Hydrogenated amorphous and microcrystalline silicon deposited from silane-hydrogen mixtures

deposited from

silane-hydrogen mixtures

Ronald van Oort

TR diss

1709

deposited from

deposited from

silane-hydrogen mixtures

Proefschrift

ter verkrijging van de graad van doctor

aan de Technische Universiteit Delft,

op gezag van de Rector Magnificus,

prof.drs. P.A. Schenk, in het openbaar

te verdedigen ten overstaan van een com

aangewezen door het College van Dek,

op dinsdag 11 april 1989 te 14.00

door

RONALD CORNELIS VAN OORT

geboren te Delft

scheikundig ingenieur

r

TR diss

1709

1. General introduction 1 1.1. Historical background 1 1.2. Properties of hydrogenated amorphous 2

silicon 1.2.1. Introduction 2 1.2.2. Structural properties 3 1.2.3. Optical properties 5 1.2.4. Electrical properties 6 1.3. Preparation methods 7 1.4. Scope of this thesis 10

References 12

2. The rf glow discharge preparation of 15

hydrogenated amorphous silicon films

2.1. Introduction 15 2.2. The deposition apparatus 16

2.3. Fabrication of devices in a single 19 chamber reactor

References 23

3. Plasma deposition of hydrogenated amorphous 24

silicon: Effect of rf power and si lane dilution with hydrogen

3.1. Introduction 24 3.2. Description of the SiH./HL system

3.3. Growth rate of amorphous silicon as a 27 function of silane-hydrogen ratio and

rf power level

3.4. Conclusions 35

4.1. Introduction 40 4.2. Optical properties 41

4.2.1. Tauc optical gap and refractive index 41 4.2.2. Thickness dependence of the optical film 46

properties

4.3. Space-charge-limited current measurements 50 4.3.1. Measuring techniques for the density 50

of states

4.3.2. The space-charge-limited current 51 technique

4.3.3. Method 52 4.3.4. Density of localized states as a 58

function of deposition parameters

4.4. Photo-thermal deflection spectroscopy 60

4.4.1. Introduction 60 4.4.2. The photo-thermal deflection 61

spectroscopy technique

4.4.3. Experimental setup 62

4.4.4. Results 63 4.4.5. Comparison of space-charge-limited 66

current measurements and

photo-thermal deflection spectroscopy analyses

4.5. Reduction of the microstructure of 68 hydrogenated amorphous silicon

4.5.1. Introduction 68 4.5.2. Experimental setup 68

4.5.3. Results 69 4.5.4. Discussion 71

4.6.4. Explanation of the improved stability 79 with regard to the existing models

which explain the Staebler-Wronski effect

4.7. Conclusions 81

References 84

5. The deposition of rf glow-discharge 88

hydrogenated micro-crystalline silicon

5.1. Introduction 88 5.2. The deposition of hydrogenated microcrystalline 89

silicon

5.3. The role of hydrogen 93 5.4. Etching of hydrogenated amorphous silicon films 94

5.5. Etching of hydrogenated microcrystalline 101 silicon films

5.6. A growth model for hydrogenated 102 microcrystalline silicon 5.7. Conclusions 106 References 108 Summary 111 Samenvatting 114 List of publications 117 Acknowledgements 119 Curriculum vitae 120

1. GENERAL INTRODUCTION.

1.1. Historical background.

For a long time it has been thought that amorphous semiconductors could not be substitutionally doped. Normally amorphous films were prepared by sputtering or evaporation, yielding material with a high density of states and a void structure. These films are not suitable for electronic applications. The current wave of activities started in 1969 when Chittick [1.1] showed that amorphous silicon and germanium could be prepared by the rf glow discharge deposition technique, using hydrides and a substrate temperature of about 275 °C, and that doping was possible. These films had a low density, of gap states and were highly photoconductive [1.2].

In 1975 W.E. Spear [1.3] demonstrated that amorphous silicon could be doped both n- and p-type. The room temperature conductivity could be made to vary over six orders of magnitude. Doping of the material opened up possibilities for the fabrication of thin film electronic devices such as solar cells.. This meant the turning point for the investigations carried out in this field. A widespread effort to characterize the material structurally and chemically began, almost six years after the work on electronic properties commenced. It was discovered that amorphous silicon contained large amounts of bonded hydrogen, responsible for the low density of gap states. Normally glow discharge a-Si:H contains a lot more hydrogen

than is necessary to saturate the dangling bonds. Therefore amorphous silicon must be seen as a binary alloy of silicon and hydrogen. The delay in studying the structure of the a-Si:H material can be attributed to the failure to look back into the literature. Chemists such as Ogier (1880), Emeleus and Stewart (1945) and Jolly (1950-1965) have published on the glow discharge production of silicon subhydrides and that should have compelled an early study of the hydrogen issue [1.4]. Carlson [1.5] reported the first amorphous silicon solar cell with a p-i-n configuration in 1976 with an efficiency of 2.4%.

The interest in the a-Si:H material is based not only on its technological applications, but also on the possibility that it may provide us with a deeper understanding of the physics of amorphous semiconductors in general.

1.2. Properties of hydrogenated amorphous silicon.

1.2.1. Introduction.

It is generally accepted that hydrogenated amorphous silicon (a-Si:H) is indeed amorphous. Evidence for this comes from transmission electron microscopy (TEM), x-ray, and neutron diffraction, extended x-ray absorption fine

29

structure spectroscopy (EXAFS), Si-nuclear magnetic resonance and Raman scattering [1.4]. The properties, structural, electrical and optical, are strongly dependent on the deposition technique and the deposition conditions used. Under specific conditions (see Chapter 5) microcrystalline films are formed. These films contain crystals with a diameter in the range of 30- 300 A.

1.2.2. Structural properties.

In real a-Si:H the experimental evidence seems to indicate important structural inhomogenities. Nuclear magnetic resonance (NMR) studies indicate a two phase structural inhomogenity in which the hydrogen atoms are found in two kinds of environment [1.6]. The material can be described by means of a two-phase model: a phase, a, of low defect density and with hydrogen bonded mainly in the form of SiH (x=l) groups. This phase is imbedded in a second phase, b, of poor quality, i.e., a high defect density. In phase a a fraction of the monohydride is atomically dispersed, the remainder is clustered. Phase b contains large amounts of hydrogen, which is bonded not only in the form of SiH (x=l), but also as SiH (x>l) and as (SiH-) (n>l) chains. In good quality material phase i dominates. With TEM no structure can be resolved at a scale exceeding 10 - 30 A for films with a thickness not exceeding 1000 A [1.4], although two configurations for hydrogen still exist as indicated by NMR [1.6]. Low quality material has a columnar character, using TEM both phases are clearly distinguishable [1.7]. Scanning electron microscopy reveals a textured surface, regardless the quality of the film. The texture becomes more pronounced in film thickness. This can be explained by regarding the columns as cones [1.8]. With increasing thickness the cones overlap while the fraction of interfacial material decreases [1.9, 1.10]. The deposition conditions effect the coalescence of the columns and the amount of phase b in the intercolumnar region.

Infrared spectroscopy is used to identify the presence of SiH (x>l) groups and (SiH,,) (n>l) chains, and to determine the total hydrogen content [1.11]. Good quality material contains 2 - 1 5 atomic percent hydrogen only bonded as SiH

dihydride and polymer-like chains increases, as does the amount of SiH (x=l) clustering, while the film properties deteriorate. Generally one refers to the clustering by using the term microstructure.

A well known property of a-Si:H is the so-called Staebler-Wronski effect: under the influence of light, defects are created [1.12]. It is believed that Si-Si and Si-H bonds are broken. The breaking of an Si-H bond is followed by diffusion of the free hydrogen atom through the a-Si:H material. This free hydrogen atom is able to break weak Si-Si bonds.

Not only hydrogen but also fluor and chlorine atoms are able to saturate dangling bonds. Alloys such as a-Si:H,F are deposited for example from SiF./hL mixtures [1.13].

Adding dopants also influences the morphology of the films. With the dopant concentration the total hydrogen content increases [1.14].

(I)

(ID

Figure 1.1: Schematic drawing of a two-dimensional crystalline (I) and amorphous (II) silicon network.

A schematic illustration of a two-dimensional a-Si:H network is given in Figure 1.1. By comparing the a-Si:H network with a crystalline one the random character of the structure of a-Si:H is clearly revealed.

1.2.3. Optical properties.

The optical absorption edge, the sharp rise in the optical absorption at a characteristic energy, of amorphous semiconductors is difficult to determine experimentally. There is a strong dependence of the optical absorption on sample preparation for a specific amorphous semiconductor. Despite this, there remains a remarkable agreement among the shape of the absorption edge for different amorphous semiconductors. One can distinguish a shoulder in the optical absorption « (« < 1 cm" ) , an exponential rise (1 <

3 - 1 4 - 1 « < 10 cm ) , and a slowly varying regime (« > 10 cm )

[1.15].

From reflection and transmission measurements, using a method described by Tauc, an optical gap (Tauc optical gap) can be deduced. The Tauc optical gap is nearly equal to the mobility gap of aSi:H, with a value in the range of 1.6 -1.8 eV.

The hydrogen concentration influences the value of the optical gap. Generally it increases with the hydrogen content of the film [1.16]. A linear increase of the optical gap with hydrogen content up to 10 atomic % is found. For hydrogen concentrations above 10 atomic % the optical gap is found to be independent of the hydrogen concentration. Infrared analysis strongly suggests that the hydrogen which does not influence the optical bandgap is present in the material as SiH, [1.17].

Measurements for photon energies below the optical gap are performed using techniques like photoconductivity measurements or photo-thermal deflection spectroscopy (see chapter 4 ) , because of the low absorption in this region

[1.18]. Here the absorption is determined by band-to-band transitions from the valence band tail to the extended conduction band states and from the conduction band tail to the extended valence band states.

Direct optical transitions without involving a phonon are possible in a-Si:H due to the lack of long-range order [1.15]. As a result, the band-to-band absorption of a-Si:H is about an order of magnitude larger in the visible range compared to that of crystalline silicon. The strong optical absorption, the possibility of depositing a-Si:H over large areas and the low deposition temperature make a-Si:H an attractive material for solar cells and other large area electronic devices.

1.2.4. Electrical properties.

The structural and electrical properties of a-Si:H are closely related. Impurities, internal stress, structural defects such as vacancies leaving dangling bonds, variations in bond lengths and bond angles are the cause of the existence of states in the gap of the material. These states govern the electrical properties. A distinction is made between tail states and gap states which lie deeper. There are no sharply defined valence and conduction band edges. The gap states can be regarded as modified valence and conduction band states which overlap in the gap. Their existence is ascribed to the non-ideal, disturbed silicon matrix of a-Si:H (variation in bond lengths and bond angle distortion in the amorphous network).

A division is made in localized tail states and non-localized (extended) states. The mobility of charged carriers in localized states is zero, while carriers in the extended states are free. The mobility edge is defined as the energy separating localized and non-localized states in the conduction and valence band respectively.

Dangling bonds are the most important defects, giving rise to states deeper in the gap. The coordination number of a silicon atom is four. If it is bonded to less than four neighours, there is a free electron, the so-called dangling bond. Amorphous silicon without hydrogen has a large density of gap states. The gap states act as recombination centres, and the tail states as traps for electrons and holes. Reducing their concentration increases respectively the lifetime and the carrier mobility. Hydrogen is an effective terminator of dangling bonds, giving a-Si:H semiconductor properties and making doping possible. By adding dopants to a-Si:H, extra carriers condense in deep traps and cause a corresponding shift in the Fermi-level, enlarging the conductivity. Effective doping is possible only if the density of gap states is sufficiently low. About 3 0 % of the impurity atoms like boron and phosphorus are four fold coordinated and contribute to the change in the conductivity [1.19, 1.20]. The low doping efficiency must be ascribed to the ease with which impurities are able to adopt their optimal coordination in amorphous material.

1.3. Preparation methods.

The most commonly used deposition method is the radio frequency glow discharge technique for the production of a-Si:H thin films. A description of depositing a-a-Si:H films by means of the rf glow discharge method is given in chapter

two. There are numerous other methods of obtaining a-Si:H films. The electrical and optical quality of the films produced by other methods is at best comparable with rf glow discharge a-Si:H film. The methods are: reactive sputtering in a hydrogen atmosphere [1.21], reactive ion beam sputtering (RIBS) [1.15], DC [1.22], microwave [1.23] and VHF glow discharge [1.24], chemical vapor deposition (heterogeneous CVD) of silane [1.25] or higher silanes [1.26], homogeneous CVD (HOMOCVD) [1.27], photo-CVD [1.28] and hydrogen remote plasma CVD (HRPCVD) [1.29].

Reactive sputtering, dc and rf glow discharge films suffer from ion bombardment during growth, which leads to films exhibiting a columnar structure and microvoids [1.6]. Generally in pyrolytic decomposition there is no ion bombardment. However, with CVD a-Si:H grown at 420-530 C from silane, post-hydrogenation is necessary to eliminate dangling bonds. When disilane is used at 380-500 C, a higher hydrogen content in as-deposited films is obtained. In HOMOCVD the construction of the reactor is such that the gas is decomposed at a higher temperature (600-700 °C) while the substrate temperature can independently be controlled, and is mostly between 30-300 °C [1.30].

To promote CVD reactions at temperatures below 300 C at the expense of PVD reactions to obtain a material with a high density (no columnar growth), low defect concentration and a good step coverage, HRPCVD is developed. Hydrogen radicals are produced in a seperate chamber and led to the reaction chamber where mixing with silane occurs. The hydrogen radical concentration can be controlled independently and the growing film is not in contact with the plasma. In photo-CVD reactive fragments are produced by mixing mercury with the gas and UV radiation of a mercury lamp is used for initialisation of the reactions.

RIBS is a new technique for film growth. In a high vacuum a target is bombarded by a high flux ion beam. The substrate can be placed outside the plasma environment. Due to the low pressure, gas phase reactions are eliminated. The reactive gas pressure can be varied over a wide range.

Glow discharge deposition is sometimes carried out under UHV conditions in a multi-chamber system to study the effect of impurities such as oxygen, carbon, and nitrogen on the Staebler-Wronski effect and of cross-contamination of dopants. [1.31]. The light induced degradation has been

20 shown to be a bulk property below impurity levels of 10

3

atoms per cm . Thus, deposition in an UHV system is not enough to prevent degradation [1.32, 1.33].

One of the major problems attached to solar cells and other thin film devices based on a-Si:H in becoming economically attractive is the low deposition rate for good quality material (1-2 A/s). A target would be around 20 A/s [1.34]. The deposition rate can be increased by raising the rf power input and/or the SiH. gas flow, by replacing SiH. by higher silanes, by mixing silane with inert noble gases or by using novel deposition techniques such as the plasma confinement method and the photo-CVD technique. So obtained films show in general poor structural order and poor opto­ electronic properties due to gas phase polymerization, stronger ion bombardment damages and other undesired side-effects such as large area non-uniformity. Knights [1.7], Jackson and Amer [1.18] have shown that the best quality films (films with a low defect density) are obtained at the lowest possible rf power using pure silane. These are widely known standard conditions for obtaining good quality a-Si:H. Most plasma systems operate at 13.56 MHz. For this frequency the equipment is readily available and the permitted radiation levels are higher. Studies on microwave (2.45 GHz) and VHF (25-150 MHz) glow discharge a-Si:H film show that

growth rates of one order of magnitude and larger can be obtained. Changing the plasma excitation frequency is a promising way of increasing the deposition rate. By changing the excitation frequency the electron density in the plasma is increased. The quality is still not comparable with the best rf glow discharge a-Si:H films.

1.4. Scope of this thesis.

This thesis consists of several parts which are closely related. The aim is to promote a better understanding of the deposition of hydrogenated amorphous and microcrystalline silicon, especially the role of hydrogen in the deposition process, and of the relationship between the structure of amorphous silicon and its electrical and optical properties. The importance of this relationship is frequently underestimated. Most of the investigations carried out so far have only dealt with electrical or structural properties and not with both.

First an outline is given o f t h e deposition of a-Si:H by means of the rf glow discharge technique and attention is paid to the fabrication of devices in a single-chamber deposition system as is currently being used by our group. Contamination by dopant gases of the intrinsic layers of devices, such as solar cells, might severely limit the device performance.

A description is given of the influence of diluting the silane feed gas with hydrogen in combination with variations in the rf power density. Increasing the latter parameter opens the opportunity to raise the deposition rate of a-Si:H, but the quality of the film deteriorates. It will be shown that by adding hydrogen it is possible to deposit good quality material with an increased growth rate. This is

important in producing a-Si:H thin film devices such as solar cells in an economical way. The a-Si:H solar cells suffer from light-induced degradation decreasing the conversion efficiency. Using silane-hydrogen mixtures leads to the formation of a more stable a-Si:H. An explanation will be given based on the structure of the material. Infrared spectroscopy opens the possibility of studying the microstructure of a-Si:H. It will be shown that the results obtained from the optical and electrical characterization, combined with structural analyses, leads to the conclusion that the microstructure is an important factor determining the properties of a-Si:H.

Microcrystalline p and n+ contacts are used in solar

cells to raise the efficiency. These films are deposited under conditions widely different from the conditions under which a-Si:H is deposited. A high rf power level and silane strongly diluted in hydrogen are necessary to obtain uc-Si:H material. The growth process of uc-Si:H is generally described as a near chemical equilibrium situation. The growth rate of uc-Si:H films is often used as an indication that the system is close to a chemical equilibrium. It will be shown that this is not a clear description. Both a- and uc-Si:H are deposited with the same growth rate. For this reason it is better not to define the conditions under which uc-Si:H is formed as being only a situation close to chemical equilibrium and using the growth rate as an indication that the system is close to an equilibrium. Based on etch rate measurements of a- and uc-Si:H in a hydrogen rf plasma the growth of uc-Si:H is described.

References.

[1.1]. R.C. Chittick, J.H. Alexander, H.F. Sterling, "The preparation and properties of amorphous silicon", J. Electr. Chem. S o c , 116, 1969, 77.

[1.2]. W.E. Spear, P.G. LeComber, "Investigation of the localised state distribution in amorphous Si films", J. Non-Cryst. Solids, 8-10, 1972, 727.

[1.3]. W.E. Spear, P.G. LeComber, "Substitutional doping of amorphous silicon", Solid State Comm., 17, 1975, 1193.

[1.4]. J.C. Knights in Topics in Applied Physics, vol. 55, Springer Verlag, Berlin, vol. ed. J.D. Joannopoulus, G. Lucovsky, 1983, chapter 2.

[1.5]. D.E. Carlson, C.R. Wronski, "Amorphous silicon solar cell", Appl. Phys. Lett., 28, 1976, 671.

[1.6]. H. Fritzsche, "Heterogenities and surface effects in glow discharge deposited hydrogenated amorphous silicon", Thin Solid Films, 90, 1982, 119.

[1.7]. J.C. Knights, R.A. Lujan, "Microstructure of plasma-deposited a-Si:H films", Appl. Phys. Lett., 35(3), 1979, 244.

[1.8]. R.C. Ross, A.G. Johncock, A.R. Chan, "Physical microstructure in device-quality hydrogenated amorphous silicon", J. Non-Cryst. Sol., 66, 1984, 81.

[1.9]. P. Rava, "Influence of thickness and hydrogen concentration on optical and electrical properties of a-Si:H thin films", J. Vac. Sci. Technol. A, 4(5), 1987, 1795.

[1.10]. R. Messier, J.E. Yehoda, "Geometry of thin film morphology", J. Appl. Phys., 58(10), 1985, 3739. [1.11]. P.J. Zanucchi in Semiconductors and Semimetals, vol.

21 B, A.P., New York, vol. ed. J.I. Pankove, 1984, chapter 4.

[1.12]. D.L. Staebler, C.R. Wronski, "Optically induced conductivity changes in discharge-produced hydrogenated amorphous silicon", J. Appl. Phys., 51, 1980, 3262.

[1.13]. J. Mort, F. Jansen, "Plasma deposited thin films", CRC Press Inc., 1986.

[1.14]. M. Pruppers, Thesis, RUU, Utrecht, 1988.

[1.15]. P. Cody in Semiconductors and Semimetals, vol. 21 B, A.P., New York, vol. ed. J.I. Pankove, 1984, chapter 2.

[1.16]. J. Shirafuji, S. Nagata, M. Kuwagaki, "Effect of hydrogen dilution of silane on the opto-electronic properties in glow-discharged hydrogenated silicon films", J. Appl. Phys., 58(9), 1985, 3661.

[1.17]. C. Maessen, Thesis, RUU, Utrecht, 1988.

[1.18]. N.M Amer, W.B. Jackson in Semiconductors and Semimetals, vol. 21 B, A.P. ,New York, vol. ed. J.I. Pankove, 1984, chapter 3.

[1.19]. W. Beyer, H. Overhof in Semiconductors and Semimetals, vol. 21 C, A.P. New York, vol. ed. J.I. Pankove, 1984, chapter 8.

[1.20]. P.G. LeComber, W.• Spear in Topics in Applied Physics, vol. 36, Springer Verlag, Berlin, vol. ed. M.H. Brodsky, 1979, chapter 9.

[1.21]. W. Paul, A.J. Lewis, G.A.N. Connel, T.D. Moustakis, "Doping Schottky barrier and p-n junction formation in amorphous germanium and silicon by rf sputtering", Solid State Comm., 20, 1976, 969.

[1.22]. R. Robertson, A. Gallagher, "Mono- and disilicon radicals in silane and silane-argon dc discharges", J. Appl. Phys., 59(10), 1986, 3402.

[1.23]. S. Kato, T. Aoki, "High rate deposition of a-Si:H using electron-cyclotron resonance plasma", J. Non-Cryst. Sol., 77 & 78, 1985, 813.

[1.24]. H. Curtins, N. Wyrsch, A.V. Shah, "High rate deposition of amorphous hydrogenated silicon: Effect of plasma excitation frequency", Electr. Lett., 23(5), 1987, 228.

[1.25]. J. Magarino, D. Kaplan, A. Friederich, A. Deneuville, "Doping effects on post-hydrogenated chemical vapour-deposited amorphous silicon", Phil. Mag. B, 45, 1982, 285.

[1.26]. S.C. Gau, B.R. Weinberger, M. Akhtar, Z. Kiss, A.G. MacDiarmid, "Preparation of amorphous silicon films by chemical vapour deposition from higher silanes", Appl, Phys. Lett., 41, 1982, 1146.

[1.27]. B. A. Scott, J.A. Reimer, P.A. Longeway, "Growth and defect chemistry of amorphous hydrogenated silicon", J. Appl. Phys., 54(12), 1983, 6853.

[1.28]. A.E. Delahay, "High rate photochemical deposition of amorphous silicon from higher silanes", J. Non-Cryst. Sol., 77 & 78, 1985, 833.

[1.29]. S. Oda, "The role of hydrogen radicals in the growth of a-Si and related alloys", Jap. J. Appl. Phys., 25(3), 1986, L188.

[1.30]. M. Akhtar, V.L. Dalai, K.R. Ramaprasad, S. Gau, J.A. Cambridge, "Electronic and optical properties of amorphous Si:H films deposited by chemical vapour deposition", Appl. Phys. Lett., 41, 1982, 1146. [1.31]. S. Tsuda, "Preparation and properties of high

quality a-Si films with a super chamber (separated ultra-high vacuum reaction chamber)", Jap. J. Appl. Phys., 26(1), 1987, 33.

[1.32]. C.C. Tsai, J.C. Knights, M.J. Thompson, " 'Clean' a-Si:H prepared in a UHV-system", J. Non-Cryst. Solids, 66, 1984, 45.

[1.33]. M. Stutzmann, W.B. Jackson, C.C. Tsai, "Light-induced metastable defects in hydrogenated amorphous silicon: A systematic study", Phys. Rev. B, 32, 1985, 23.

[1.34]. D.E. Carlson in Semiconductors and Semimetals, vol. 21 D, A.P. Press, New York, vol. ed., J.I. Pankove, 1984, chapter 4.

2. THE RF GLOM DISCHARGE PREPARATION OF HYDROGENATED AMORPHOUS AND MICROCRYSTALLINE SILICON FILMS.

2.1. Introduction.

The most commonly used deposition method is the radio frequency glow discharge technique for the production of a-Si:H thin films. To obtain a better uniformity over a larger area capacitative coupled parallel plate reactors are used. This type of reactor is easily scaled up and it is even possible to deposit a-Si:H using a continuous system. The deposition is carried out at relatively low temperatures, making a variety of substrates suitable.

Often silane is used as the gas from which intrinsic a-Si:H films are grown. This yields films which are slightly n-type. By mixing the silane gas with dopant gases such as arsine, phosphine and diborane, one is able to deposit n-and p-type films. Mixed alloys with bn-andgaps varying from 1.1 to 5.2 eV are obtained by adding gases such as germane (decreased energy gap), methane (increased energy gap) or ammonia (increased energy gap) to the silane feed gas. The bandgap is dependent on the gas ratio used. It is believed that through an adjustment of the gas composition the electronic properties of the films can be tailored for device applications. However defect saturisation by hydrogen is far less effective in mixed alloys than in a-Si:H.

An example of the use of alloys is found in solar cells. By stacking cells with successively narrower bandgaps the solar spectrum is utilized more efficiently compared to

single bandgap cells [2.1]. A cell with the structure glass/TCO/p+-uc-Si:H:F/i-Si:H:F/n+/P+-Si:H:F/i-SiGe:H:F/n+

-Si:H:F (TCO is transparant conductive oxyde) had a reported efficiency of 12.5% [2.2].

The p -uc-Si:H window side layer is a special feature of this cell. Microcrystalline silicon is a material consisting of two phases, microcrystals with a diameter of about 10 nm embedded in an amorphous network of a-Si:H. The uc-Si:H material, doped up to 1% with boron, does not suffer from bandgap narrowing caused by doping. For comparison, the optical gap of a-Si:H:B is 1.35 eV while for uc-Si:H:B the optical gap has a value of 1.65 eV. The built-in electric field for solar cells is larger when uc-Si:H contact layers are used compared to conventional solar cells. The electrical conduction and transmission of light for doped uc-Si:H is high compared to n and p a-Si:H contact layers. Also the refractive index is smaller which reduces the surface reflectance of incident light [2.3, 2.4]. All these aspects lead to an enlargement of the efficiency of solar cells when uc-Si:H is used as a material for the contact layers. The properties of uc-Si:H are determined by the microcrystals, by the amorphous phase and the boundary region between the microcrystals and the amorphous web. The crystals do not contain hydrogen. Hydrogen is located on the surface of the crystals, mainly in the form of polyhydrides [2.5].

2.2. The deposition apparatus.

The apparatus used for the deposition of a- and uc-Si:H films is schematically given in Figure 2.1. The system consists of three main parts:

-The reaction chamber, fabricated by the Japanese firm Samco, model PD 10.

-The vacuum pump system, in our case a simple rotary pump 3

with a capacity of 45 m /h which keeps the pressure during deposition at around 0.3 torr.

-The gas handling system, fabrication ASM. The gas flow rates are controlled by electronic mass flow controllers with an accuracy of 1%. Mixing of the gases takes place in the common gas tube leading from the gas manifold to the reactor.

The combined leak/desorption rate of the system is below 2.10 torr.l./s.. This gives a background contamination level in the gas below 5 parts per million when a total gas flow rate of 100 standard cubic centimeter per minute (seem) is used.

The reactor is an rf glow discharge parallel plate reactor. The 13.56 MHz generator (ENI ACG 5) is connected to the upper electrode through a matching network and the substrate is placed on the heated grounded lower electrode. Also the walls of the reactor are grounded which enlarges the surface of the substrate electrode A ~ considerably.

sf J

Between the bulk of the plasma and the electrodes there is a sheath voltage. Due to the electric field in the sheaths the positive ions in the sheath regions are accelerated towards the electrodes. The resulting sputtering deteriorates the quality of the a-Si:H film. Between the sheath voltages V and the electrode areas there is a relation [2.6]. It is written as:

f*=£¥ (2.D

vrf Msf

The surface of the powered electrode is presented by A -. Due to the construction, both the walls of the reactor and

the substrate electrode are grounded, A - » A - and V - is small. The influence of the ion bombardment is minimized.

The volume of the stainless steel reactor is 17 1. The total flow rate, pressure, substrate temperature and electrode distance were kept constant at 100 seem, 0.3 torr, 523 K and 5 cm. respectively for the films on which this thesis reports, unless otherwise stated. The power level in

2

cm is calculated with reference to the surface of the upper electrode. rf-generator matching network P * substrates Y///A] Y///A / I < ni IS it-let - T ^ ' pump sj /St em -«■ -*m _^ ««■ L - — SiH4 SiH4 PH3 B2H6 B2H6 gas system 100% (2% in H2) (1% in SiHJ (1% in SiH4) (200 ppm in H2) Ar H2

2.3. Growing of devices in a single chamber system.

The Samco-reactor consists of a single chamber. When producing devices, such as n-i-n and p-i-p structures, the intrinsic layer is formed after the deposition of doped layers. This leads to an intrinsic layer being more or less contaminated with dopants. Also the n/i or p/i interface will be less sharp. The impurity profile is determined by three factors. First by residual dopant gas in the reactor. Secondly by diffusion of dopant from the doped contact layers into the i-layers (auto-doping). These two factors determine the sharpness of the interface between the doped and undoped layer. The third factor, desorption of gases from the chamber walls, leads to a nearly constant level of contamination of intrinsic layers.

Autodoping can be reduced by lowering the deposition temperature which reduces the diffusion of dopant from the doped layer into the intrinsic layer. However reducing the deposition temperature strongly and negatively affects the quality of a-Si:H. A second possibility is reducing the deposition time by raising the growth rate of the film.

The quantity of impurities present in the gas phase can be reduced by gas flushing. As a first approximation the reactor can be considered as a completely stirred tank reactor due to the low pressure in the system. The concentration C of dopant gas in the deposition chamber as a function of time is given by [2.7]:

C=CQ exp(-t/T) (2.2)

where CQ is the initial concentration and r the residence

time of the gases in the reactor. Rewriting equation (2.2) in the logarithmic form:

t=-Tln(C/C0) (2.3)

it is seen that the number of times the gas must turn over depends on the initial and final concentration of dopant gas. Reducing the phosphine or diborane level by a factor of

5

10 in a 17 1 system with a flush gas flow of 100 seem system takes 32 seconds.

However, this calculation does not take into account the release by desorption from the reactor walls of dopant gas. This proces is determined by unknown desorption times. For this reason the following procedure was executed. In the reactor a rotatable shutter was mounted. After the deposition of doped layers a flush sequence was carried out followed by a dummy growth using pure si lane and a relatively high rf power density of 36 mW/cm . Thus all the contaminated inner parts of the reactor were covered with an intrinsic layer of a-Si:H.

By secondary ion mass spectrometry (SIMS) the dopant profiles of p/i interfaces were measured for films grown under different conditions, see Figure 2.2. SIMS is a highly sensitive (detection level about 1 ppm.) analysis method, but it is not able to distinguish between ' P and SiH fragments because of the equal mass of the species. Therefore p/i interfaces were used. The dopant efficiency of PH, and B-Hg is not equal and thus the structure of the p-and n-type film will differ slightly as well, but the information obtained is useful for a general evaluation of the shutter concept. From Figure 2.2 it is clear that with the shutter the p/i interface is very sharp. The different deposition conditions of the films 1 and 2 have no influence on the dopant profile.

i-layer p-layer

0.30 0A5 thickness

(|im)-0.60

Figure 2.2: SIMS dopant profile of p/i-interfaces deposited under different conditions. Solid line: the deposition rate of the i-a-Si:H layer is l.OA/s. Dashed line: 4.8 A/s. The thickness of the p-layer is 1500 A. Along the y-axis the boron concentration is given.

The SIMS analysis made it possible to estimate the level 19 of contamination with carbon and oxygen, which is 3.10

3 19 3 atoms per cm for carbon and 6.10 per cm for oxygen, independent of the deposition conditions. This knowledge is important because contamination with carbon and oxygen might play a role in the light-induced degradation of a-Si:H films.

The shutter was also used to shield the substrates from the plasma when ignition took place. After 1 minute it was assumed that the plasma was stabilized and the substrates

were brought into contact with the plasma. It takes about half a minute to reach by diffusion a new dynamic equilibrium in the concentration of gas-phase species inside and outside the plasma-zone. After this period it is expected that a stable discharge has been reached [2.8]. In this way it also was possible to reduce the number of defects in the film near the substrate, as proven by a combination of surface photo voltage measurements and determination of the diode quality factor of i-n devices

References.

[2.1]. Amorphous Semiconductor Technologies & Devices, North-Holland, Amsterdam, ed. Y. Hamakawa, 1984. [2.2]. S. Guha, J. Yang, P. Nath, M. Hack, "Enhancement of

the open-circuit voltage in high-efficiency amorphous silicon solar cells", Appl. Phys. Lett., 49(4), 1986, 218.

[2.3]. Y. Uchida, T. Ichimura, M. Ueno, H. Haruki, "Microcrystalline Si:H film and its application to solar cells", Jap. J. Appl. Phys., 21(9), 1982, L586.

[2.4]. A. Matsuda, "Boron doping of hydrogenated silicon films", Jap. J. Appl. Phys., 20(3), 1981, L183. [2.5]. C.C. Tsai, R. Thompson, C. Doland, F.A. Ponce, G.B.

Anderson, B. Wacker, "Transition from amorphous to crystalline silicon: effect of hydrogen on film growth", Proc. 1988 MRS Spring Meeting, Symposium Amorphous Silicon Technology, vol. 118, Pittsburgh, USA, 1988, 49

[2.6]. B. Chapman in Glow Discharge Processes, J. Wiley and Sons, New York, 1980.

[2.7]. J. Levenspiel, "Chemical reaction engineering", J. Wiley and Sons, New York, 1980.

[2.8]. Y. Nakayama, T. Ohtsuchi, M. Nakano, T. Kawamura, "Initial transient phenomena in the plasma decomposition of silane", J. Non-Cryst. Sol. 77 & 78, 1985, 757.

[2.9]. J.C. van den Heuvel, R.C. van Oort, B. Bokhorst, M.J. Geerts, J.W. Metselaar, "Influence of the defect density of amorphous silicon at the substrate

interface on the Schottky-barrier characteristics", Proc. of the 15th International Conference on Defects in Semiconductors", Budapest, Hungary, 1988, in press.

3. PLASMA DEPOSITION OF HYDROGENATED AMORPHOUS SILICON: EFFECT OF RF POWER AND DILUTION OF SILANE WITH HYDROGEN.

3.1. Introduction.

Mass spectrometry [3.1], optical emission spectroscopy, infrared absorption and emission [3.2] are used for studying the plasma chemistry and ellipsometry for the in situ studying of the growing film surface [3.3]. These studies have led to a better understanding of the deposition of a-Si:H. Still, the growth process is not fully understood, which makes it necessary to determine the most favourable deposition conditions empirically to obtain films with good optical and electronic properties. The properties of glow-discharge a-Si:H are dependent on the deposition parameters. The influence of parameters on the quality of a-Si:H, such as rf power level [3.4], pressure [3.5], substrate temperature [3.6], substrate bias [3.7], electrode distance

[3.8] and the use of a triode instead of diode reactor [3.9] has been well established, just as has the dilution of silane gas with inert carrier gases, such as the noble gases [3.10].

The rf power level must be just sufficient to maintain the glow discharge. A rise in the rf power level increases the number of defects such as microvoids and promotes a columnar film morphology. High rf power levels lead to strong ion bombardment of the growing film.

In general pressures in the range of 0.05 - 1.0 mbar are used. A low pressure, below 0.05 mbar, increases . the

influence of ion bombardment. Electron and ion energies are limited primarily by collisions with the substrate and reactor surfaces and not by collisions in the gas phase. Pressures which are too high lead to gas phase polymerization resulting in the formation of dust particles which give rise to pinholes in the film.

High quality films are obtained with substrate temperatures in the range 200 - 300 °C. A higher substrate temperature promotes the surface mobility of species adsorbed on the growing film surface. At substrate temperatures above 300 °C hydrogen evolves from the film which results in an increase of the density of gap states because fewer dangling bonds are passivated. A low substrate temperature, just as a high pressure and rf power level, leads to the formation of films with a high total hydrogen concentration. Not only SiH groups, but also SiH2 groups and

(SiH„) (n>l) chains are present, which gives rise to a material with poor electronic properties.

A low background concentration of contaminants in the films, which originates from leakages and desorption, is obtained by a high gas flowrate. Choosing the flow rate too high results in transportation of active species outside the reactor before diffusion can take place to the substrate surface, which results in a decrease of the growth rate. Also, the extent of the departure of the system from the quasi chemical equilibrium strongly influences the deposition kinetics and the attainment of a quasi-equilibrium depends on whether or not the residence time of

species is smaller than the characteristic time of the reaction or the overall reaction time constant. Normally a flow rate is choosen in the range 10 - 100 seem.

Diluting the silane gas with inert gases such as He, Ne, Kr, and Ar deteriorates film properties. The influence of the carrier gas is dependent on the molecular mass of the

gas molecules. The deterioration is stronger for dilution with gases having a higher molecular mass.

The investigations carried out so far have led to conditions widely known to yield good quality material. The conditions are the use of pure silane and an rf power level just sufficient to maintain the plasma. The growth rate of the films formed under these circumstances is about 1 A/s, a value useful under laboratory conditions, but a bottleneck for an economic application of a-Si:H, as has already been mentioned in Chapter 1.

In this thesis a description is given of the effect of the dilution of silane with hydrogen in combination with the rf power level. The result of the study is given in this and the following chapter. The hydrogen concentration is of major importance. Extra hydrogen forms a tool to influence the deposition of intrinsic a-Si:H films. By dilution and by adjusting the rf power level one can significantly alter the density of activated species, thereby changing the surface reactions and plasma chemistry. These parameters strongly affect the deposition rate [3.11] and the opto-electronic properties [3.12]. Also, hydrogen is generated during the deposition process, through which a certain degree of self dilution occurs, dependent on the degree of conversion of the process gas. In this chapter the influence of the studied process parameters on the deposition rate is described in terms of shifts in quasi-chemical equilibrium

[3.13] and is described by using only the most significant reactions occuring in the plasma which contribute to the film growth process [3.14, 3.15].

3.2. Description of the SiH./Hp system.

Though the plasma deposition of a-Si:H from silane is a rather complex process, the net reaction can be simply expressed by the following equation:

i

SiH4 (g) ;;;> Si:Hx(s) + (2-x/2) H2(g) (3.1) 2

The forward reaction (1) means depositing a layer and the backward reaction (2) etching. Reaction (1) is controlled by the silane concentration and the rf power level. The decomposition reaction rate r, of the forward reaction is

given by [3.16]:

r: = kj ne [SiH4] (3.2)

with:

n = electron density

e J

k, = rate constant for the dissociation reaction [SiH.] = silane concentration

The rate constant is equal to:

kj = (2/me)1/2|0°E1/2f(E)(71(E)dE (3.3)

with:

m = electron mass e

E = energy of electrons present in the gas phase f(E) = normalized electron distribution function

a, = cross section for the forward reaction of

From formula (3.3) it can be seen that a change in the pressure, influencing the electron temperature, primarily affects the rate constant. The partial pressure of the reactant gas determines the magnitude of [SiH.] in equation (3.2). The electron density n is related to the electric field. The electron density should scale linearly with increasing discharge power R^ [3.17]:

Rf

ne = — 2 <3-4>

e 7T IT h 2a.[n,]

J J

with:

R = the radius of the reactor h = the height of the reactor

a.= the rate at which electrons lose their energy <J

to species j per unit volume [n.] = the density of species j

Thus, using formula (3.2) it is evident that r, can be increased by increasing the rf power or the partial pressure of SiH. or by decreasing the total pressure.

Under the conditions of chemical equilibrium no layers are formed. The net reaction rate is zero. Any departure from chemical equilibrium can be controlled by the process parameters. The dependence of the growth rate Rg can be described in terms of shifts in chemical equilibrium. This is in line with Vepreck's suggestion [3.18] that a situation near chemical equilibrium, a quasi-equilibrium situation with high but almost equal reaction rates r, and r2 , is

necessary to grow microcrystalline silicon.

The neutral dissociation of the silane exceeds the dissociative ionisation by a factor of 50-100 due to a lower threshold energy. The rate constants for these processes in

-10 -9 3 an rf glow discharge are respectively 10 - 10 cm /s and 10"1 2- 1 0 "1 1 cm3/s [3.19]. The main channels for

dissociation of silane into neutral particles by electron impact are [3.19, 3.20]:

SiH4 -?-> SiH3 + H (3.5)

SiH4 ---> SiH2 + 2H (3.6)

The channel yielding SiH2 + H? has a smaller probability

compared to the channels (3.5) and (3.6) because in the presence of a considerable excess of energy the hydrogen molecules can be expected to dissociate. The hydrogen atoms have little opportunity to form an HL molecule [3.20]. Dissociation products like SiH + 3H and Si + 4H are energetically unfavorable [3.20].

The generation rate of SiH3 and SiHp radicals can be

altered by participation of the radicals in gas-phase and surface reactions. The most important reactions following the neutral dissociation (3.5) and (3.6) are:

H + SiH4 ---> SiH3 + H2 (3.7)

SiH2 + SiH4 ---> Si2H6 (3.8)

SiH3 + SiH3 ---> Si2H6 (3.9)

The majority of the radicals take part in these reactions because these reactions are more rapid than other radical -silane or radical-radical reactions. The rate constants of (3.7), (3.8) and (3.9) respectively are 4xl0"1 3- 8xl0~1 2

cm3/s, 5 x l 0 "H- lxlO"10 cm3/s and 7xl0"12- 4 x l 0 "n cm3/s

[3.19] for typical glow-discharge conditions.

The electrode diameter of 20 cm is considerably larger than the electrode distance of 5 cm. A one-dimensional equation can be used to calculate the time-average flux J of radical a to the electrode surfaces :

L

J, = J [G.(x) - R (x)] dx (3.10)

a 0 a a

with:

G,(x) = the generation rate of radicals of type a

a

R (x) = the recombination rate of radicals of type a

a

L = the electrode distance.

The actual growth rate is dependent on the sticking coefficient of species a at the surface of the growing film.

3.3. Growth rate of amorphous silicon as a function of si lane-hydrogen ratio and rf power level.

Amorphous silicon films were deposited using a range of 2

rf power levels (15-482 mW/cm ) and various si lane-hydrogen mixtures (20-100 vol.% Siti.). The deposition rate Rg of

intrinsic a-Si:H films was measured by weighing (accuracy 10%) and by optically measuring the transmission and reflection (accuracy 1%) of intrinsic a-Si:H films deposited on Corning 7059 glass substrates.

The growth rate dependence of the silane-hydrogen ratio at low power is quite different compared to the behaviour at higher levels, giving an indication that major reaction pathways have changed, see Figure 3.1.a. Energy barriers, preventing certain reaction pathways occurring, are passed at higher rf power levels, thus altering the plasma chemistry and/or the surface reactions. Due to the complexity of the plasma chemistry and the surface processes it is not possible to indicate exactly the changes that are occurring.

At a power level of 15 mW/cm an increasing silane concentration from zero up to 55 % results in a higher Rg,

the net result of the forward and backward reaction rate of reaction (3.1). For silane concentrations of 55% and more, decreasing Rg means, in terms of Vepreck, going back closer to chemical equilibrium. Departures from the point of chemical equilibrium are dependent on the silane dilution. The existence of the broad maximum in curve 3 is in agreement with P.E. Vanier et al. [3.11] although they found a much higher maximum at a concentration of 25% silane. This can be ascribed to the different deposition systems used. The maximum was attributed to changes in the electron energy distribution in the plasma. At higher power no maximum is found at medium concentrations. Equation (3.2), indicating a linear relationship between the silane partial pressure and the growth rate, is then more or less obeyed for all gas mixtures.

Figure 3.1.a. The influence of rf power level and dilution of silane with hydrogen on the growth rate Rg of intrinsic

2 2 amorphous silicon. Curve 2: 62 mW/cm , curve 2: 36 mW/cm ,

2

O 60 120 180 240 300 360 420 480

2

rf power (mW/cm )

»-Figure 3.1.b. The influence of rf power level and dilution of silane with hydrogen on the growth rate Rg of intrinsic amorphous silicon. Curve 4: 100 vol.% SiH^, curve 5: 55 vol.% SiH., curve 6: 33 vol.% SiH..

The contribution of the SiH? radicals to the amount of

material deposited is limited to less than 2% [3.20]. The main contributors are the SiH3 radicals, more than 98%. The

SiH, radicals do not react rapidly with silane. Their characteristic reaction distance during diffusion is:

( D / krN )1 / 2 » L (3.11)

with:

D = the diffusion coefficient k = the reaction rate constant N = the concentration of silane. L = the electrode distance

The SiH3 radicals should diffuse to the substrate without

reacting with silane. The term R,(x) in (3.10) can be a

disregarded. Therefore the growth rate Rg has to be proportional to the generation rate G, of SihU radicals. However, the reaction (3.9) yielding Si?Hg becomes more

important with increasing power because of the increasing SiH, density in the plasma. This results in infringement of the linear dependence between Rg and G3. Due to the linear

relationship between G, and the rf power level [3.17] this results in a less than linear increase of Rg with rf power as experimentally observed, see Figure 3.1.b. A second possibility for explaining the infringement is depletion of silane occuring in the gas phase. Due to the enhanced growth rate the actual concentration of silane in the reactor will be smaller than the silane concentration in the feed gas. The influence of depletion of silane is larger at higher rf power densities.

2

The decrease at a power exceeding 295 mW/cm might be caused by a greater increase in the hydrogen density compared with the silane radical density. The hydrogen atoms

produced by reactions (3.5) and (3.6) rapidly react with silane through reaction (3.7). At the same time a relative large diffusion coefficient for H is expected. The diffusion coefficient of the H atom in silane is not available in literature and hence it is not possible to evaluate the characteristic reaction distance during diffusion. But knowledge of the value of the diffusion coefficient is not essential. It seems a reasonable assumption that under typical discharge conditions most H atoms convert the SiH. in the gas phase to SiH., [3.20]. For H atoms which do reach the surface there are two possibilities to react with a-Si:H. The abstraction of an H atom from a surface Si-H bond and and the breaking of an Si-Si bond thereby creating volatile SiH3 or SiH.. In the first case the dangling Si

bond abstracts an H from SiH. producing an SiH, radical in the gas which can contribute to the film growth. If the second channel occurs the growth rate is decreased due to etching of the surface. The quasi-equilibrium shifts to the left.

It also is important that the reaction (3.9) is quadratic in the SiH., radical density while reaction (3.7) is linear in the H atom density. This means that the term R (x) in (3.10) is increased more strongly for SiH, than for H with increase in their densities. This causes an increase in the H flux with respect to the SiH3 flux with increasing

rf power. At high power this ratio can be so large that the growth rate decreases because of enhanced etching of the a-Si:H surface. This probably is the explanation of the experimentally observed maximum in the growth rate-power

2

dependence around 295 mW/cm . This etch effect is also a third possibility for explaining the observed infringement.

The above explanation of the effect of hydrogen is consistent with the fact that the hydrogen content of the films deposited with pure silane increases with the rf

power, see Table 3.1. A saturation of the film by hydrogen does appear, being the first step in the film etching mechanism, see Chapter 5. Also the refractive index of the films deposited at high power is less than the refractive index of the films deposited at low power, see Section 4.2. Microcrystalline silicon has a smaller refractive index and this material is obtained under conditions that etching of the growing film surface is of major importance.

7ab7e 3.1: Total hydrogen content of a-Si:H films deposited using pure silane as a function of the growth rate Rg. The hydrogen content has been measured by means of nuclear reaction analysis (NRA).

film no. Rg (A/s) [H]tot^a t-% * ° 5 a t - %)

465 467 483 0.8 6.0 10.5 8.0 12.5 23.0 3.4. Conclusions.

The growth rate of a-Si:H as a function of the rf power level and degree of dilution of silane with hydrogen has been measured. The results are described in terms of shifts in quasi-chemical equilibrium. An infringement of the theoretical expected linear dependence of the growth rate from the rf power level is observed. This is due to a stronger recombination of the SiFU radicals producing SipHg at higher rf power levels and depletion of silane in the gas phase caused by the enhanced conversion of silane in

amorphous silicon at higher rf power densities. Hydrogen is able to etch a-Si:H films thus lowering the growth rate. This etch effect is also a third possibility for explaining the observed infringement. A common maximum is found for the growth rate near 295 mW/cm for various si lane-hydrogen mixtures. At still higher rf powers the growth rate starts to drop. This is ascribed to an enhanced influence of the hydrogen.

References.

[3.1]. N. Hata, A. Matsuda, K. Tanaka. " Spectroscopie diagnostics of plasma-chemical vapor deposition from si lane and germane", J. Appl. Phys., 61(8), 1987, 3055.

[3.2]. J. Perrin, Thesis, University Paris VII, Paris, 1983.

[3.3]. A.M. Antoine, B. Drevillion, P. Rocca i Cabarrocas, "In situ investigation of the growth of rf glow discharge deposited amorphous germanium and silicon films", J. Appl. Phys., 61(7), 1987, 2501.

[3.4]. R.A. Street, J.C. Knights, D.K. Biegelsen, "Luminescence studies of plasma-deposited hydrogenated silicon", Phys. Rev. B, 18(4), 1978,

1880.

[3.5]. J.E. Potts, E.M. Peterson, J.A. McMillian, "Effects of rf power and reactant gas pressure on plasma-deposited amorphous hydrogenated silicon", J. Appl. Phys., 52(11), 1981, 6865.

[3.6]. R.W. Collins, J.M. Cavese, "Effect of substrate temperature on the nucleation and growth of glow discharge hydrogenated amorphous silicon", Appl. Phys. Lett., 49(18), 1986, 1207.

[3.7]. R.G. Pyon, M. Aozasa, K. Ando, "Bias effect on the morphology and growth rate of glow discharge a-Si:H films in triode system", Jap. J. Appl. Phys., 25(7), 1986, 944.

[3.8]. R.C. Ross, J. Jaklik, Jr., "Plasma polymerization and deposition of amorphous hydrogenated silicon from rf and dc si lane plasmas", J. Appl. Phys. 55(10), 1984, 3785.

[3.9]. A. Gallagher, "Amorphous silicon deposition rates in diode and triode discharges", J. Appl. Phys., 60(4), 1986, 1369.

[3.10]. J.C. Knights, R.A. Lujan, M.P. Rosenblum, R.A. Street, D.K. Biegelsen, J.A. Reimer, "Effects of inert gas dilution of silane on plasma deposited a-Si:H films", Appl. Phys. Lett., 38(5), 1981, 331.

[3.11]. P.E. Vanier, F.J. Kampas, R.R. Cordermann, G. Rajeswaran, "A study of hydrogenated amorphous silicon deposited by rf glow discharge in silane-hydrogen mixtures", J. Appl. Phys., 56(6), 1984, 1812.

[3.12]. J. Shirafuji, S. Nagata, M. Kuwagaki, "Effect of hydrogen dilution of silane on opto-electronic properties in glow discharge hydrogenated silicon films", J. Appl. Phys., 58(9), 1985, 3661.

[3.13]. R.C. van Oort, M.J. Geerts, J.C. van den Heuvel, H.M. Wentinck, "The effect of silane dilution with hydrogen on the deposition of glow-discharge deposited silicon films", Proc. of the Seventh EC Photovoltaic Solar Energy Conference, Seville, 1986, 480-483.

[3.14]. V.I. Kuznetsov, R.C. van Oort, J.W. Metselaar, "Influence of the discharge power on PECVD amorphous silicon film properties", Proc. of the 1988 International Topical Conference on "Hydrogenated Amorphous Silicon Devices and Technology", New York, USA, 1988, 18.

[3.15]. V.I. Kuznetsov, R.C. van Oort, J.W. Metselaar, "Plasma deposition of hydrogenated amorphous silicon: effect of rf power", accepted for publication in the Journal of Applied Physics, to appear in January 1989.

[3.16]. B. Chapman in Glow-discharge Processes, J. Wiley, New York, 1980.

[3.17]. M.J. Kushner, "A kinetic study of the plasma etching process: A model for the etching of Si and Si0? in CnFn /H2 a n d CnFn/°2 P1 a s m a s" ' J- AP P]• P hys-> 53t4)>

1982, 2923.

[3.18]. S. Vepreck, Z. Iqbal, H.R. Ostwald, F.A. Sarrot, J.J. Wagner, "Parameters controlling the deposition of amorphous and microcrystalline silicon in Si/H plasmas", J. Phys., 10, 1981, C4-251.

[3.19]. M.J. Kushner, "On the balance between silylene and silyl radicals in rf glow discharges in silane: The effect on deposition rates of a-Si:H", J. Appl. Phys., 62(7), 1987, 2803.

[3.20]. A. Gallagher, "Neutral radical deposition from silane discharges", J. Appl. Phys., 63, 1988, 2406.

4. OPTICAL, ELECTRICAL AND STRUCTURAL PROPERTIES OF HYDROGENATED AMORPHOUS SILICON DEPOSITED FROM SILANE-HYDROGEN MIXTURES.

4.1. Introduction.

The optical, electrical and structural properties of a-Si:H films deposited from si lane-hydrogen mixtures are the subject of this chapter. A distinction has been made between three categories: electrical, optical and structural properties. Attention has been paid to all three categories because there is a close relationship between them. The electrical properties have been investigated by means of space-charge limited current measurements. The optical properties by means of transmission, and reflection measurements and photo-thermal deflection spectroscopy providing data for calculation of the Tauc optical gap, the refractive index, the value of the Urbach edge and the sub-bandgap absorption. The structural properties have been investigated by means of infrared spectroscopy and x-ray diffraction. The refractive index of a-Si:H is both a property of the film surface and of the bulk. The intensity of the reflection signal yields information about the surface roughness of the films. From the position of the reflection extrema the bulk refractive index is obtained. All other methods and properties mentioned provide information about the bulk of the material.

In Chapter 3 the growth rate of a-Si:H films as a function of rf power for various silane-hydrogen mixtures

was determined. In this way one is able to deposit films of equal thickness. This calibration is necessary to avoid possible thickness dependent film properties. Thickness-dependent film properties were reflected in the determined thickness dependence of the Tauc optical gap.

Material properties, such as the density of states (DOS) in the energy gap are of major importance. The DOS is extremely sensitive to the deposition conditions. A lower DOS means a better quality material with regard to the electrical properties. The states in the energy gap originate from tailing of band edges because of absence of long range order, from dangling bonds, defects and impurities in the material. The deposition conditions (rf power level, degree of silane dilution with hydrogen) have been systematically varied to study the effect of these parameters on the DOS as measured by the space-charge limited current method and the sub-bandgap absorption of intrinsic a-Si:H [4.1, 4.2]. Sub-bandgap absorption measurements were performed to establish whether a different optical measurement method confirmed the results of the electrical analyses [4.3]. To obtain information about the structure of the material infrared spectroscopy analyses were performed [4.4]. The results have been related to the electrical and optical characterization and were used to clarify the influence of silane dilution with hydrogen on the properties of a-Si:H.

4.2. Optical properties.

4.2.1. Tauc optical gap and refractive index.

The optical measurements used in Chapter 3 to obtain the growth rate of a-Si:H also provided the optical gap E , by

means of the Tauc plot, and the refractive index n. According to reference [4.5] the optical gap and the structural order of the films are correlated.

The influence of the rf power level and degree of dilution of the silane on the optical gap E of the films is given in Table 4.1, while the refractive index n of the films is presented in Figure 4.1.

_ c 4.00 a 8 6 ■ a 7 2 ■

ass

a3o

Figure 4.1 a: The refractive index of i-a-Si:H films grown under different process conditions. The same layers as used in Figure 3.1. Films A, B, C, and D are deposited using pure

o silane and an rf power of 15, 36, 62, and 194 iêl/cm respectively.

.D C 430 4.14 3.98 3.82 -3.66 • 3.50 0.60 i. . 0.67 i \.~ 0 . 7 4 0.81 wavelength (firn 0.88 0.95

Figure 4.1 b: The refractive index of i-a-Si:H films grown under different process conditions. The same layers as used

in Figure 3.1. Films A, B, C, and D are deposited using a gas mixture of 45 vol.% si lane and 55 vol.% hydrogen, and an

2

c E 4.06 <U ? 3.92 > O a 7 8 i -a 6 4

aso

V

V

Figure 4.1 c: The refractive index of i-a-Si:H films grown under different process conditions. The same layers as used in Figure 3.1. Films B, C, and E are deposited using a gas mixture of 33 vol.% siiane and 67 vol.% hydrogen, and an rf

2

For comparison: the refractive index as found in reference [4.5] for a-Si:H has a value of 3.5 and for microcrystalline silicon (uc-Si:H) a value of 2.9 at a wavelength of 0.9 /zm. It is clear that with increasing rf power n decreases with an accompanying blue shift in E , indicating a rise in the structural disorder. At lower silane concentrations the shifts are less pronounced or absent suggesting a positive influence of hydrogen admixture.

The data presented in Figures 3.1 and 4.1, in combination with Table 4.1, is useful in determining the direction of a serie of experiments based simply on the results of optical measurements. For example, in the case of using pure silane:

If one determines for deposited a-Si:H films a value of the optical gap of 1.73 eV and a refractive index of 3.4 at a wavelength of 0.9 /zm, then one must lower the rf power density to obtain a better quality a-Si:H.

Table 4.1 The Tauc optical gap (in eV ± 0.01 eV) of i-a-Si:H films deposited under different process conditions. The same films as used in Figure 3.1. The Tauc optical gaps given are measured using films with a thickness exceeding 0.6 \im.

rf (mW/cm2)--> 15 3 6 . 62 194 295 482

vol.% SiH4

100 1.65 1.70 1.74 1.73 1.74 1.70 55 1.68 1.70 1.70 1.71 1.71 1.71 33 1.68 1.69 1.70 1.70 1.71 1.71

4.2.2. Thickness dependence of the optical film properties.

It is generally recognized that thin films prepared by physical vapor deposition show a wide range of microstructures and film properties, both of which are highly dependent on the preparation conditions [4.6]. This has led to the development of the structured zone model (SZM), which relates the microstructure of a thin film to the film thickness, the pressure during deposition and the reduced temperature T/T . T is the actual film temperature during deposition and T is the melting point of the film, both in Kelvin. In this model, micro-structural development is controlled by shadowing effects (shielding the flow of gas phase particles by film surface features such as columns due to the linear motion of the gas phase particles.) in the range T/T <0.5 (zone 1 ) , surface diffusion in the range 0.5<T/T <0.8 (zone 2 ) , and bulk diffusion in the range T/T >0.8 (zone 3 ) . This classification appears to explain why the microstructure is independent of the particular vapor deposition method used to prepare the film. The SZM was developed to explain the relation between microstructure and preparation conditions in the case of crystalline metallic films. However there have been several recent reports of its application to both amorphous and crystalline ceramic films and amorphous semiconductor films [4.6]. The structure is not unique to the deposition technique but has been found in all vapor-deposited films, as well as electro-deposited films. The only common link between all these structurally anisotropic materials is that they were prepared under conditions of low mobility (T/T<0.5). For the specific case of a series of a-Si:H films prepared at a T/T =0.27 under low ion bombardment conditions, as a

' m ' function of film thickness, five distinct levels of physical