molecules on Si(1 0 0):H . . . . . . . . . . . . . . . . 4.3. Hydrogen on silicon carbide surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3.1. Hydrogen on hexagonal SiC surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3.2. Hydrogen on cubic SiC surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.4. Hydrogen on diamond surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.4.1. Introduction: properties of diamond surfaces . . . . . . . . . . . . . . . . . . . . . . 4.4.2. Atomic-scale imaging of diamond surfaces . . . . . . . . . . . . . . . . . . . . . . . . 4.4.3. Conductivity of hydrogenated diamond surfaces . . . . . . . . . . . . . . . . . . . . 4.4.4. Desorption of hydrogen from hydrogenated diamond surfaces . . . . . . . . . 4.4.5. Electronic structure of hydrogenated diamond surfaces . . . . . . . . . . . . . . . 4.4.6. Reactivity of hydrogenated diamond surfaces . . . . . . . . . . . . . . . . . . . . . . Hydrogen on other semiconductor surfaces. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1. GaAs(1 0 0) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.2. Ge(1 0 0)-2 · 1 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3. Si(1 1 1)-7 · 7. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3.1. Introduction: the adsorption of molecules on the clean Si(1 1 1)-7 · 7 surface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3.2. Desorption of hydrogen on the Si(1 1 1):H surface . . . . . . . . . . . . . . . . . . . Summary and perspectives . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.1. Preparation of hydrogenated semiconductor surfaces . . . . . . . . . . . . . . . . . . . . . 6.2. Electronic structure and conductivity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.3. STM and laser induced desorption of hydrogen . . . . . . . . . . . . . . . . . . . . . . . . . 6.4. Reactivity of hydrogenated surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.5. Future work . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.5.1. Atomic-scale engineering . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.5.2. Hydrogenated surfaces as substrates for molecular nanomachines . . . . . . . Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

5 5 6 7 7 8 8 9 9 9 10 10 10 12 14 14 14 16 17 20 21 26 26 26 27 27 28 31 32 34 36 37 37 37 38 38 38 39 39 39 40 40 41 41 41 42 42

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Acronyms STM scanning tunnelling microscope SR synchrotron radiation LEED low energy electron diffraction HREELS high-resolution electron loss spectroscopy UHV ultra-high vacuum DAS dimer-adatom-stacking fault L Langmuir ML monolayer ARPES angle-resolved photoemission spectroscopy ESD electron stimulated desorption PSD photon stimulated desorption DB dangling bond TPD temperature-programmed desorption LITD laser-induced thermal desorption SHG second-harmonic generation XRD X-ray diffraction LDOS local density of states ARELS angle-resolved electron energy loss spectroscopy TOF time-of-flight VUV vacuum ultra-violet XPS X-ray photoemission spectroscopy CVD chemical vapour deposition AFM atomic force microscope FTIR Fourier-transform infrared adsorption spectroscopy UPS ultra-violet photoemission spectroscopy IRAS infrared adsorption spectroscopy NEA negative electron affinity RHEED reflection high energy electron diffraction SC surface conductivity NEXAFS near-edge X-ray adsorption fine structure spectra

1. Introduction The hydrogenation of semiconductor surfaces has many important aspects. Hydrogen modifies both the chemical reactivity and the electrical conductivity of the surfaces on which it is adsorbed. Hydrogenated surfaces are passivated, that is, they are less reactive compared to the clean surfaces. The physical properties of these hydrogenated surfaces are also changed; hydrogenation not only removes the surface states by terminating the surface dangling bonds, hydrogen is also present in the atomic layers just below the surface. This changes the electrical conductivity of the surface with respect to the bulk. For diamond, the hydrogenated surface is semiconducting whereas the clean diamond surface is an insulator. However, for example, the clean silicon(1 0 0) surface is semiconducting while the corresponding hydrogenated silicon(1 0 0) surface much less so.

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There are many applications of hydrogenated surfaces. The electrical properties are important because the elimination of electronic defects can improve significantly semiconductor interfaces in microelectronic devices. Also, if zones of clean and hydrogenated surface can be controlled, these surfaces will be suitable for developing electronic contacts on passivated surfaces. The relative chemical inertness of hydrogenated surfaces means that they have the potential to be used in far more areas than the clean surfaces could be, for example, in contact with liquid environments containing biological species. The key objective is to control these hydrogenated surfaces. Can the electronic and chemical properties be controlled and can this be done at the atomic level? The advent of the scanning tunnelling microscope (STM) has made the observation and manipulation at the atomic scale a reality. Combining the STM with other surface techniques which are sensitive to the electronic changes enables a more complete characterisation of these surfaces. Synchrotron radiation (SR) and other electron diffraction methods such as low energy electron diffraction (LEED) and high-resolution electron loss spectroscopy (HREELS) are particularly appropriate. So, the possibilities for applications in nanotechnology, molecular electronics and biology are vast. From this point of view, it is important to review the state of the art. Several reviews have been given over the last 15 years on the interaction of hydrogen with surfaces. Christmann [1] presented a review of hydrogen on metal surfaces with an emphasis on the energetics and kinetics of the adsorption process and the structural properties of the hydrogen adsorbed layers. Schaefer concentrated on the vibrational studies of hydrogenated semiconductor surfaces [2] and aspects of the atomic geometry and electronic structure. Boland’s review concerns the preparation and characterisation of silicon surfaces using STM and the thermal properties of these hydrogenated surfaces [3]. Finally, Oura et al. have made a thorough review of hydrogenated semiconductor surfaces and their reactivity with regard to the adsorption of metals [4], where studies using electron diffraction techniques such as HREELS received particular attention. In this review article, so as not to duplicate previous work, we will concentrate on presenting an overview of atomic-scale STM and SR studies on hydrogenated semiconductor surfaces with a particular focus on manipulating hydrogenated surfaces at the atomic scale. We will consider primarily the semiconducting surfaces of silicon (Si), germanium (Ge), silicon carbide (SiC), and diamond (C). Other surfaces will be mentioned such as GaAs. The preparation of these surfaces will be described in some detail in Section 2, in particular, in ultra-high vacuum (UHV) and using chemical methods. Different surface orientations will also be considered. An introduction to the different desorption methods, that is electronic, photonic and thermal desorption, will be given in Section 3. The reason for this is to prepare partially hydrogenated surfaces in a controlled manner with which molecules can react. This can be used to control the chemical and electronic properties of the surfaces such that the reactions can be very selective. Section 4 will give a thorough examination of the experimental and theoretical studies that have been made. Three important themes will be developed; taking the four most used surfaces in turn: Firstly, the manipulation of surface hydrogen atoms using the STM or synchrotron radiation (SR), secondly, the reactivity of the surfaces with respect to the adsorption of atoms or molecules, and thirdly, the properties of the hydrogenated surfaces such as electrical conductivity and thermal stability. A thorough conclusion will be given in the summary along with perspectives for these surfaces (Section 6).

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2. Preparation of hydrogenated surfaces on semiconductors We will look at the preparation of hydrogenated surfaces of silicon (Si), germanium (Ge), silicon carbide (SiC), and diamond (C). Other surfaces will be mentioned such as GaAs. One question will determine the preparation procedure to be followed: Is the objective a partially hydrogenated or a fully hydrogenated surface? One idea lying behind this question is whether or not interactions between neighbouring hydrogen atoms have an impact on the ensuing surface properties and surface chemistry. This will be expanded on in more detail in the context of STM and SR manipulation in Section 4. Additionally, the crystallographic structure of the semiconductor must also be considered (cubic or hexagonal) as it has an impact on the nature (and quality) of the hydrogenated surfaces. An illustration of this problem is that isolated hydrogen atoms can be easily adsorbed on the Ge(1 1 1) surface in UHV. However, the fabrication of a completely passivated surface with a monolayer of hydrogen does not work in UHV since Ge islands form within the hydrogenated areas. A chemical technique is required to obtain a fully hydrogenated Ge(1 1 1) surface. In contrast, a perfect monohydride Si(1 0 0) surface can be obtained easily in UHV. The three principle preparation methods will be discussed below. 2.1. Hydrogenated surface preparation in UHV This is probably the most common method of preparation for those interested in surface science, particularly if atomic-scale studies such as STM, photoemission, LEED or HREELS experiments are to be carried out subsequent to hydrogenation. Often, the sample is transferred under vacuum into an adjacent analysis chamber which guarantees that the surfaces remain clean at the atomic scale. The hydrogenation procedure is relatively simple in its conception. A hot tungsten filament is placed in front of the clean surface while the chamber is filled with pure H2 gas. The hot metal filament cracks the molecular hydrogen forming atomic H which reacts with the clean surface. The sample is usually heated at the same time. As it is not usually possible to measure the exact dose of hydrogen, the exposure is expressed as the product of the pressure and the time and is given in Langmuir (L) where 1 L corresponds to 1 · 10
6 Torr for 1 s. Another term, frequently used to describe the surface coverage of the adsorbing species, is monolayer (ML). One monolayer corresponds to the occupation of every surface site by an adsorbed atom. There are 6.8 · 1014 atoms/cm2 on the Si(1 0 0) surface. 2.1.1. (1 0 0) Surfaces 2.1.1.1. Silicon. A survey of the literature over the last 15 years shows that the recipes for preparing hydrogenated silicon surfaces in UHV have varied little since the seminal work of Boland [3,5–7] in the early 1990s. The adjustable parameters are the temperature of the tungsten filament, hydrogen pressure, duration, filament–sample distance, and surface temperature. Deuterium covered surfaces have been studied to a much lesser extent [8]. In general, in this review, the discussions will focus almost entirely on hydrogenated surfaces while reference will be made to deuterated surfaces where necessary. In Table 1, the different parameters are listed from various references in the literature. So as not to disrupt the flow of this review, the references in the table have been added a posteriori. One can see from the table that certain parameters vary over a seemingly large range. Experimentalists

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Table 1 A non-exhaustive selection of parameters used in the preparation of the Si(1 0 0)-2 · 1:H surface from various references in the literature (in chronological order) Year

1983 1984 1990 1993 1994 1996 1999 2003

Author

Maruno Ciraci Boland Jiang Lyding Hashizume Lin Soukiassian

T(filW) Distance H/D Exposure (C) (cm) Pressure Time (s) 1700 2000 1523 1800 1500 1500 1523 1600

20 5 2 8 6 3 5 3

H H H D H H H H

7 · 10
6 360 1 · 10
6 60–1000 6 · 10
8 60 1 · 10
6 2 · 10
8 1 · 10
6 2 · 10
6

Langmuir (L)

2500 60–1000 3.6 <10,000 400–1200 400–1200 600 12 200 200 700 1400

T(Si) Reconstru- Reference (C) ction

350 250 377 127 377 390 317 370

2·1 2·1 2·1 3·1 2·1 2·1 2·1 2·1

[40] [41] [5] [8] [90] [105] [47] [101,102]

faced with this, are left to choose the parameters at their convenience and to make adjustments. We found that the critical moment occurs when terminating the hydrogenation. In the case of silicon(1 0 0)-2 · 1, in particular, one has to stop heating the sample, then stop the hot filament and finally shut off the hydrogen to the chamber. If all three were stopped at the same time, half the hydrogen adsorbed on the surface would be thermally desorbed since the silicon surface was at 375 C. On the other hand, if more than 30 s elapses between cooling the sample and cooling the hot filament, the 3 · 1 reconstruction forms in significant quantity (since this is stable at a lower temperature) rather than the 2 · 1 surface which was desired. 2.1.1.2. Cubic silicon carbide. Note that the hydrogenated hexagonal SiC(0 0 0 1) will be discussed in Section 4.3. The 1 0 0 surfaces of cubic silicon carbide have been the source of much interest over recent years by virtue of the importance of this material in high temperature and high power electronic components [9,10]. The various surface reconstructions are quite complex due to the possibility of varying the surface concentration of silicon or carbon depending on the preparation conditions. Though the various reconstructions of these surfaces have been thoroughly studied using STM and other surface-sensitive techniques [11–17], the hydrogenation of silicon carbide has been achieved only recently [18,19] and only on the silicon-rich SiC(1 0 0)-3 · 2 surface. Again atomic hydrogen is produced by cracking the molecular hydrogen with a hot tungsten filament placed in front of the SiC surface. If the SiC surface is maintained at room temperature, a 3 · 1:H phase is observed in LEED and angle-resolved photoemission spectroscopy experiments (ARPES) [18]. On the other hand, if the SiC surface is heated to 300 C, the 3 · 2 structure is kept while the surface becomes metallic [19]. The STM images show that the SiC surface is much more reactive than that of silicon as demonstrated by the small quantities of hydrogen needed. Partial coverage is achieved with 10 L of H2 and an almost complete monolayer with 20 L of H2 [19]. This reactivity is probably due to the highly strained surface. 2.1.2. (1 1 1) Surfaces 2.1.2.1. Germanium. The hydrogenation of germanium is not so simple. A complete uniform and well-ordered monolayer is difficult to achieve. As Boland showed [20], prolonged exposure of the c(2 · 8) surface [21] to hydrogen in UHV results in hydrogenation of the

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rest layer forming a 1 · 1 structure with clusters of atoms from the ad-layer dispersed across the surface. This is because the hydrogenation is driven by the thermodynamics of the reaction and the relaxation of the surface. Weak or strained bonds are broken first. In order to facilitate the desorption of GeH2 and GeH3 species, the sample must be heated to around 225 C while exposed to the atomic hydrogen flux [22]. The surfaces prepared in this manner show clean 1 · 1 LEED patterns which is an indication though not a proof that the surface is well reconstructed at the atomic scale. If a sub-monolayer coverage is desired, then the hot filament method works well. As we have shown [23], a small exposure of 1–10 L of hydrogen using a hot filament is sufficient to adsorb isolated H atoms on the Ge surface without any major disruption of the c(2 · 8) surface reconstruction. In fact one can use the ion gauge as a hot source to crack the hydrogen if a very low coverage is required [24]. 2.1.2.2. Silicon. Silicon(1 1 1) surfaces have been far more widely studied than those of germanium, in particular the 7 · 7 reconstruction. It should be noted that the 7 · 7 surface reconstruction extends through four layers based on the model dimer-adatom-stacking fault (DAS) structure proposed by Takayanagi [25]. This surface is different in that the terminating adatoms form a surface state at the Fermi level which gives this surface a metallic character. Under UHV hydrogenation attacks the surface layers and different stoichiometries are observed as the hydrogenation proceeds. As a first step, atomic hydrogen reacts with the silicon surface saturating the adatom and rest atom dangling bonds [26]. Exposure of this hydrogenated surface to more atomic H results in the etching of the surface layers and the formation of SiH2 and SiH3 species on the surface. Heating the surface during the hydrogenation is not sufficient to remove all the SiHx species and so these surfaces are not perfect due to the SiHx species remaining on the surface forming adatom islands [27,28]. A more detailed description of the processes involved and the different techniques used to understand the interaction hydrogen with the Si(1 1 1) surface can be found in Oura’s review [4]. 2.2. Hydrogenated surface preparation by chemical methods Chemical preparation of hydrogenated surfaces is most often used in applications such as the formation of interfaces, deposition of molecular organic layers (Langmuir Blodgett films) and experiments in liquid or ambient environments. One advantage of the wet chemical technique is that the passivated surface can be studied by a battery of different techniques that do not have to be UHV compatible. The passivation of silicon surfaces is based on a series of oxidation and stripping cycles using hydrogen peroxide/sulphuric acid and HF/water, respectively. This has been a speciality of Chabal’s group [29–31]. Studies have shown that the SiH antibonding and bonding states are roughly 3 eV above and below the Fermi level, respectively [32,33]. It is this large gap that makes the surface passive and stable. In fact, it has been realised for some time that the passivation can be used to eliminate electronic defects at interfaces more effectively than for the Si/SiO2 interface [34]. 2.3. Hydrogen plasma A hydrogen plasma is required to produce hydrogenated surfaces under certain circumstances – usually when the two previous methods fail. This applies especially to

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diamond surfaces. The procedure involves the formation of a hydrogen plasma in contact with the sample by microwave heating of hydrogen gas at a relatively high pressure [35]. Three necessary processes occur during the plasma treatment; the hydrogen etches the surface – it removes the roughness, it removes graphite-like carbon, and it passivates the surface [36]. In fact, the hydrogenated diamond surfaces prepared in this manner are quite stable in air for several weeks. In our experiments, we used the following conditions: microwave power, 950 W at 2.65 GHz; H2 pressure, 60 Torr; H2 flow rate, 200 sccm; duration, 1 h. The plasma forms just above the sample surface which heats it to about 800 C. This is followed by cooling the sample down before stopping the flow of hydrogen and transferring the sample [37,38]. These seem to be fairly standard conditions, although the parameters can be varied to some extent, for example, Williams et al. [39] used the following conditions: 2.45 GHz, 800 W, 40 Torr, 5 min, with a sample temperature of 500 C. 3. Hydrogen desorption methods In this section, we will outline three methods for desorbing hydrogen atoms from the semiconductor surfaces. The reason for this is to prepare partially hydrogenated surfaces in a controlled manner with which molecules can then react. This can be used to control the chemical and electronic properties of the surfaces such that the reactions can be very selective. The following questions, amongst others, will be touched upon: What are the products of the desorption method? Is the desorption controllable? The desorption mechanisms are not the same depending on the method used. For example, single hydrogen atoms can be desorbed using electrons or photons but thermal desorption desorbs only pairs of hydrogen atoms. The desorption mechanisms will be discussed in detail for each surface in Section 4. 3.1. Electronic excitation There are several ways of inducing hydrogen desorption by electronic excitation. Electron stimulated desorption (ESD) has been used to characterise the electronic structure of surfaces through the desorption of hydrogen and other species. However, it has not been used with the direct intention of manipulating atoms. ESD does have the advantage of being more energy selective – one can tune the energy window of the electrons impacting the surface. Also one can choose the incident angle of the incoming electrons with respect to the surface normal and achieve an angular resolution. However, the zone irradiated can vary from a few 100 nm2 to several mm2, depending on the ability to focus the electron beam. The limited spatial resolution means that the energy spectrum is an average of several thousands of desorption events. The STM has the advantage of providing an atomic-scale source of electrons and the result of the desorption event can be seen with an atomic resolution – one can selectively remove a single hydrogen. However, it is more difficult to select an energy window; electrons have any energy between the Fermi level of the tip and the surface (which is determined by the applied bias). The majority of electrons leave the tip at the Fermi level and so the effective energies are that of the applied bias. Also angle-resolved studies are not possible; electrons arrive from the tip normal to the surface.

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3.2. Photonic excitation Photon stimulated desorption can be achieved using either a synchrotron radiation source or lasers. While the size of the photon spot limits the spatial localisation of the irradiation, bond breaking by photons has several advantages. Synchrotron radiation has been used for the most part to characterise the electronic structure of surface over a certain energy range (10 to a few 100 eV). However with the new generation of synchrotrons, one should have enough light flux to allow a small energy range to be selected. This is the case for lasers where one can select a photon energy with a small width and thus probe particular excitations. The photon beam can easily be polarised and hence selectively excite bonds having a certain angle with respect to the surface normal. Finally, one can play with the temporal nature of the laser beam and so investigate the dynamics of the excitation reaction. Taking the example of the Si–H bond, early studies have indicated that the localised electronic transition at 8 eV corresponds to the Si–H r–r* transition [40,41]. This has led to a number of studies in the hope of evidencing a direct electronic excitation. Experiments indicate that this is indeed the case both on the hydrogenated Si(1 1 1) [42] and Si(1 0 0) [43] surfaces. A more detailed treatment will be given in Section 4.2. 3.3. Thermal desorption As the heading suggests, hydrogen can be desorbed from a semiconductor surface by heating the sample. Thermal desorption is obviously not a local technique as the entire sample is heated. In STM images of the hydrogenated Si(1 0 0)-2 · 1 surface isolated silicon dangling bonds (DB) are observed. Even at room temperature these are observed to hop – in fact it is the hydrogen next to the DB that hops and this can be facilitated by choosing the appropriate tunnelling conditions [44]. Thermal annealing to around 300–350 C can induce diffusion. It is found that the number of isolated DBs is greatly reduced; they have paired up. This is driven by the restoration of the more energetically stable clean silicon p bond [45]. This has also been observed directly for the deuterated Si(1 0 0) surface [46]. Annealing to higher temperatures (450–500) results in the desorption of hydrogen in pairs [45]. Again, this is driven by the reformation of the p bond of the Si–Si dimer. The kinetics of the thermal desorption process has been studied using STM [47]. As an alternative to temperature-programmed desorption (TPD), hydrogen can be removed using laser-induced thermal desorption (LITD) [48–50] or second-harmonic generation (SHG) [51]. In contrast to direct photon excitation of the semiconductor–hydrogen bond described in the previous subsection, LITD and SHG use lower energy photons (<2.0 eV) that are adsorbed by the substrate which causes local heating and hence thermal desorption of the hydrogen. 4. Experimental and theoretical studies We will now examine the experimental and theoretical studies that have been made on hydrogenated surfaces. Three important themes will be developed, though not necessarily in the same order, for the four most commonly used surfaces (Si, Ge, SiC, and diamond). These themes are: (i) the manipulation of surface hydrogen atoms using the STM or synchrotron radiation (SR), (ii) the reactivity of the surfaces with respect to atoms or molecules, and (iii) the properties of the hydrogenated surfaces such as electrical

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conductivity and thermal stability. The theoretical aspects will focus on the atomic and electronic surface structure, the electrical properties of the surfaces, the properties of nanostructures created by manipulation, and the adsorption configurations of adsorbed atoms or molecules. 4.1. Hydrogen on germanium(1 1 1) 4.1.1. Introduction: the clean Ge(1 1 1)-c(2 · 8) surface A brief description of the structure of the Ge(1 1 1)-c(2 · 8) surface will be given. For a more detailed discussion see the review on semiconductor surfaces by Kubby and Boland [52]. Since the first STM images were obtained of the c(2 · 8) reconstructed surface [53], a debate followed as to the local atomic surface structure. Angle-resolved photoemission spectroscopy (ARPES) and X-ray diffraction (XRD) measurements both suggested an adatom structure [54,55] but without the presence of dimers as had been proposed for the Si(1 1 1)-7 · 7 reconstruction [25]. This suggested a simple adatom layer directly on top of a bulk layer. This was confirmed by detailed calculations [56–58]. The clean surface is best prepared by ion sputtering the Ge surface at 640 C with 1 kV Ar+ ions [21,54,55]. By cooling the surface down very slowly, large single domains of up to 100 · 100 nm can be obtained. Under good UHV conditions the Ge(1 1 1) surface remains clean for several days. This is due, in part, to the relatively low reactivity of the surface to adsorbing species compared to the similar silicon surfaces. For example, in order to achieve a substantial sub-monolayer coverage of oxygen on the surface, more than 100 L is required on Ge(1 1 1) [59] whereas 1 L is enough on the Si(1 1 1) surface [60]. This is probably due to the difference in electronic structure of the surface states between the two surfaces. The Ge(1 1 1) surface has restatom and adatom states about 0.7 eV and 1.3 eV below the Fermi level, respectively [54,61]. The Si(1 1 1) surface has adatom states at the Fermi level which facilitates the adsorption of molecular oxygen [62]. 4.1.2. Adsorption sites of hydrogen on Ge(1 1 1)-c(2 · 8) The adsorption of hydrogen on the Ge(1 1 1)-c(2 · 8) surface has been studied using a variety of techniques. Most of the earlier studies have focused on the electronic structure of the hydrogenated surface using photoemission [61] or infrared methods [64]. The general conclusion was that a new state was observed 5 eV below the Fermi level corresponding to the formation of the Ge–H bond [61]. The bond energy of Ge–H is around 3.0 eV [63]. The mono-hydride (Ge–H) could be preferentially obtained at elevated temperatures (>150 C) while the mixed hydrides (GeH2 and GeH3) were formed at lower temperatures (<
100 C) [64]. Both these studies used very large quantities of hydrogen which severely disrupted the Ge surface. Given the STM images observed under similar conditions [20], it is not surprising. Few studies have looked at the low coverage regime. An EELS study combined with TPD found that hydrogen causes the work function to change significantly at low coverages [65]. They found that the work function decreased by
0.05 eV for a coverage of less than 0.05 ML and then increased by 0.1 eV as a further 0.2 ML was added. The STM study by Klistner and Nelson was the first to show that the adsorption of hydrogen takes place on the restatom sites and not on the adatom sites [66]. In addition they suggested that the adsorption of hydrogen causes a very local transfer of charge from the restatom to the nearest-neighbour adatoms. This is the reason for the characteristic triangle and

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square sites (three and four nearest-neighbour adatoms) that are observed in the STM images (Fig. 1) [23]. In addition to the adsorption of hydrogen on the restatom sites (triangle and square), we found a new site (called a zip site) which was the result of a local rearrangement of the c(2 · 8) structure [67]. This zip site could be selectively created at low voltage by the displacement of a hydrogen atoms rather than its removal [24,67], provided there was another hydrogen atom along the same row of atoms. From the observation on a macroscopic scale of the change in work function [65], hydrogen can cause either an increase or a decrease of the work function. It is important to realise that the work function is a thermodynamic property which is defined at a macroscopic and mesoscopic

Fig. 1. STM topographs of a Ge(1 1 1)-c(2 · 8) surface showing; in (a) a triangle and (b) a square site of a hydrogen atom adsorbed on a Ge rest atom surrounded by three and four adatoms, respectively, and in (c) a zip site produced by two hydrogen atoms. The left-hand STM topographs show the unoccupied states (I = 1 nA, Vs = 1 V) and the corresponding right-hand STM topographs show the occupied states (I = 1 nA, Vs =
1 V) of the same areas. To the right, the schematics of each site are shown where the Ge adatoms and rest atoms are represented as large and small light circles, respectively. Adsorbed H atoms are shown as dark circles and the effective charges are also indicated [67].

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scale and therefore is normally considered to be non-local [68]. At the atomic scale the barrier height is modified, there is local electrostatic potential which could be considered as a localised variation in the work function. We wished to explore this concept using the STM at the level of a single atom [67] and in this discussion will use the term ‘‘local work function’’ as in the article rather than ‘‘local electrostatic potential’’. By considering the change in the effective charge on the various surface atoms, one can understand the redistribution of the local work function at the atomic scale. The adsorption of a hydrogen atom on a Ge restatom, giving rise to a triangle or square site, is equivalent to a positive charge located at the restatom site [67,69]. This effective positive charge induces a local screening by negative charges and consequently a local decrease of the work function [67,68]. However, for the zip site, the hydrogen atoms break adatom back-bonds. This creates two equivalent negative charges at either end of the zip site which induces a local screening by positive charges, resulting in an increased local work function. We observed that at very low coverage, only triangle and square sites were visible but as the quantity of hydrogen increased above 0.05 ML, the zips sites became predominant. Thus the STM study enabled us to show that the different hydrogen sites were the cause of the change in work function. We were also able to show that by careful manipulation of the hydrogen, the local work function could be toggled [67] (see Section 4.1.4 below). The substrate temperature also has an effect on the STM topography of the hydrogen sites. We found in the STM images that the H sites did not appear the same at low temperature (30 K); the three or four adatoms surrounding the H have the same intensity as the other surface adatoms [24]. This suggested that the surface states lying within the bulk band gap are decoupled from the bulk and frozen as the temperature is lowered whereas the tunnel electronic channels lying outside the bulk band gap seem unaffected by the drop in temperature. Furthermore, the zip sites did not form at low temperature and could not be induced even by manipulation suggesting that the formation of the zip site requires thermal as well as electronic activation. Detailed STM spectroscopy measurements by Feenstra have confirmed the influence of the temperature on the surface states [70,71]. This can be explained in part by tip-induced band bending and by non-equilibrium effects created by local disorder such as defects or hydrogen adsorption sites. 4.1.3. STM manipulation of individual hydrogen atoms On the Ge(1 1 1)-c(2 · 8) surface, at low surface coverages, isolated individual hydrogen atoms are adsorbed. This is a very different situation from the Si(1 0 0)-2 · 1:H surface where every surface silicon atom is bonded to a hydrogen atom. There are several advantages in studying the desorption of isolated hydrogen atoms. First, we can be sure that only one hydrogen atom is removed at a time. Second, any influence of the presence of neighbouring hydrogen atoms on the desorption processes can be ignored which is not the case for hydrogen on silicon. For example, electronic coupling or steric interaction between hydrogen atoms could play a role. The manipulation procedure can be operated in two ways: the first is in ‘‘constant current’’ mode and the second is in ‘‘constant height’’ mode. The first method works well if there is a significant change in the local density of states (LDOS) between the adsorbed atom site and the clean surface. This is the case for hydrogen on silicon and will be discussed in the next section. In the constant height mode, the STM tip is placed above the hydrogen atom under computer control, the feedback loop is opened, the tip is dis˚ or approaching) from the surface placed a fixed vertical distance (either retracting 2–20 A

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Fig. 2. The diagram shows a section of the pulse applied to an adsorbed hydrogen atom. The ‘‘write’’ pulse is 1 ms, the ‘‘read’’ pulse is 10 ms, by applying a voltage of +6 V, the current required to remove the hydrogen during the ‘‘write’’ pulse is 3 nA. The ‘‘read’’ pulse corresponds to the initial conditions of
1 V and 1 nA [72].

and a positive sample bias (2–10 V) applied for a short time period (200 ms) during which the tunnel current is measured. After this, the applied bias is removed, the tip is returned to its initial z position, the initial bias reapplied and the feedback restored. The surface is then imaged to verify that the extraction of the hydrogen has taken place and the process repeated on another H atom. The removal of hydrogen from the Ge surface gives rise to a small change in the current during the pulse which was not possible to detect using the constant current mode and difficult using the constant height mode. So a modified constant height mode had to be employed. A pulse composed of a train of about 100 short ‘‘write’’ and ‘‘read’’ pulse was created. Between each ‘‘write’’ and ‘‘read’’ pulse, the feedback loop was not restored, so that after the ‘‘write’’ pulse, a ‘‘read’’ pulse was made by applying the initial voltage bias and the current measured permitting the removal of the H atom to be clearly detected. By applying a sequence of a hundred or so pulses (duration 0.1–2 ms), each followed by a measure of the current at the image bias, before re-establishing the feedback loop, it was possible to deduce the current required and the time at which the change occurred. A section of such a sequence of pulses is shown in Fig. 2 where the removal of the hydrogen occurs during the fourth pulse [67,72]. Detailed studies showed that the inelastic electronic excitation of the Ge–H bond was strongly voltage dependent [23]. The probability of removing a single hydrogen atom was close to 1 for voltages above 5 eV and decreased sharply for voltages below 5 eV. Given that the duration of the voltage pulses was very long compared to the characteristic extraction time, we expected to detect all possible events. This could be explained if the inelastic interaction at low voltages is considered to be very localised. By assuming a lateral spread of the inelastic interaction having a Gaussian distribution, we could determine that in the field emission regime (above 4.5 eV), the spread was about 150 pm whereas in the tunnel regime, the tip needed to be positioned to within 20 pm (at 2 V) of the centre of the hydrogen in order to efficiently remove it. When the tip is moved away laterally from the centre of the H site, the total tip–surface interaction is not strongly modified since the electrons can tunnel through the elastic channels of the substrate. However, the tip interaction with the localised inelastic channels of the Ge–H bond decreases exponentially as the tip moves away laterally.

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In this study [23], we found that the desorption yield increased for sample voltages above 4 eV. The desorption yield was found to be independent of the tunnel current and of the tip–surface distance. This suggested that we were observing a pure electronic excitation not involving multi-vibrational processes (unlike silicon) and that the electric ˚. field could be neglected. We estimated that the electric field never exceeded 0.3 V/A Furthermore, the Ge–H bond energy of 3.0 eV [63] is low enough that a single electron is sufficient to break the bond in the voltage range 3–10 V. The known electronic excitation processes are resonant electron attachment of an electron in the r*(Ge–H) antibonding orbital (unoccupied) at 3.5 eV above the Fermi level [61] and the direct r–r* (Ge–H) transition at 8.5 eV [65]. To explain the increase in the yield above 4 eV, we have to consider that the tunnel electrons attach to the r* Ge–H orbital. The inelastic electronic channels were found to have the same efficiency at 300 and 30 K for the vertical manipulation (removal) of hydrogen atoms [24]. However, when the hydrogen displacement is more complex such as the formation of the zip site discussed above, the inelastic electronic channels are completely inhibited at 30 K. 4.1.4. Reactivity The manipulation of hydrogen was then used to demonstrate the promotion of the surface reactivity at a selected site [67]. This takes advantage of the modification of the local work function by the hydrogen atoms. It is well known that the adsorption of promoter atoms can enhance the surface chemical reactivity [73]. For example, sodium atoms on the Ge(1 1 1) surface have been used to render the surface reactive to oxygen [74] which is normally completely inactive [59]. In our experiment, a hydrogen atom was transferred from the surface to the tip using a positive sample voltage and then, after displacing the tip ˚ , re-deposited at a selected restatom site (by reverslaterally across the surface some 100 A ing the bias). This site is reactive to oxygen due to the increased local density of states on the adatoms surrounding the Ge–H restatom site. After an exposure to 20 L of O2, an adatom next to the H site had reacted while the clean surface remained intact (Fig. 3). This shows that the modified electronic structure of the surface due to the presence of the H atom is very localised. 4.2. Hydrogen on silicon(1 0 0) 4.2.1. Introduction: the clean Si(1 0 0) surface Silicon is the primary material used in semiconductor technology and has been studied widely for many years. As such, silicon surfaces have been extensively studied with an emphasis on the surface structure and chemical reactivity. The reader is advised to consult the review of Waltenberg and Yates on silicon surfaces [75]. Their review is an ideal starting point for an introduction to the surface structure and thermal processes involving the adsorption and desorption of a variety of small molecules. The review contains a very thorough bibliography. Only a brief sketch of the clean silicon surface is given here. The clean silicon(1 0 0) surface has a 2 · 1 periodicity and is composed of rows of silicon dimers. The first LEED observations led to the proposal of the dimer model [76,77] which was confirmed by later experiments [78,79] and calculations [80]. The first STM images of the silicon dimers [81] simply added to the ongoing debate as to whether the dimers were symmetric or asymmetric at room temperature [82,83]. With the ability to study now semi-

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˚ 2) of a Ge(1 1 1) surface Fig. 3. Series of STM topographs (I = 1 nA, Vs = 1 ) showing the same area (165 · 90 A after (a) adsorption of an hydrogen atom into a triangle site, (b) controlled displacement of the hydrogen atom to another predetermined triangle site, and (c) exposure to 20 L of oxygen. The natural dislocation in all three topographs serves as a marker [67].

conductor surfaces at very low temperatures (5 K), the debate has now shifted to determining the most stable structure at these low temperatures [84,85]. The clean surface is produced easily in UHV by degassing the sample at 700 C for at least 12 h followed by a series of short flashes to temperatures in the range 1050–1120 C. It is better to flash p-doped samples at the lower end of this range and n-doped samples at the higher end. This is because p-doped samples tend to form facets on the surface more easily. These show up as large pyramids on the surface in the STM images. Silicon surfaces are very sensitive to the residual vacuum; the cleanest surfaces are obtained with around 1% defects if the initial background pressure in the vacuum chamber is in the 10
11 Torr range and the pressure rise is kept well below 1 · 10
9 Torr. Hydrogenation of the clean surface is done using a hot filament to crack molecular hydrogen which reacts with the heated surface as described in Section 2. For most experiments where the hydrogen has been manipulated using STM or synchrotron radiation, the 2 · 1 mono-hydride phase is used since the 2 · 1 phase corresponds to 1 ML coverage and all Si–H sites are equivalent.

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This is not the case for the 3 · 1 structure where rows of mono-hydride and di-hydride are formed which can have an impact on the desorption processes. 4.2.2. Electronic structure of hydrogenated and partially hydrogenated Si(1 0 0) surfaces Non-local techniques such as synchrotron radiation (SR), ESD, and EELS, are able to give complementary information on the electronic structure of surfaces vital to the understanding of the physical processes involved in hydrogen desorption. In particular, the thermodynamics and kinetics of the hydrogen desorption can be obtained. When the techniques can be combined with STM either directly or indirectly, they become powerful methods of understanding and controlling the manipulation of hydrogen. HREELS and angle-resolved electron energy loss spectroscopy (ARELS) give us an insight into the electronic structure of the surface. In particular, they are sensitive to the vibrational coupling and plasmon modes. Early measurements [40] concentrated on the structure of the hydrogenated surface and showed that the Si DBs on the clean surface are occupied by the hydrogen in the mono-hydride phase and that the hydrogenated dimers are symmetric in contrast to the asymmetric dimers on the clean surface. More recent HREELS measurements have shed more light on the inelastic scattering of electrons with the Si–H bond and the underlying silicon [86]. It was found that the bending and scissor modes of the adsorbed hydrogen are excited predominantly by dipole scattering. However, the stretching modes showed an increase in intensity at 3 and 5 eV. These were interpreted as an enhancement of the inelastic vibrational loss due to a negative ion resonance or impact scattering, respectively. The trapping of an electron in the r* orbital forming a negative ion resonance leads to an enhancement of the energy loss. The electronic structure of the partially hydrogenated Si(1 0 0) surface has been studied using synchrotron radiation [87]. The partially hydrogenated surfaces were prepared using photon stimulated desorption (PSD) with unmonochromatised SR (up to 250 eV) and by controlled thermal desorption following the procedure applied by Boland [3,45] – a 1 min anneal at 700 K. Valence-band photoemission spectra were then recorded at a photon energy of 40 eV and normalised to the beam intensity. The spectra of the thermally annealed surface show a decrease in the intensity of the Si–H bands (at 3.6, 4.2, 5.8, and 10.1 eV) as a function of the annealing temperature combined with a simultaneous appearance of surface states associated with the silicon dangling bonds (Fig. 4). The spectra could be fitted well with a surface state at 1.0 eV, even at high but non-saturation coverage, corresponding to the p orbital of isolated Si dimers [87]. These dimers should be free from any interdimer coupling effect which is observed on the clean surface [88,89]. The spectrum of the thermally prepared surface could not be fitted if the peak corresponding to the isolated DBs was included. This can be explained by the fact that during the cooling down process hydrogen diffusion occurs causing the dangling bonds to pair up and reform the more energetically favourable p-bonded Si–Si dimer. This is in agreement with studies where the dangling bonds have been counted using STM [3,45]. Indeed, in the STM images only clean Si dimers were observed. Using variable-temperature STM, dangling bond pairs have been observed to break up and diffuse across the surface before recombining at a later time [46]. A statistical analysis of STM images obtained as a function of time at elevated temperatures [47] showed that DB pairs can attract one another but that entropy (due to the temperature) drives them apart. The spectra of the partially hydrogenated surface prepared by PSD [87] showed a state at 0.7 eV below the Fermi level in addition to the peak at 1.0 eV. This additional state was

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Fig. 4. Photoemission spectra of the Si(1 0 0)-(2 · 1):H surface as a function of (a) annealing temperature and (b) irradiation dose (annealed clean surface is shown for comparison). The irradiation dose was calculated by multiplying the irradiation time with the photoemission current of a gold mesh monitoring the photon flux [87].

observed only on the irradiated surface. This could only be explained by the presence of isolated silicon dangling bonds on the Si(1 0 0):H surface. An analysis of the relative intensities of the bands as a function of hydrogen coverage could be fitted very well by assuming a random distribution of isolated DBs, pairs of DBs (dimers) and coupled dimers. Irradiation of the clean surface induces no heating of the surface so diffusion of the dangling bonds cannot occur. 4.2.3. STM desorption of hydrogen atoms The hydrogenated silicon(1 0 0) surface is the most studied semiconductor surface. This is due to the importance of silicon in microelectronics, the relative ease of preparation of the hydrogenated surface and the nature of the Si–H bond. STM desorption of hydrogen atoms has been the subject of much study over the last 15 years. There are three main methods for desorbing the hydrogen with electrons from the STM tip. In the constant current mode, the STM tip is placed above the hydrogen atom whereupon both the set-point voltage and current are changed to a value different (usually larger) from the initial setting (for example,
2 V and 1 nA to +3 V and 10 nA). Since the feedback loop is on, the tip– surface distance changes. After a fixed time, the tip is returned to the initial voltage and

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current settings. The desorption event is detected by the presence of a sharp peak (or depression) in the trace. At this moment, an abrupt change in current induces the feedback loop to adjust the tip–surface distance to re-establish the fixed current value. The constant height mode is not so appropriate on the Si(1 0 0):H surface as the tunnel current increases greatly when the hydrogen is removed and since neighbouring hydrogen atoms are under the tip, multiple desorption events can occur during the same pulse. This makes the interpretation of the results particularly difficult. Finally, rather than placing the tip over a single hydrogen and applying a pulse, the tip can be scanned over the surface in the constant current mode with modified voltage and current settings (with respect to the imaging conditions) thereby tracing lines on the surface. This was the first method to be widely used ever since the initial experiment by Lyding et al. [90]. The advantage of this method is that it is quicker and so patterns can be created easily on the surface. The early studies on this surface were pre-occupied with understanding the desorption mechanism. The results showed the presence of two desorption regimes with a threshold around 6.5 eV. The high energy regime corresponds to a direct r–r* excitation of the Si–H bond [91,92]. Indeed, the r-bond of Si–H can be easily broken and the desorption yield is high (2.4 · 10
6 H-atoms per electron). This is related to the long vibrational lifetime of the excited state of the Si–H bond on the order of 10 ns [93] and is due to the fact that the Si–H bond can only relax via phonon coupling to the surface [94] whereas on metals, electron–hole pair formation is more favourable and therefore more rapid. In the low voltage regime, a very different behaviour was observed. The desorption yield was several orders of magnitude lower and more importantly manifested a power law with the current and voltage applied during the pulse. Between 2 V and 4 V, the yield increased by more than a factor of 10 as the current was increased from 1 to 3 nA. The current dependence was explained by the desorption resulting from an electron attachment in the r*(Si–H) orbital via vibrational heating of the Si–H bond where 10 or more electrons were needed – each electron giving only a small fraction of its energy to the Si–H bond (1 quanta) [90,95–97]. The same mechanism was evoked for hole excitation [98] and the behaviour of the Si–H desorption as a function of temperature [99,100]. We carried out studies using both the stationary mode and line scan desorption methods over a larger current range (1–10 nA) and with more data points [101]. Since the energy of the Si–H bond is around 3.5 eV, we chose to apply a voltage below this (2.5 eV) so as to be sure that two or more electrons were involved. The desorption yield showed the same weak current dependence for both n- and p-type samples and for both the stationary and line-scan modes (Fig. 5). The desorption yield had the same absolute value as that found in previous studies but we found a very small dependence of the desorption yield on the current where only two electrons are required to break the Si–H bond [101]. In fact, the line scan method used by Shen et al. [95] is not without a number of practical complications. The whole line receives a certain dose of electrons so depending on the speed, several electrons can interact with a single hydrogen atom or between hydrogen atoms, which renders an understanding of the physical process of desorption more difficult. In a further study, we found that the role of the tip is important. A large number of hydrogen atoms are removed in a short space of time so the tip is easily passivated which can change the efficiency of the desorption process [102]. We observed that the lines drawn by the tip were segmented where from time to time no hydrogen was removed. It was as if the tip had ‘‘on’’ and ‘‘off’’ modes.

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Fig. 5. Desorption yield as a function of the tunnelling current in stationary mode for p-type samples (squares) and in scanning mode for p-type samples (circles) and n-type samples (up triangles). The solid lines are the corresponding least squares fit to a power law I a. The exponents are a = 0.3 ± 0.1, 1.3 ± 0.3, and 0.8 ± 0.3, respectively. For our scanning mode measurements we have plotted as error bars the standard deviation of our statistical ensemble for a given current. Some data points (at 3, 6, and 10 nA) are slightly shifted along the horizontal axis to improve the display. The values of the yield from previous studies as a function of the tunnelling current ([95] (down triangles) and [97] (diamonds)) and the respective least squares fit to a power law I a. The exponents are a = 15 and a = 10, respectively [101].

The major problem with the results produced by Lyding, Avouris, Stokbro and others [90,95–100] was the lack of precision. It was difficult to determine the exact number of electrons involved since the tip is scanned over the surface. The total dose is known but the electrons interact not only over the hydrogenated dimers but between them as well, so the inelastic coupling will vary from site to site. As a consequence, it seems hard to justify the mechanism given the uncertainties and especially the lack of experimental data points. Our results [101,102] show that the vibrational heating mechanism is no longer a valid description of the hydrogen extraction process. Instead, the most appropriate mechanism involves the coherent excitation of the Si–H bond as was proposed by Salam et al. [103]. In this model, the electron attaches to the r*(Si–H) orbital forming a negative ion resonance, then as the electron leaves (to the surface) it transfers a large part of its energy to the Si–H bond (several quanta) thereby climbing the vibrational ladder several levels at a time. Thus only two electrons are needed to induce hydrogen desorption. The silicon dangling bond structures which can be created by electronic-induced desorption of hydrogen atoms have many potential uses in nanolithography as was suggested in Lyding’s first paper [93]. A variety of two-dimensional patterns (Fig. 6) can be created [93,102,104,105]. Calculations of the electronic properties of the dangling bond structures suggest that they will behave as conducting wires [106], although as yet no in situ measurements have been made. The adsorption of molecules on these structures will be discussed in a later section (4.2.5). Several groups have noted that these silicon dangling bond lines have an interesting internal structure where an electronic relaxation causes only every other DB to be visible in the STM images [103,107–109]. This is an atomic-scale example of a Peierls distortion [110] or Jahn–Teller relaxation [111]. This will obviously have an effect on the transport properties of DB wires. In addition, single hydrogen atoms

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Fig. 6. An STM topography (20 · 20 nm) of the Si(1 0 0)-2 · 1:H surface showing the pattern of an OR gate formed from Si DBs created by extracting the hydrogen atoms with the STM tip. The image was taken at
1.5 V and 0.5 nA [102].

on a silicon dimer (next to a dangling bond) can diffuse [44] which means that these structures are not perfectly stable. Another use of the desorption of hydrogen from Si(1 0 0):H surfaces has been to study the isotope effect. Indeed, Avouris and co-workers found that the desorption yield for deuterium was 100 times less than that for hydrogen [92]. This was then used by Hersam et al. to study the effect of the substrate temperature during the simultaneous adsorption of hydrogen and deuterium [112]. By counting the number of desorbed hydrogen atoms, they found that the ratio of D to H varied from 5 to 50 as the substrate temperature during passivation was varied from 350 K to 650 K. 4.2.4. Desorption of hydrogen on Si(1 0 0):H with lasers As yet, there have been only a few atomic-scale studies on the manipulation of hydrogen where a laser and STM have been combined. The first study carried out by Heinz and co-workers [50] used a 532 nm 7 ns laser to induce local heating of the Si(1 0 0):H surface. The idea behind this approach is to study high energy barrier processes such as the desorption of hydrogen. The pulsed laser irradiation is capable of providing the heating rates in excess of 1011 K/s (1000 K in 10 ns) necessary for the process to be observed. The result was the production of molecular hydrogen where the H atoms came from neighbouring dimers and not the same dimer. This indicated that the thermal desorption mechanism involved hydrogen atoms on two neighbouring dimers (interdimer pathway) [113] rather than the same dimer (intradimer pathway) that had been observed and proposed by Boland [6,45] using resistive thermal heating. The problem of understanding the dynamics of the non-adiabatic processes induced by ultra-fast photon excitation and the associated excited states has been highlighted in a recent review by Kolasinski [114].

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In a different experiment, Vondrak and Zhu investigated hydrogen desorption using direct optical excitation at 157 nm [43,115]. The photon energy of 7.9 eV corresponds to the direct r–r* transition of the Si–H bond. From the detection of atomic hydrogen by time-of-flight (TOF) measurements using both p-polarisation and s-polarisation they were able to deduce that the transition dipole moment was at 18 to the surface normal which agrees very well with the calculated Si–H bond angle [116]. This led them to conclude that hydrogen desorbs via a direct optical excitation of the dipole. However, the lack of sensitivity in the TOF detection of the hydrogen ions desorbed by the laser obliged them to use a high irradiation dose of 300 J/cm2. In this regime, other indirect process could occur which might not have been detectable. We have recently carried out combined laser and STM experiments using a laser at 157 nm to induce hydrogen desorption [117,118]. We chose low irradiation doses from 1 to 23 J/cm2 so as to avoid thermal heating. This was done for three different energy pulse densities of 0.7, 2.8, and 4.1 mJ/cm2 for both n-type and p-type samples. Through a statistical analysis of the STM images obtained after irradiation, the number of new individual silicon dangling bonds produced was counted. We found that the desorption yield was three times higher for the n-doped samples than the p-type samples (Fig. 7). In addition, on the p-type surfaces, local modifications of the surface were observed in the STM images. These were ascribed to inhomogeneous laser-induced atomic-scale charging. This could be explained by the presence of B–H complexes in the sub-surface region [119] which deactivates the boron dopant in the case of p-type samples [120]. This positive charging of the surface explains the reduced photodesorption cross-section of p-type samples. Our results suggest that considering only the direct photodesorption mechanism is an oversimplified view of vacuum ultra-violet (VUV) laser photodesorption. Complex surface and sub-surface processes such as local charging should be taken into account if we are to fully understand VUV photochemistry on the Si(1 0 0):H surface. Rather than use photons to desorb hydrogen, one can use the STM to desorb hydrogen atoms and then induce the emission of photons from the tunnel junction [121,122]. Thirstrup and co-authors have used this capability to study the light emission from atomic-scale patterns on the hydrogenated surface [123]. They obtained spatial maps and spectroscopy of silicon dangling bond patterns on the 3 · 1 surface. By observing that the wavelength of the emitted photons changed as a function of the bias voltage on the tip, they proposed that the light emission involved optical transitions between a tip state and localised surface states. They found that the spatial maps were comparable to the STM images and hence deduced that the photons are emitted from a quasi-point source corresponding to the dangling bonds. In a subsequent study on the deuterated 2 · 1 surface the switching of individual silicon dangling bonds could be observed [124]. 4.2.5. Reactivity: the adsorption of molecules on Si(1 0 0):H The interaction of the hydrogenated surface with metals [4] and atoms of a large number of elements [75] has already been reviewed. In the case of metals, there is often a process of substitution in the surface between the metal and the silicon. The passivity of the surface by hydrogen does not cause any substantial restructuring of the surface especially the 2 · 1 surface. In addition, the hydrogen renders the surface chemically inert with respect to the clean surface. It can be expected that the interaction of organic molecules with the hydrogenated surface will be relatively weak – more van der Waals interactions than chemical bonding. Hydrogen also induces a wider band gap by eliminating the p

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Fig. 7. Si(1 0 0)(2 · 1):H STM topographies from n-type samples acquired at Vs =
2.5 V and It = 150 pA. (a) 40 · 40 nm2, hydrogenated surface before irradiation. (b) 40 · 40 nm2, surface after irradiation at 157 nm with a dose of 6 J/cm2 at 2.8 mJ/cm2 shot. (c) 16 · 16 nm2, illustration of hopping dangling bonds observed after irradiation noted as DB1. (d) 16 · 16 nm2, illustration of fixed sites observed after irradiation noted as DB2. NC are not counted sites. (e) Variation of the number of isolated dangling bonds by created VUV light (157 nm) on the hydrogenated Si(1 0 0) surface as a function of the irradiation dose for three different fluences and for two types of dopant [117].

surface states. This makes these surfaces useful for the study of large nano-objects such as CdSe nanocyrstals which have interesting optical properties. The adsorption of these nano-objects can be controlled on the hydrogenated surface (Fig. 8) using the pulsed valve technique [125,126]. We found that the adsorption characteristics were very dependent on

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Fig. 8. (a) STM topography (13 · 14 nm2) of a TOPO capped CdSe nanocrystal on the hydrogenated Si(1 0 0)2 · 1 surface recorded at Vs =
3.5 V and I = 0.1 nA. The inset shows the induced diffusion of the NC. The cross-section along the black arrow indicates the height and width of the nanocrystal. (b) STM topography (20 · 20 nm2) of a stearate capped CdSe nanocrystal [125].

the organic ligand used to passivate the nanocrystal core. CdSe nanocrystals with trioctylphosphine ligands (TOPO) were observed to diffuse on the hydrogenated surface under the influence of the STM tip whereas this did not happen if stearate ligands were used. This shows that the longer alkane chain of the stearate ligands (17 carbon atoms compared to eight carbon atoms for the TOPO ligands) enables the nanocrystal to be more strongly bound to the surface. From these studies we can conclude that large nano-objects such as nanocrystals can be stabilised on the Si(1 0 0):H surface because they can provide a sufficiently large van der Waals interaction energy due to the large number of ligands. The pulsed valve technique has been used to deposit other very large nano-objects on the hydrogenated surface. Poly(3-hexylthiophene) (P3HT) molecules have been adsorbed showing long polymer chains of up to 50 nm in length [127]. This group has also succeeded in adsorbing isolated multi-wall nanotubes on the Si(1 0 0):H surface and obtained STM images and spectroscopy of the nanotube [128]. To adsorb smaller organic molecules on the Si(1 0 0):H surface, we must selectively modify the surface in order to fix these smaller organic molecules. This can be efficiently done by using STM desorption to create active sites with which to position the molecules on the surface. The idea of using STM desorption of hydrogen atoms to create reactive

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sites for subsequent reaction with other atoms or molecules was first put forward by Lyding and co-workers [91,129] when the initial H desorption experiments were performed. However, individual dangling bonds are not selective in the sense that they are all identical. To select molecules by their size or orient the molecules on the surface requires the controlled construction of groups of dangling bonds as will be discussed in the next paragraph. The isolated Si dangling bonds have been used as reactive sites for molecules where one molecule is adsorbed per site. A variety of complex organic molecules have been adsorbed in this fashion, for example, Norbornadiene, copper phthalocyanine (CuPc) and C60 [130]. However, the molecules do not always attach themselves in the same way to the dangling bond. For instance, the CuPc can attach via the central metal atom or through a p interaction with one of the pyrrole groups. The Si dangling bonds have also been used as nucleation sites from which organised one-dimensional nanostructures can be grown [131]. The exposure of the Si(1 0 0):H surface to oxygen induces an oxidation reaction to propagate along a dimer row from a single dangling bond site. Another example is that of styrene molecules forming a 1-D chain via a self-propagating chain reaction [132]. The styrene molecule adsorbs on the dangling bond creating a radical on the vinyl group which picks up a hydrogen atom from a neighbouring dimer and so creates another DB site which can react with the next styrene molecule. The 1-D growth is determined by the dimer-row structure of the hydrogenated Si(1 0 0) surface. Other studies of styrene chains adsorbed on the Si(1 0 0):H surface [133] showed the presence of charge effects through the molecular chains from the initial Si DB which could be reduced by using low-doped silicon surfaces. In addition, the selective adsorption of a different molecule at one end of the chains could alter the conductivity of the chains. The controlled creation of groups of silicon dangling bonds can be used to not only select the position of molecular adsorption but can also select the molecule by its size and orient the molecule on the surface with respect to the underlying surface. These groups of dangling bonds behave as molecular molds [134]. At room temperature, the biphenyl molecule adsorbs on the clean surface across two adjacent dimers along the same row [135,136] in two possible chemisorbed configurations: bistable (pivoting about a rotation axis) and fixed. The biphenyl molecule could be trapped in either of its bistable configurations at 5 K. STM manipulation of this bistable biphenyl at 5 K showed that the molecule could be reproducibly switched from one configuration to the other and back again many times [137]. The switching yield could be determined for different positions of the excitation within the molecule and was observed to vary by a factor 200 over a distance of 400 pm. In addition, the dynamics of the molecular movement showed the presence of a transition state whose duration could be selectively adjusted by selecting different excitation positions within the molecule. At room temperature, we found that biphenyl molecules could be selectively adsorbed into the dangling bond molds on the hydrogenated surface provided that the mold was at least the size of the molecule i.e. four or more adjacent DBs on two or three neighbouring dimers (Fig. 9) [134]. The biphenyl was not observed to adsorb on 1, 2 or 3 DBs. Also the dynamics of the adsorbed molecule were modified by the presence of the hydrogen atoms surrounding the mold. The pivoting molecule in the mold was seen to fix itself spontaneously unlike on the clean surface where electrons from the tip are needed to fix the molecule in its stable site [135,136]. We tried to adsorb a larger poly-phenyl molecule called Trima (1,400 -paraterphenyldimethylacetone) in such molds and we found that the reaction was much more complicated due to the relatively more reactive molecule. The Trima molecule contains two ketone groups, one

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˚ ,
1.5 V, and 0.5 nA) of Fig. 9. A chronological sequence of four filled-state STM topographies (each 50 · 50 A the same area of the hydrogenated Si(1 0 0)-2 · 1 surface. The left-hand column shows, (a) the clean hydrogenated surface, (c) the surface after fabricating the molecular mold, (e) a moving biphenyl molecule adsorbed in the mold, and (g) the biphenyl molecule is fixed in the mold. These four stages are shown schematically in the righthand column (b), (d), (f), and (h). Note in (f), only one of the bistable states of the molecule is shown. The structure of the molecule is shown at the bottom [134].

at each end of a triphenyl chain, which are very reactive towards the clean surface. There was evidence that the Trima molecule reacted immediately with the hydrogenated surface

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before attaching to or in a mold since new DB sites appeared several tens of a˚ngstro¨ms away from the mold. The biphenyl molecule just diffuses across the surface into the mold without any apparent modification of the nearby surface. 4.3. Hydrogen on silicon carbide surfaces 4.3.1. Hydrogen on hexagonal SiC surfaces Silicon carbide surfaces are increasingly used in high temperature and high power applications because of the enhanced physical properties over silicon [9,10]. In this context, SiC surfaces have been the subject of much study over recent years (see Section 2). The reaction of hexagonal SiC surfaces (6H and to a lesser extent 4H) with hydrogen has not been so easy. The wet chemical treatment by buffered HF which works so well on Si(1 1 1) leaves OH groups on the SiC surface [138]. Similarly, in early experiments the hydrogen plasma treatment using the conditions applied to diamond produced an ordered oxide layer [139,140]. Combining the two techniques – a chemical dip in HF followed by hydrogen plasma treatment in UHV – seems to produce uncontaminated surfaces where no residual oxygen could be detected with X-ray photoemission spectroscopy (XPS) [141]. A simple hydrogenation procedure involves placing the sample at 1550 C in a hot-wall chemical vapour deposition (CVD) reactor in 1 atmosphere of hydrogen for 30 min [142]. In this study, the samples were then transferred to UHV and flashed to 950pC for p 3 min before imaging with the STM. The STM images show flat surfaces with a 3 · 3 reconstruction. Flashing the SiC surfaces in UHV before studying them in the STM means that they are not hydrogenated but the STM images suggest that hydrogen etching is efficient at the elevated temperature used thereby reducing the surface roughness of the substrate since the surfaces after treatment are much flatter than the surfaces shown in the atomic force microscope (AFM) images prior to any treatment. Effective hydrogen termination of the hexagonal surface was achieved by heating the SiC surface to 1000 C in pure hydrogen. Fourier-transform infrared adsorption spectroscopy studies (FTIR) have been carried out on hydrogenated surfaces using this procedure. They indicate the formation of only Si–H with very little SiH2 and SiH3 and no residual oxygen in the XPS, which suggests that a monohydride is formed [143,144]. Subsequent LEED studies [145] indicate that the surfaces prepared in this manner have a relatively low number of surface defects as evidenced by the clear 1 · 1 LEED pattern and low background diffraction. A further study [146] confirmed that hexagonal SiC surfaces are well hydrogenated by showing that the etching rate was minimal in the range 950–1100 C (5–50 atomic layers per hour). 4.3.2. Hydrogen on cubic SiC surfaces One of the critical physical aspects is the conductivity of these surfaces. This is amply demonstrated in the first STM studies of the hydrogenated SiC(1 0 0)-3 · 2 surface [19] where the reaction with atomic hydrogen causes the surface to become metallic. The hydrogenated SiC surface was observed using STM, ultra-violet photoemission spectroscopy (UPS) and infrared adsorption spectroscopy (IRAS). The STM images showed that the top Si–Si dimers became saturated with hydrogen as for Si(1 0 0). However, this alone could not be responsible for the metallisation. It was the combination of the three techniques that showed that hydrogen also attached to silicon atoms having a neighbouring carbon atom. Since the 3 · 2 reconstruction comprises three silicon layers on top of the

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first carbon layer [11] then this implies that the silicon atoms in the third layer have reacted. In this third layer there are very weak Si–Si dimers that are reactive to hydrogen. These are broken; however, there is only space for one hydrogen atom leaving a newly formed dangling bond on the other silicon atom. These dangling bonds interact with the continuum states to form a delocalised, and therefore, metallic band. Surface metallisation by hydrogen is a significant modification of the surface electronic properties given that the surface band gap of the clean cubic SiC(0 0 1) surface is around 2.4 eV. In addition, this is a very different behaviour from the hydrogenated silicon surface where the hydrogen saturates the surface states causing the surface band gap to increase. Several studies have shown that the different reconstructions of cubic SiC react very differently to hydrogen; the Si-rich 3 · 2 surface reacts with atomic hydrogen causing the surface to become metallic as discussed above, yet it is un-reactive to molecular hydrogen [147]. On the other hand, the SiC(0 0 1)-c(4 · 2) surface is very reactive to molecular hydrogen [147] – the sticking probability of molecular hydrogen was estimated to be several orders of magnitude higher than that on the clean Si(1 0 0) surface. In this experiment, the relaxation of the strained surface induced by hydrogenation produced a change to the 2 · 1 structure and a metallisation of the surface. Indeed, the relaxation of the clean SiC(1 0 0)-c(4 · 2) to a metallic (2 · 1) reconstruction can be temperature induced and has already been observed in STM using I(V) spectroscopy [148]. In a subsequent study, hydrogen was adsorbed on a pre-oxidised SiC(1 0 0)-3 · 2 surface [149]. Valence-band photoemission measurements of the Si 2p level and XPS measurements of the C 1s core level showed a build-up of intensity at the Fermi level accompanied by a small shift to lower binding energies in the Si 2p surface components. This indicated a hydrogen induced metallisation of the partially oxidised surface. This could be explained by a hydrogen-induced charge transfer to the top three silicon layers without much effect on the carbon layer just below. Early studies of the oxidation of the SiC(1 0 0)-3 · 2 surface using valence band photoemission had shown that the initial stages of oxidation occurred in the Si back bonds on the surface leaving the 3 · 2 reconstruction relatively intact [150]. STM studies of the initial stages of oxidation showed that some of the clean Si dimers had a high local density of states and that these sites acted as nucleation centres [151]. The oxygen reacted with these ‘‘bright’’ dimers causing neighbouring dimers to become bright in turn. Thus, the oxidation proceeded by auto-propagation forming oxidised islands on the surface. Therefore, it would seem that hydrogen-induced metallisation of cubic SiC [150] occurs at sites not involved in the oxidation. This is interesting because it means that native oxides do not have to be removed before a surface can be metallised, which would be particularly useful in device technology or at the interface with biological systems where sources of oxygen are unavoidable. 4.4. Hydrogen on diamond surfaces 4.4.1. Introduction: properties of diamond surfaces Diamond surfaces are receiving more attention nowadays in surface science because of their potential applications in nano- and bio-technology. This is due to the exceptional physical properties of diamond. It has a large gap (5.5 eV) which gives pure samples a transparency at optical wavelengths. Diamond has a high mobility of both electrons and holes, and is thermally very stable which is attractive to microelectronic applications. Diamond surfaces are also bio-compatible. Hydrogenated diamond surfaces share these

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properties; the surface is thermally very stable since hydrogen desorbs only above 900 C (compared to 400 C for silicon). Hydrogenated surfaces are very stable in air for many weeks. Hydrogenated diamond has one very special property; the surface manifests a negative electron affinity (NEA) [152,153] because the conduction band is higher in energy than the vacuum level. This is because the C–H bonds terminating the diamond surface create a surface dipole which causes a significant upward band bending of the valence and conduction band, pushing the conduction band minimum above the vacuum level. The consequence of this is that when irradiated with light, electrons are spontaneously emitted. Thus, hydrogenated diamond has been used in applications as electron-emitters or cold-cathodes [154,155] as had been predicted earlier [156,157]. Another important aspect motivating recent studies is that clean diamond and hydrogenated diamond have very different electrical properties. Clean diamond is an almost perfect insulator whereas hydrogenated diamond behaves as a reasonably conducting p-type semiconductor [158,159]. STM observations of both the hydrogenated and clean surface will be discussed in Section 4.4.2 and the conductivity in the Section 4.4.3. As described in Section 2, hydrogenated diamond surfaces can be prepared in a hydrogen plasma or by homo-epitaxial chemical vapour deposition (CVD) with a hydrogen/ methane mixture. We will describe the atomic-scale studies on single crystal diamond surfaces prepared by plasma treatment in the next paragraph. The highest quality samples are obtained by plasma treatment of single mono-crystals because CVD growth introduces grain boundaries and point defects which cause surface roughness and are detrimental to the optical properties of diamond. Nevertheless, if homo-epitaxial growth is performed on undoped single crystals surfaces then usually the surface is of good quality. A review article by Kawarada [160] gives a detailed survey of hydrogenated diamond surfaces. His review places more emphasis on the surface preparation methods and conditions, in particular, using chemical vapour deposition (CVD) processes. The surface morphology of diamond surfaces is presented as observed by reflection high energy electron diffraction (RHEED), and high resolution electron energy less spectroscopy (HREELS). The electron transport properties of hydrogenated diamond surfaces are discussed within the context of metal–diamond contacts. The few STM experiments presented in Kawarada’s review article were carried out in air [161–163] and on polycrystalline films [164,165]. 4.4.2. Atomic-scale imaging of diamond surfaces In a recently published review of STM studies of hydrogenated diamond surfaces [166], we presented the case for studying single crystal diamond surfaces in UHV, by virtue of their transparency for future optical applications and the low surface roughness and lack of defects for STM studies. In recent work, we have shown that one can produce high quality surfaces using the hydrogen plasma method described in Section 2 on single crystal diamond. Indeed, our STM images (Fig. 10) show atomically resolved hydrogenated surfaces [37]. However, even though it is possible to obtain high quality atomic-scale images of diamond surfaces produced in a CVD reactor [167–169], it is not clear if the diamond samples retain the optical properties of single crystals. As mentioned in the introductory section above, while hydrogenated surfaces have conducting p-type semiconductor characteristics (see next section), the clean diamond surface is an insulator. This is a major obstacle if one wishes to observe the atomic structure of the surface. We were able to obtain the first atomically resolved STM topographies of the clean diamond surface (Fig. 11) by using a novel imaging method [170]. This was achieved

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Fig. 10. STM topographies of the hydrogenated diamond C(1 0 0)-(2 · 1):H surface: (a) Ubias =
1.5 V, It = 1.5 nA (unoccupied states), and (b) Ubias = +1.5 V, It = 1.0 nA (occupied states). The bright lines on the top topography indicate the C–C dimer rows in the vicinity of the step (SA) [37].

by injected electrons from the STM into the conduction band of the diamond. A high bias of 6 V had to be applied. Under these conditions, there is a small vacuum gap above the surface through which the electrons pass and standing-wave resonances are formed in the potential well in this vacuum gap just above the surface as the electrons are reflected by the surface (Fig. 12). One could observe these resonances using Z(V) spectroscopy. Barrier resonances and surface reflectivity have been considered from a theoretical point of view as a function of the surface potential and tip structure [171]. The Z(V) spectra

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Fig. 11. A STM topography of the clean diamond C(1 0 0)-(2 · 1) surface recorded by resonant injection of the electrons into the conduction band (6 · 6 nm, Ubias = +5.9 V, It = 1.0 nA) [170].

Fig. 12. Potential well for the electrons scattered on the diamond surface. The potential barrier in the vacuum ˚ tip–sample separation and a 0.4 V/A electrostatic field in the vacuum gap. gap was constructed assuming a 10 A The actual shape of the potential barrier was obtained by a superposition of the electric field and the image potential of the surface, modified to include the surface potential contribution. For both the tip and the sample, the positions of the vacuum level, Evac, and the Fermi level, EF, are indicated. In addition, the surface gap Eg, the positions of the valence band maximum, VBM, and the conduction band minimum, CBM, are labelled. In the potential well, the reflection coefficients, rB, rC, and the dephasing, uB, uC, of the barrier and the diamond surface respectively, are noted at the first resonance, n = 1 [170].

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of the clean diamond surface showed a series of peaks between 5 and 10 V. However, atomic resolution could only be obtained at a certain bias which corresponded to the first resonance. This is because the first resonance energy varies as a function of the position of the tip above the surface – it changes by an estimated 11 meV between a position on top of the dimer row and between the dimer rows. Most importantly, however, the first resonance is exceptionally intense and narrow in the Z(V) spectrum. This can be explained ˚ /V, it is an order of magby the unusually large dZ/dV slope at the first resonance. At 20 A nitude larger than in other materials [172–174]. Fundamentally, the first resonance has a very narrow width (0.15 eV). Such fineness is expected when the reflectivity of the electron wave function on the surface is high [172]. This is the case here, for two reasons. First, the band gap of diamond is large, which places the bottom of the conduction band very close to the energy of the resonance (see Fig. 12). The large change in the k vector when the electrons penetrate into the diamond at the bottom of the conduction band results in this high surface reflectivity. Second, the energy of the first resonance is only slightly above the energy of the conduction-band minimum (see Fig. 12). It follows that the off-resonance tunnel current should be very small compared to the on-resonance current, resulting in a large Dz step at the resonance. It should be emphasised that both effects, the narrow resonance width and the large Dz step, apply only to the resonance closest to the bottom of the conduction band. As a consequence, the fineness of the lowest resonance is much higher than that of the higher lying resonances. This resonant electron injection technique was then applied to the hydrogenated diamond surface and the clean silicon surface [175]. There we showed that both hydrogenated diamond and clean silicon have these standing-wave resonances in the vacuum gap when the bias applied is above the vacuum level of the surface. Atomically resolved images could be obtained on both the hydrogenated diamond and clean silicon surfaces. However, the corrugation was far less due to the fact that the first resonances for these two surfaces were much less intense and broader. This confirmed that the exceptional nature of the first resonance on the clean diamond surface is responsible for the atomic resolution. 4.4.3. Conductivity of hydrogenated diamond surfaces The conductivity and transport properties of hydrogenated diamond surfaces have been the subject of much discussion. One of the difficulties has been to understand the nature and origin of the conductivity of diamond surfaces. Surface conductivity (SC) is due to the presence of several factors: the dopant atoms present in the bulk and sub-surface regions, defects, adsorbates and the existence of surface states. In the literature, over the last decade or more, the various factors just mentioned have been put forward. Boron is most commonly used to provide intrinsic p-type conductivity in bulk diamond [176] and recently, n-type doping with phosphorus has been reported [177]. However, it has been proposed that surface hydrogen plays the role of the dopant by being directly responsible for the hole accumulation layer by forming shallow electron acceptors [178]. Furthermore, it has been proposed that the acceptors form a buried layer 30 nm below the surface [179]. An experimental study carried out around the same time by Maier et al. [180] provided evidence suggesting that chemisorbed hydrogen is a necessary but not sufficient prerequisite for surface conductivity and that the electron acceptors are provided by atmospheric adsorbates. They proposed a mechanism in which a redox reaction in an adsorbed water layer provides an electron sink for the sub-surface hole accumulation. This model has been supported in subsequent studies [181,182]. Underlying all these factors is the type of

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sample and the preparation method used. By nature, polycrystalline surfaces contain a large proportion of defects, both structural and elemental. Also, both hydrogen plasmas and CVD growth introduce hydrogen into the sample. Other studies have supported the variation of surface conductivity with the sample preparation (annealing, exposition to air, and other factors) [183]. In general, the investigations mentioned above have measured the surface conductivity using the point contact method. I(V) spectroscopy using the STM provides a measure of the surface band gap and the local density of states (LDOS) near the valence band and conduction band edges. However, one cannot necessarily draw the same conclusions from the surface conductivity and STM spectroscopy measurements as they are two intrinsically different methods; the first is macroscopic and the second nanoscopic. Both boron and hydrogen atoms participate as doping atoms in the conductivity of diamond. Both act as p-type dopants; however, the boron doping level is about 0.4 eV above the Fermi level, which makes the surface conductivity more sensitive to temperature variations. In a recent study, we investigated the conductivity of both a type IIa (undoped) and a type IIb (boron-doped) diamond sample [38]. Using STM and I(V) spectroscopy, we found that several elements affected surface conductivity: adsorbed species such as water or CO2 [180], sub-surface hydrogen, surface hydrogen and bulk dopant atoms. We measured I(V) spectroscopy curves with the STM as a function of thermal treatment (no heating, heating to 400 C, 750 C and 950 C). For the undoped IIa sample, STM images of the sample could be obtained if it was not heated. After heating the undoped IIa sample to 400 C no image could be obtained, but it was possible to restore the sample conductivity by exposing it to air again. For the doped IIb sample, filled-state STM images could be obtained under normal tunnelling conditions (
1.5 V and 0.5 nA) at all temperatures but not after heating to 950 C. At 950 C the diamond surface is clean and images can only be obtained using the resonant injection method [170] discussed in the previous section. The I(V) curves showed a tendency for the gap to enlarge, suggesting that the surface species are desorbed during the thermal treatment. These results suggest that the level of boron doping actually plays a relatively important role in the surface conductivity as observed in the I(V) spectroscopy curves. It is clear from this discussion that the basic mechanisms of hydrogenated diamond surface conductivity are still far from being completely understood. 4.4.4. Desorption of hydrogen from hydrogenated diamond surfaces The negative electron affinity of the hydrogenated diamond surface [152,153] has been applied to make high efficiency electron emitters [154,155]. Thus, the manipulation of hydrogen on hydrogenated diamond surfaces using the STM could provide a way of fabricating atomic-scale electron emitters. With this in mind we attempted to desorb individual hydrogen atoms from the hydrogenated diamond surface using the STM [184]. Using a similar procedure to that used on silicon, we applied a series of one hundred 2.5 V voltage pulses, each 2 ms in duration and regularly spaced along a line 20 nm long (Fig. 13). This is not a continuous scan line at constant current rather a series of stationary mode pulses with a short pause (20 ms) between each to allow the tip to displace horizontally and return to the same tip–surface distance before the next pulse. We chose three different pairs of voltage (V) and current (I): 2.7 V and 50 nA, 5.5 V and 1 nA, 7 V and 5 nA. The lowest voltage corresponds to the tunnelling regime (voltage less than the work function), whereas the higher voltages correspond to the field emission regime (voltages above the

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Fig. 13. Desorption of hydrogen from the hydrogenated diamond C(1 0 0)-(2 · 1):H surface. The corresponding STM topographies (15 · 19 nm) (a) before and (b) after the desorption procedure were recorded at Ubias =
1:5 V, It = 1.0 nA. The dashed line in (b) indicates the line traced by the tip across the surface during the manipulation. The bright features in (b), scattered around the dashed line and highlighted by the white circles, represent the dangling bonds after desorption of individual hydrogen atoms [184].

work function of the surface). By counting the number of bright dangling bond sites created during the pulse and dividing by the total number of electrons (I · 200 ms), one can estimate the desorption yield. We found yields of 2 · 10
11, 7 · 10
8, and 2.7 · 10
8 H-atoms per electron. On silicon, the desorption yields are 2 · 10
9, 7 · 10
8 and 2.4 · 10
6 H-atoms per electron for the same applied voltages [95,99] (though not the same current). Immediately, one can see that the desorption process on diamond at 2.7 V is three orders of magnitude lower than in the field emission regime. Indeed, in the field emission regime, hydrogen atoms could be removed easily however, there was no localisation of the incident electron beam from the STM tip and hydrogen atoms were desorbed from a zone over 10 nm wide centred on the line of voltage pulses applied to the tip. At 2.7 V, the electrons emitted from the STM tip cannot reach the conduction band of diamond because of the large band gap; the electrons probably pass into the sub-surface region which is conducting and therefore relinquish their energy in breaking the hydrogen bond. However, the inelastic scattering of the electrons is very different in the case of diamond when compared to silicon (or germanium). In silicon, the electrons pass directly into the conduction band. A quantitative comparison between the three surfaces (diamond, silicon, and germanium) is rather difficult to make for several reasons. First, the bond energies are not the same, C–H = 4.2 eV, Si–H = 3.5 eV and Ge–H = 3.0 eV. Second, it is not always easy to position the energy of antibonding r* orbital in each case and hence, thirdly to know the energy of the r–r* transition. Another way to desorb the hydrogen from the hydrogenated diamond surface is to use photons to desorb the hydrogen atoms from the diamond surface. In fact, in a similar fashion to the experiments on the silicon surface described in Section 4.2.2, partially hydrogenated surfaces can be prepared in several ways: by photon irradiation, thermal heating and the adsorption of hydrogen on the clean diamond surface using a hot tungsten filament. These surfaces have been studied using synchrotron radiation. Some have looked at the electronic structure (see the next section) while others have looked at the desorption mechanism. It was possible to detect for the first time the C(1s)-r*(C–H) resonance at

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287.5 eV which was enhanced in the H+ yield [185]. By comparing the PSD H+ yield with the NEXAFS spectra of low energy electrons (8 eV and 35 eV), a sharp peak was visible at 287.5 eV in the PSD spectrum. This resonance was hardly visible in the NEXAFS spectrum using 35 eV electrons and not at all in the spectrum using 8 eV electrons – the onset was seen at 289 eV. Eight eV electrons probe the volume whereas 35 eV electrons are more ˚ which corresensitive to the surface; nevertheless their mean free path is around 10 A + sponds to several atomic layers. These results show that the H ions are the most sensitive to the surface. This provided evidence for two distinct processes resulting in H+ desorption, a direct electronic excitation of C–H (the peak at 287.5 eV) and an indirect excitation of C–H by secondary electrons in the bulk (above 289 eV) [186]. In addition, H
desorption is observed to be an indirect process involving secondary electrons and is not observed on amorphous diamond films [186]. Ion-implantation can be used to introduce a controlled level of defects in the near-surface region. For ion-damaged diamond films produced in this way, the dominant desorption mechanism is modified with respect to un-damaged films; photon stimulated H desorption via the direct C(1s)-r* excitation of C–H is more efficient than for undamaged films [187]. 4.4.5. Electronic structure of hydrogenated diamond surfaces The electronic structure of the hydrogenated diamond surface has also been examined using SR. Photoemission studies of the valence band give access to the bonding orbital energies just below the Fermi level whereas NEXAFS of core levels excites electrons into the antibonding orbitals. Therefore, information from both experiments is required to obtain a more complete picture of the electronic structure of species adsorbed on surfaces. Early studies using valence-band photoemission on the fully hydrogenated diamond surface showed evidence for bulk states [188,189] at around 8 eV below the Fermi level, while theoretical calculations suggested that no surface states should exist in the band gap of the hydrogenated surface [190,191]. Only recently has a photoemission study of the valence band shown the existence of surface states corresponding to C–H [192]. This was achieved by recording the photoemission spectra at grazing angles to improve the sensitivity to the surface. The spectra obtained on the plasma-hydrogenated surface were compared with those of the clean surface and the in situ filament hydrogenated surface. C–H surface states were observed at 8.8, 11.4, and 16.9 eV which disappeared when the hydrogen was removed and re-appeared when the clean surface was hydrogenated in situ with a hot filament. It was possible to distinguish between the bulk bands and the surface states because the bulk bands show a strong dispersion as a function of photon energy whereas the surface bands did not. The partially hydrogenated diamond surface can be produced either by photodesorption or by thermal annealing or by hydrogenation in vacuum of the clean diamond surface. The electronic structure of this surface has been studied using SR for each of the preparation methods [193]. In the NEXAFS spectra the surface resonances were investigated, that is resonances appearing in the band gap below the bulk core excitation threshold at 289.3 eV (Fig. 14). For the hydrogenated surface, a peak is observed at 287.2 eV and as the surface is irradiated progressively with photons, this peak diminishes and shifts to slightly higher energy. As the irradiation dose increases, first a new surface state appears at 282.5 eV (Ex2), then a second state at 283.8 eV (Ex1) for higher doses followed by a weak peak at 286 eV (Ex3). When the clean surface was exposed to atomic hydrogen (from a hot tungsten filament) the intensity of each peak followed the reverse behaviour as the

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Fig. 14. In (a) C 1s photoemission spectra and (b) NEXAFS spectra of the fully hydrogenated diamond C(1 0 0)(2 · 1):H surface as a function of the synchrotron irradiation dose: curve A, no irradiation was done; curve B, 51 lA min, curve C, 210 lA min; and curve D, 344 lA min. (a) The open circles and dotted lines represent the original C 1s spectra and the background of secondary electrons, respectively. The thin solid lines represent the de-convoluted C 1s components (see Appendix A). The thick solid lines (almost coincident with the original spectrum) represent the superposition of the de-convoluted C 1s components. The shadow areas represent the SC component. For each irradiation dose the hydrogen coverage (hH) was calculated the equation described in the article. (b) The diamond bulk exciton is shown in the insert [192].

amount of hydrogen was increased. By considering the binding energy of the C(1s) hole, it was possible to position the different peaks with respect to the Fermi level and hence assign the peaks. The surface state appearing at 282.5 eV (Ex2) was assigned to single carbon dangling bonds on the surface, and the states at 283.8 eV (Ex1) and 286 eV (Ex3) to clean p-bonded dimers. However, both Ex1 and Ex2 are considered to be core-exciton states where the electron and hole remain paired whereas the peak at 286 eV (Ex3) was assigned to an electronic transition into a p* state within the band gap as had been suggested by Graupner et al. [194]. An analysis of the peak intensity versus hydrogen coverage as a

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function of the method used to produce the partial hydrogenated surface showed that there was no preferential pairing of hydrogen atoms on the C–C dimers at room temperature but that thermal annealing did induce pairing. 4.4.6. Reactivity of hydrogenated diamond surfaces There have been a number of studies on Schottky contacts formed by the deposition of thin metal films on diamond surfaces. This has been reviewed earlier by Kawarada [160]. In general, the Schottky barrier heights are in the range 1–2.5 eV depending on the metal with no correlation with the metal work function. It appears that the best contacts are formed on the hydrogenated (1 0 0) surface as adsorbates on the (1 1 1) surface reduce the quality of the metal contacts. As yet there have been very few studies on the adsorption of molecules on the hydrogenated surface. FTIR experiments have studied the 2 + 2 Diels– Alder cyclo-addition reaction between cyclo-pentene and the clean diamond surface [195]. Synchrotron radiation has been used to study the adsorption of molecular oxygen on the partially hydrogenated diamond surface [196]. It was found that oxygen does not react with either the clean diamond surface or the fully hydrogenated surface. The clean surface is unreactive because of the stability of the p-bond of the C–C dimers [197]. However, oxygen was

Fig. 15. Valence band photoemission spectra of the C(1 0 0)-2 · 1 surfaces recorded at a photon energy of 50 eV. (a) The fully hydrogenated surface (thick curve) followed by the in situ irradiation (thin curve). The irradiation dose is 26 lA min. (b) The difference between the two curves shown in Fig. 15(a). (c) The in situ irradiated surface (thin curve with an irradiation dose of 123 lA min) followed by the in situ oxygen adsorption (thick curve). Oxygen exposure is 25 L. (d) The difference between the two curves shown in Fig. 15(c) [196].

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observed to react with the partially hydrogenated surface that had been prepared either by photon stimulated desorption or by in situ deposition of atomic hydrogen (Fig. 15). This means that oxygen reacts with the single dangling bonds on the partially hydrogenated surface. Since thermal desorption produces clean carbon dimers, photon desorption must be used to produce the single dangling bonds necessary to react with the oxygen. 5. Hydrogen on other semiconductor surfaces In this section we will give a brief overview of studies on other semiconductor surfaces, where these surfaces have been only rarely studied from the point of view of the manipulation of hydrogen, such as GaAs or Ge(1 0 0). The exception is the Si(1 1 1) surface which has been widely studied in terms of understanding the structure of the hydrogenated surface but very little has been done towards manipulating hydrogen on this surface. 5.1. GaAs(1 0 0) The hydrogenated GaAs surface has been studied because of the importance of chemisorption processes and the rearrangement of surface atoms. A variety of surface sensitive techniques have been used to investigate the interaction of atomic hydrogen with GaAs surfaces. High-resolution electron energy loss spectroscopy [198] has shown that H adsorbs on both the Ga and As sites on a GaAs(1 1 0) surface. STM studies on the GaAs(1 0 0) surface showed that one or two hydrogen atoms adsorbed on the As sites and at high coverage AsH3 formed which then desorbed breaking up the As dimer structure [199]. ESD has been used to desorb both H+ and H
from GaAs surfaces suggesting that super-excited molecular complexes containing a core hole can lead directly to the ejection of negative ions [200]. Valence-band excitation has induced H+ desorption from a H2O layer on GaAs(1 1 0) surface suggesting that a core level to conduction band excitation followed by Auger decay was responsible [201]. PSD of H+ from GaAs around the core-level energies of Ga and As atoms has provided evidence of a direct desorption process [202]. The importance of hydrogen in the preparation of flat well-reconstructed GaAs surfaces has been highlighted [203]. Significant reductions in surface roughness and defect density can be achieved. The quality of GaAs surfaces after hydrogen cleaning has been studied with STM [204]. STM studies of the adsorption of hydrogen on the As-rich GaAs(0 0 1)-c(4 · 4) surface [205] show that at low temperatures (around 50 C) hydrogen produces a disordered (1 · 1) surface covered with AsH2 and AsH3 clusters. At higher temperatures (150–400 C) there is formation of mixed c(2 · 2) and c(4 · 2) surface domains with H adsorbed on surface Ga atoms. This confirms earlier calculations [206]. The Ga is exposed because H induces the loss of As from the surface. At the highest temperature (480 C) a disordered (2 · 4) reconstruction was formed due to thermal desorption of As from the surface. The results are consistent with the loss of As from the surface, either through direct thermal desorption or as a result of the desorption of volatile compounds which form after reaction with H. 5.2. Ge(1 0 0)-2 · 1 STM has been used to study the hydrogen passivated Ge(1 0 0) surface in an ambient nitrogen atmosphere [207]. The Ge surface had been passivated in HF solution. However,

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only large micron sized images with no atomic resolution were presented. Micron-sized structures were produced by electron stimulated desorption using the STM tip. These structures were then oxidised. The vibrational excitation mechanisms on hydrogenated Ge(1 0 0) surfaces has been studied with HREELS [86]. Similar to the Si(1 0 0):H surface it was found that dipole scattering dominates the excitation of the bending and scissor modes and that the stretching mode is enhanced between 2 and 6 eV. A maximum at 2.5 eV was assigned to a negative ion r resonance. 5.3. Si(1 1 1)-7 · 7 5.3.1. Introduction: the adsorption of molecules on the clean Si(1 1 1)-7 · 7 surface A short discussion of the Si(1 1 1)-7 · 7 surface will be given in this section. There are several reasons why this has not been included in the main section of the paper. The first is that both the clean and hydrogenated Si(1 1 1) surfaces have been the subject of much study over the last few decades and therefore have been reviewed before [3,4,52,75]. Secondly, a large number of studies have been performed on the STM manipulation of atoms and molecules on the clean surface. However, contrary to the hydrogenated Si(1 0 0) surface, the manipulation of atoms or molecules has not been investigated on the hydrogenated Si(1 1 1) surface. We will take the opportunity to illustrate the variety of studies that have been made on the Si(1 1 1)-7 · 7 surface using the studies by our research group as a starting point. Silicon adatoms have been desorbed from the clean Si(1 1 1) surface using both electronic excitation [208–213] and direct contact between the tip and the surface [214]. The direct contact mode was first used to extract individual Ge atoms from the clean Ge(1 1 1)c(2 · 8) surface [215]. A large number of small molecules have been studied on the 7 · 7 surface as the review in 1995 shows [75]. Since then further studies have been carried out. As an example, NO has been manipulated using STM [216] and valence-band and core-level photoemission with selective ion photodesorption [217]. The adsorption and manipulation of oxygen has received even more attention due to its importance in microelectronic device technology. Again both valence-band and core-level photoemission [218– 222] and STM [60,223–229] have been used. A variety of larger organic molecules have been studied using synchrotron radiation by our group: benzene [230–232], formic acid [233], methanol [234–236], phenol [237] and toluene [238]. The adsorption of benzene has been studied by STM [239,240] as has geranyl acetone using both STM and valence-band photoemission [241]. One of the reasons why the Si(1 1 1)-7 · 7 surface is not ideal for studying molecular manipulation is that the interpretation of the manipulation results is problematic because one cannot see any molecular structure in the STM images, only a change in the intensity of the adatoms or restatoms is observed. It can be done if the manipulation of a molecule that produces easily distinguishable reaction products as in the case of chlorobenzene [242–244]. 5.3.2. Desorption of hydrogen on the Si(1 1 1):H surface The preparation of the hydrogenated Si(1 1 1) surface has been discussed in Sections 2.1.2 and 2.2, based on the work by Chabal and co-workers [29–31]. While the structure of the fully hydrogenated Si(1 1 1) surface has been investigated with synchrotron radiation [32,245–248], STM [3,27,249], and theoretical calculations of the electronic structure

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[248,250], there have been few studies of the manipulation of hydrogen on the fully hydrogenated Si(1 1 1) surface. Manipulation of hydrogen has been achieved using photon stimulated desorption with synchrotron radiation [251]. Here, it was found that H+ ion desorption could be produced by both valence band and Si(2p) core level excitation. In both cases, this could be explained by a multi-electron process. Electron stimulated desorption using the STM produced an interesting restructuring of the surface when hydrogen was desorbed from a large area [33]. The clean area of the surface transformed into a 2 · 1 structure rather than the expected 7 · 7 structure which suggested that the 2 · 1 surface is thermodynamically more stable. First-principles calculations suggest that a row of dangling bonds on the (1 1 1) surface is stable [252], however, this cannot be directly compared with the STM observations as they did not calculate a p-bond structure. Early manipulation studies by Wintterlin and Avouris [63], confirmed that the basic desorption mechanism involved electronic excitation and was not a field effect. Recently, hydrogen atoms have been desorbed individually producing small groups of silicon dangling bonds [214] using the direct contact method [215]. 6. Summary and perspectives 6.1. Preparation of hydrogenated semiconductor surfaces In this review article, we have described the preparation procedures for the various semiconductor surfaces. It is clear that each surface has a different recipe that works best, for example, monohydride formation on the Si(1 0 0) surface is best done with a hot filament to crack the hydrogen whereas plasma hydrogenation of the diamond surface is the only option. However, there are several fundamental questions to be resolved. First of all, there are as yet no experimental atomic-scale STM studies on the hydrogenated hexagonal SiC surface. These will be important given the increasing interest in studying wide-band gap surfaces. Secondly, hydrogenation of the semiconducting surfaces leads to the presence of hydrogen in the sub-surface layers of the substrate. The role of this hydrogen is not well understood and has direct consequences on the electrical properties of the substrates. While the manipulation of atoms and molecules on fully hydrogenated surfaces is a rich field of investigation, it is quite apparent that partially hydrogenated surfaces are of great interest as the surface properties are modified with respect to both the hydrogenated and the clean surfaces. Moreover, these modifications can be done at the atomic scale with both the STM, lasers and synchrotron radiation. 6.2. Electronic structure and conductivity Understanding the electronic structure and conductivity of hydrogenated surfaces remains a fundamental requirement if one is to optimise the use of hydrogenated surfaces. Several fundamental physical aspects need still to be determined more clearly especially as they are interrelated. Firstly, the surface band gap needs to be characterised. As semiconductor surfaces have a different atomic structure with respect to the volume, the surface band gap is often different from that of the volume. In particular, we need to know the surface states and band edge positions relative to the Fermi level. Here, the band bending is crucially important especially as it can be modified at the atomic scale by the electric field due the presence of the STM tip, for example.

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A second important aspect concerns the unoccupied electronic states. This applies both to the surface and any adsorbed species on the surface. If we wish to activate a molecular function (for example, a movement or photon emission), it is easiest at the moment to inject electrons into the unoccupied states forming a negative ion resonance. Whether these are the unoccupied states of the molecule or the surface, in both cases, we need to know their energy with respect to the Fermi level and in the case of the molecule, its electronic and conformational structure. The information provided by STM spectroscopy and synchrotron radiation is difficult to interpret. In the first case because the tip states contribute to the spectrum and in the second, inverse photoemission is not easy and core-level excitation creates an electron–hole pair which shifts the energies from their absolute values. The surface conductivity is the third aspect of importance. Again, the surface band gap and band bending play a crucial role as do the presence and positions of surface states. In addition, the effect of any sub-surface species such as hydrogen on the surface conductivity is poorly understood and still a subject of much debate. Likewise, doping atoms (which are necessary for the semiconductors to have bulk conductivity) also influence the surface conductivity and the band structure since the dopant states lie just above the valence band maximum (p-type doping) or just below the conduction band minimum (n-type doping). 6.3. STM and laser induced desorption of hydrogen The major conclusion that should be underlined is that the processes of hydrogen desorption are still poorly understood both for electronic excitation with the STM and laser-induced desorption with photons. Several aspects should be highlighted. First, the interaction between the STM tip and the hydrogen on the surface is very localised especially at low voltages in the tunnel regime as was shown for the desorption of isolated hydrogen atoms on the Ge(1 1 1) surface. This renders the manipulation procedure relatively inefficient. Other complications arise in the case of a fully hydrogenated surface such as Si(1 0 0):H where interactions between neighbouring hydrogen atoms can influence the desorption process. Secondly, it is not always easy to identify the electronic state involved in the excitation for example; a voltage threshold might be observed that is difficult to attribute to a particular electronic state. Thirdly, the desorption process is very tip dependent. In fact this is probably the biggest unknown since ideally we would like to know the structure of the tip which we do not. In the laser experiments where the hydrogenated surface is irradiated with photons, one problem is that the direct electronic excitation process is not always dominant. Other secondary process involving the substrate can occur, for example on the p-doped Si(1 0 0):H local surface charging at the atomic scale is observed. In both cases, further studies and detailed calculations will be necessary to advance our understanding. The problem of understanding the dynamics of the non-adiabatic processes induced by ultra-fast photon excitation and the associated excited states has been the subject of a recent review. 6.4. Reactivity of hydrogenated surfaces Hydrogenated semiconductor surfaces are ideal for adsorbing molecules or other nanoobjects. Depending what experiment is to be carried out, one can carefully choose the

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appropriate molecule on an appropriate surface. For example, fully hydrogenated surfaces tend to create weaker bonds with molecules due to the predominant van der Waals forces over the chemical bonds that from on a clean surface. Hence diffusion processes may be studied. Partially hydrogenated surfaces are potentially very useful because individual sites of molecular molds can be created at chosen locations on the surface and therefore the adsorption site and eventually the orientation of the molecule can be selected. Other properties are equally important, for example, hydrogenated diamond samples are stable in air (hydrogenated SiC is also probably stable). This makes these surfaces suitable for adsorbing biological molecules or depositing molecules from solution. In addition, other investigation techniques such as atomic force microscopy and scanning electron microscopy can be used on these surfaces. Furthermore, the properties of the surface have to be taken into account in combination with the chosen deposition technique which in turn depends on the properties of the molecule (leak valve, filament, crucible, pulsed valve). We have shown that it is possible to deposit very large nano-objects (CdSe nanocrystals) using the pulsed valve technique in a controlled manner without polluting the surface so that the atomic structure is clearly visible in the STM images. 6.5. Future work There are two principle aspects that we would like to put forward for future work. 6.5.1. Atomic-scale engineering With the STM and now combined laser-STM techniques, further progress can be made in engineering at the atomic scale the surface properties by desorption. Two examples come to mind, the first is that the work function can be modified at the atomic scale (hydrogen on Ge(1 1 1)) and second, on a fully hydrogenated surface, we can engineer molecule molds where molecules can be selectively adsorbed as a function of their size and chemical properties. It is quite likely that the molecular properties will be modified to some degree by the mold, for example, rotation may become easier or made more difficult. 6.5.2. Hydrogenated surfaces as substrates for molecular nanomachines The concept of molecular nanomachines has only really appeared in the last few years as expertise in both local probe techniques and chemical synthesis of molecules has advanced. We can imagine many uses of hydrogenated surfaces towards this end. Conducting atomic lines can be created by desorption of the hydrogen atoms. Hydrogenated surfaces could be used for studying molecules in ambient conditions provided that both the molecule and the surface are stable. Hydrogenated surfaces are relatively chemically inert with respect to their clean analogues and so can be used to deposit large molecules using the pulsed valve technique, for example. Due to the weaker molecule surface interactions, diffusion should be easier than on the clean surfaces and so could be used to construct two-dimensional supra-molecular structures as has been done at low temperatures on metals. Finally, most hydrogenated surfaces have a larger band gap than their clean surface counterparts which can be used in optical experiments or indeed for reducing the electronic interaction of molecular states with the substrate. Two exceptions to this are diamond, which has a larger gap for the clean surface, and cubic silicon carbide which becomes metallic.

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Acknowledgements We would like to thank our former Ph.D. students who have contributed to the success of this research work: Franck Rose, Laetitia Soukiassian, and Mathieu Lastapis. Our present Ph.D. students: Romain Bernard and Marion Cranney, and post-doctoral students: Marilena Carbone, Kirill Bobrov, Alex Laikhtman. and Marta Martin. This work is also the result of fruitful collaborations: Patrick Soukiassian, Alon Hoffman, and Philip Bergonzo. Financial support has been provided by the CNRS and the European Union Network ‘‘Atomic and Molecular Manipulation: a new tool for Science and Technology’’ (AMMIST). References [1] K. Christmann, Interaction of hydrogen with solid surfaces, Surf. Sci. Rep. 9 (1988) 1. [2] J.A. Schaefer, Electronic and structural properties of hydrogen on semiconductor surfaces, Physica B 170 (1991) 45. [3] J.J Boland, Scanning tunnelling microscopy studies of the interaction of hydrogen with silicon surfaces, Adv. Phys. 42 (1993) 129. [4] K. Oura, V.G. Lifshits, A.A. Saranin, A.V. Zotov, M. Katayama, Hydrogen interaction with clean and modified silicon surfaces, Surf. Sci. Rep. 35 (1999) 1. [5] J.J Boland, Structure of the H-saturated Si(1 0 0) surface, Phys. Rev. Lett. 65 (1990) 3325. [6] J.J Boland, Evidence of pairing and its role in the recombinative desorption of hydrogen from the Si(1 0 0)2 · 1 surface, Phys. Rev. Lett. 67 (1991) 1539. [7] J.J Boland, Role of bond strain in the chemistry of hydrogen on the Si(1 0 0) surface, Surf. Sci. 261 (1992) 17. [8] D.T. Jiang, G.W. Anderson, K. Griffiths, T.K. Sham, P.R. Norton, Adsorption of atomic hydrogen on Si(1 0 0)-2 · 1 at 400 K, Phys. Rev. B 48 (1993) 4952. [9] V.M. Bermudez, Structure and properties of cubic silicon carbide (1 0 0) surfaces: a review, Phys. Status Solidi B 202 (1997) 447. [10] P. Soukiassian, Cubic silicon carbide surface reconstructions and Si(C) nanostructures at the atomic scale, Mater. Sci. Eng. B 96 (2002) 115. [11] F. Semond, P. Soukiassian, A.J. Mayne, G. Dujardin, L. Douillard, C. Jaussaud, Atomic structure of the b-SiC(1 0 0)-(3 · 2) surface, Phys. Rev. Lett. 77 (1996) 2013. [12] P. Soukiassian, F. Semond, L. Douillard, A.J. Mayne, G. Dujardin, L. Pizzagilli, C. Joachim, Direct observation of a b-SiC(1 0 0)-c(4 · 2) surface reconstruction, Phys. Rev. Lett. 78 (1997) 907. [13] P. Soukiassian, F. Semond, A.J. Mayne, G. Dujardin, Highly stable Si atomic line formation on the b-SiC(1 0 0) surface, Phys. Rev. Lett. 79 (1997) 2498. [14] L. Pizzagalli, C. Joachim, A.J. Mayne, G. Dujardin, F. Semond, L. Douillard, P. Soukiassian, Reconstruction of the Si-terminated b-SiC(1 0 0) surface, Thin Solid Films 318 (1998) 136. [15] V. Derycke, P. Soukiassian, A.J. Mayne, G. Dujardin, J. Gautier, Carbon atomic chains formation on the b-SiC(1 0 0) surface by controlled sp ! sp3 transformation, Phys. Rev. Lett. 81 (1998) 5868. [16] S. Hara, S. Misawa, S. Yoshida, Y. Aoyagi, Additional dimer-row structure of 3C-SiC(0 0 1) surfaces observed by scanning tunneling microscopy, Phys. Rev. B 50 (1994) 4548. [17] H.W. Yeom, Y.-C. Chao, S. Terada, S. Hara, S. Yoshida, R.I.G. Uhrberg, Surface core levels of the 3CSiC(0 0 1)3 · 2 surface: atomic origins and surface reconstruction, Phys. Rev. B 56 (1997) 15525R. [18] H.W. Yeom, I. Matsuda, Y.-C. Chao, S. Hara, S. Yoshida, R.I.G. Uhrberg, Hydrogen-induced 3 · 1 phase of the Si-rich 3C-SiC(0 0 1) surface, Phys. Rev. B 61 (2000) 2417R. [19] V. Derycke, P. Soukiassian, F. Amy, Y.J. Chabal, M.D. D’Angelo, H.B. Enriquez, M.G. Silly, Nanochemistry at the atomic scale revealed in hydrogen-induced semiconductor surface metallization, Nat. Mater. 2 (2003) 253. [20] J.J. Boland, Hydrogen as a probe of semiconductor surface structure: the Ge(1 1 1)-c(2 · 8) surface, Science 255 (1992) 186.

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