cvd diamond films nucleation and growth

CVD diamond films: nucleation and growth

S.-Tong Leea,*, Zhangda Linb, Xin Jiangc

Center Of Super-Diamond & Advanced Films and Department of Physics & Materials Science, City University of Hong

Kong, Hong Kong, China

State Key Laboratory of Surface Physics, Institute of Physics, Chinese Academy of Sciences, Beijing 100080, China

Fraunhofer-Institut fuÈr Schicht- und Oberflaechentechnik, Bienroder Weg 54E, D-38108 Braunschweig, Germany

Received 21 April 1999; accepted 10 June 1999

Abstract

In the last decade, we have seen rapid developments in metastable diamond synthesis by means of low-pressure chemical vapor deposition. Concurrently, a fast growing interest in diamond technology has emerged. Thisreview discusses the various low-pressure growth methods of diamond films. Particular attention is paid to recentadvances in the understanding of the mechanism of diamond nucleation and metastable growth. These advances arediscussed in connection with the advances in diamond heteroepitaxy, which raises hopes that single crystallinediamond films are not far beyond reach. Modern surface science techniques applied to diamond study have playedan essential role in these achievements and their contributions are discussed. # 1999 Elsevier Science S.A.All rights reserved.

1. Introduction

Diamond is considered an ideal material for many applications [1±6]. Its structure belongs to thespace group O7

h (F4, /d 32/m) with two atoms per primitive (Bravais) cell. The structure shown inFig. 1 can be viewed as two interpenetrating face centered cubic lattices shifted along the bodydiagonal by (1/4, 1/4, 1/4)a, where a is the dimension of the cubic (mineralogical) unit cell. Eachcarbon atom has a tetrahedral configuration consisting of sp3 hybrid atomic orbitals. The {1 1 1}crystallographic plane comprises 6-atom hexagonal rings arranged so that the adjacent atoms arealternately dislocated upward and downward from the plane. The stacking sequence in the h111idirections is ABC ABC ABC. . .. The lattice constant is 3.56 AÊ and the bond length is 1.54 AÊ .Natural diamond consists of 98.9% 12C and 1.1% 13C. The characteristic Raman spectroscopicsignals for diamond are 1332 cmÿ1 for 12C and 1284 cmÿ1 for 13C. Diamond has two isomers. Thefirst isomer is lonsdaleite found in meteorites. The structure of lonsdaleite is derived from diamondas shown in Fig. 2. The positioning of atoms in each plane is the same as that in the cubic structure.However, the planes are linked in a manner which results in a stacking sequence of AB AB AB. . ..Consequently, the atoms experience closer chemical bonding, with lattice constants in the a and cdirections of 2.52 and 4.12 AÊ respectively. The distance between adjacent atoms is 1.52 AÊ . Thecorresponding Raman peak is in the range of 1315±1325 cmÿ1. Another isomer is graphite, the mostcommon form of carbon. Each carbon atom has a sp2 atomic configuration and therefore, three in-plane sigma bonds. The remaining valence electron forms � bonds using a pz atomic orbital. Thus,

Materials Science and Engineering, 25 (1999) 123±154

0927-796X/99/$ ± see front matter # 1999 Elsevier Science S.A. All rights reserved.

PII: S 0 9 2 7 - 7 9 6 X ( 9 9 ) 0 0 0 0 3 - 0

*Corresponding author.

the trigonally bonded 6-carbon rings are situated in a flat plane instead of being alternately plaited asin diamond. The planes are layered in an ABABAB. . . sequence. The lattice constant in the basalplane between repeating layers is 6.707 AÊ , and the in-plane, nearest neighbor spacing is 1.42 AÊ . Thesignature Raman peak of the in-plane layers is 1580 cmÿ1.

In this review we will briefly discuss diamond properties, diamond applications, and diamondgrowth by chemical vapor deposition in Sections 2 and 3. Special attention is paid to recent advancesin the understanding of the mechanism of diamond nucleation, metastable growth, as well asheteroepitaxy in Sections 4, 5 and 6. Modern surface science techniques applied to diamond studyhas played an essential role in these achievements and their contributions will be discussed inSection 7.

Fig. 1. Face-centered cubic structure of the diamond crystal.

Fig. 2. Crystal structure of cubic and hexagonal diamonds. The difference in stacking sequence of (1 1 1) layer pairs in twostructures has been illustrated. Differences in the atomic arrangements are highlighted by the darkened bonds.

124 S.-T. Lee et al. / Materials Science and Engineering 25 (1999) 123±154

2. Properties and applications of diamond and diamond films

2.1. Properties of diamond

By virtue of its very strong chemical bonding, the structure of diamond leads to specialmechanical and elastic properties. The hardness, molar density, thermal conductivity, sound velocity,and elastic module of diamond are the highest of all known materials while its compressibility is thelowest of all materials (Ref. [4], pp. 8±9).

Diamond also has the largest Young's modulus among all materials. The dynamic frictioncoefficient of diamond is only 0.05, as low as that of teflon and the lowest among the materials ofinterest. Diamond exhibits the highest longitudinal sound velocity among all materials.

Diamond possesses the highest thermal conductivity ever known (Ref. [4], p. 4). Even so, itsthermal conductivity would increase five times more if diamond were to be fabricated usingisotopically pure carbon on account of decreased phonon scattering by differing isotopes [7,8].

The most important optical properties are the refraction index and optical absorption for a givenwavelength and temperature. Diamond has an appropriate refraction index and a small absorptioncoefficient of light from infrared to ultraviolet region [4].

The Hall mobility of holes in natural diamond is 1800 cm2/v.s, and that of electrons can reach2000 cm2/v.s. For synthetic homoepitaxial diamond, a hole mobility as high as 1400 cm2/v.s hasbeen obtained. For natural diamond, the hole and electron carrier drift velocities start to saturate at anelectric field strength of 104 V/cm. The saturation velocity is 107 cm/s for holes, and 2.0 � 107 cm/sfor electrons. Also, its electrical resistivity can reach 1015 cm.

The dielectric loss tangent is an important parameter for applications in microwave andmillimeter wave. Diamond possesses the lowest loss tangent among the compared materials (Ref.[4], p. 13±14).

The chemical properties of diamond have been reviewed in Ref. [9]. Diamond does not react tocommon acids even at elevated temperatures [3]. Treated by a hot chromic acid cleansing mixture ora mixture of sulfuric and nitric acids, graphite slowly oxidizes while diamond is chemically inert.However, diamond oxidizes (graphitizes) readily at high temperatures in an oxygen atmosphere andin air. Also, molten hydroxides, the salts of oxy-acids, and some metals (Fe, Ni, Co etc.) have acorrosive effect on diamond.

At temperatures above 870 K, diamond reacts with water vapor and CO2 [3]. The oxidation ofdiamond in potassium-containing liquid salts is twice as rapid as etching by sodium-containing salts[10]. Diamond may either chemically react with metals to form carbides, or dissolve in the metals.Metals such as tungsten, titanium, tantalum, and zirconium react with diamond to produce carbidesat high temperatures, while iron, cobalt, nickel, manganese, and chromium dissolve diamond.Because diamond dissolves in and/or reacts with iron or iron alloys (e.g. steels) at temperaturesabove 950 K, diamond tools are unsuitable for most machining operations on ferrous metals,including high speed and hardened steels.

2.2. Applications of diamond

By virtue of its excellent hardness and low coefficient of friction (in the range of 0.05±0.1),diamond can be used as cutting tools. Materials most pliable to mechanical deformation by diamondinclude aluminum, aluminum alloys, copper, copper alloys, chlorides, fluorides, polycarbonates,plastics, quartz, sapphire, NaCl, Si3N4, SiC, Ti, WC, ZnS, and ZnSe.

S.-T. Lee et al. / Materials Science and Engineering 25 (1999) 123±154 125

Diamond or diamond-like carbon is also used for coating magnetic disks to protect the headcrashes on magnetic disks, which require surface smoothness and hardness. In addition, fine grainpolycrystalline diamond films can be used as wire dies and water jet nozzles, since polycrystallineartificial diamond nozzles are isotropic in hardness and lighter in weight, the latter of which iscritical for most streamlined water cutting operations.

With a thermal conductivity of 20 W/cm/8C, diamond is unparalleled as a thermal conductor.For polycrystalline CVD diamond films, these figures are dependent upon the grain size. For a filmcolumnar in structure, the thermal conductivity can be lowered to 55% of the best value in the growthdirection and 25% of the best value in the lateral direction. With a high thermal conductivity,diamond is presumed to be the ideal heat exchange material (heat sink and heat spreader). Diamondhas been utilized as an electrically insulating thermal conductor for various electronics applications.Recently, high-power laser diodes have also been mounted on diamond in order to improve theperformance and increase the output power of the diodes.

Very large integrated circuit (VLSI) multiple chip module (MCM) compacts also use thickdiamond films as heat spreaders to increase the packaging density [11].

Diamond has the potential for both passive and active optical applications, although currentusage is only passive. Passive applications take advantage of its high thermal conductivity, corrosionresistance, and hardness, as well as its low absorption coefficient and small coefficient of friction.

The first diamond window was used for the IR emission sensor of the Venus explorer. Diamondwindows were also used in periscopes and in missiles.

Optical matching is another passive usage of diamond. Diamond has a refractive index of 2.4,which is lower than that of most semiconductors and higher than that of typical dielectrics. Diamondgenerally has a lower refractive index than materials from which infrared detectors are made,namely, silicon, germanium, group II±VI elements, and lead salts. Therefore, it is the preferredmaterial for coating applications. In addition, much progress has been achieved for longerwavelength detectors with higher refractive indices. The efficiency of silicon solar cells has beenaugmented by as much as 40% while that of germanium cells by up to 88% by diamond coatings[12].

Recently, it has been proposed that CVD polycrystalline diamond film can be used as a very fastoptical switch (� 60 ps), due to its low dielectric constant and high breakdown voltage [13,14].

Because of its high carrier mobility, breakdown field, saturation velocity, thermal conductivity,and wide band gap, diamond is considered an ideal material for electronic devices which function athigh temperatures, voltages, power-levels, frequencies, and radiation environments. For additionalinformation, readers may refer to some thorough references in this field [4,5,15±20]. The utilizationof synthetic crystals in photodetectors, light emitting diodes, nuclear radiation detectors, thermistors,varistors, and negative resistance devices has been documented by Bazhenov et al. [16]. Meanwhileseveral groups [18,20,21] have demonstrated basic field effect transistor (FET) device operation withhomoepitaxial diamond films and boron-doped layers on insulating single crystal diamondsubstrates. However, wide application of diamond solid-state devices require high quality films to bedeposited on more commonly available substrates.

Device fabrication comprises a number of process steps besides the growth of a high qualitycrystalline film. These steps include controlled doping of the film, selected area doping with p and ntype dopants, etching, formation of ohmic and rectifying contacts, deposition of dielectric films, andpassivation of the film surface. In the following we will review and discuss the state of the artmethods pertaining to these processes.

Heteroepitaxy of diamond has been attempted on a variety of single crystals including c-BN, b-SiC, Si, Ni, Pt, and Ir. Although single crystalline films have not yet been obtained, heteroepitaxial


growth of diamond on silicon has proven to be a very feasible possibility and is currently beingextensively investigated. This topic will be discussed in detail in Section 7.

Doping is an essential process in device technology. Boron is the dopant that causes p-typebehavior in diamond. Homoepitaxial and polycrystalline films can be doped with boron by chemicalvapor deposition of B2H6 gas or various solid sources [22±24] as well as by ion implantation [25±29]. While the problem of p-doping in diamond has been solved in principle, the problem of n-doping (N, P, Li, Na, K and Ru) in diamond is still open. As summarized by Popovici and Prelas intheir review article [30], the difficulty associated with n-doping is due to distortion of the latticeinduced by large n-dopants. The lattice distortion generates a relatively shallow acceptor level indiamond, which in turn compensates donor dopants. Thus a reasonable approach to activating n-typediamond is to search for a method which will not induce serious lattice distortion during doping.

Geis reported that the formation of ohmic contacts is possible on highly boron-doped diamond(about 1021 cmÿ3) via the deposition of single layers of several metals [19]. Sandhu [31] and Prins[32] found that ohmic contacts can be formed by select-area boron implantation at room temperatureusing 65 keV ions at a dose of 3 � 1016 cmÿ3. This high-dose ion implantation causes extensiveradiation damage which leads to graphitization of the surface layer during the subsequent annealingtreatment. The surface is then eliminated by a boiling solution of acids and contact is establishedwith a film of metal (Ag, Cu, or Au). At about 8008C, metallization occurs with an alloy composedof Ag, Cu and In, which provides strong mechanical bonding. A specific contact resistance of3.7 �10ÿ3 cm2 can be obtained using an as-deposited film of Au by employing a standard TLMmodel. Furthermore, a specific contact resistance with an order of magnitude of 10ÿ6 cm2 can beachieved by Au/Ti metallization and subsequent annealing at 8508C and 10ÿ6 Torr for 30 min.

Rectifying contacts effective at both room temperature and 4008C have been fabricated bydepositing heteroepitaxial Ni films on single crystal diamond [33]. At room temperature, Ni filmsdeposited at 5008C, for 625 mm diameter dot, are measured to have a reverse leakage current of 2 nAat a bias of 20 V. Also, the Ni films adhere well to the underlying diamond substrate. Another meansof producing excellent rectifying contacts on diamond is to use a composite film of co-sputtered Taand Si [34]. This has been demonstrated in the work by Shiomi et al. [21] where a rectifying contactwas positioned on an undoped diamond film, which was in turn placed on an already depositeddoped film. The result was a significant decrease in the reverse leakage current and an improvementin the breakdown voltage.

Due to the difficulty in obtaining n-type diamond, designs for bipolar transistors which utilizep±n junctions have not yet been realized. Instead, many schemes have been concentrated in metalsemiconductor field effect transistors (MESFET) or metal insulator field effect transistors(MISFET). Gildenflat et al. [35] and Zeisse et al. [36] have fabricated these FET devicesrespectively. However, the characteristics achieved thus far for this type of device are not goodenough for practical applications. The optimal designs and fabrication methods have yet to berealized, and the quality of the deposited films also needs to be improved.

3. CVD processes for diamond growth

3.1. Historical aspect of CVD diamond

The cyclic process developed by Eversole [37] was the first method to demonstrate CVDdiamond growth under low pressures. In the early 1970s, Angus et al. extended this work. They grewboron-doped diamond film on diamond grit (seed) [38]. Eversole's work was further expanded by


Derjaguin et al. who performed careful physical chemistry experiments [39]. In the cyclic pyrolysismethod, diamond was used as a substrate, and diamond growth occurred homoepitaxially. However,since cyclic, hydrocarbon pyrolysis had a very slow diamond deposition rate (�1 nm/h) and requireda diamond grit substrate, its application was unrealistic.

In 1982, Matsumoto et al. made a breakthrough in CVD diamond technology [40]. They usedhot filaments (�20008C) to directly activate hydrogen and hydrocarbon which were passed throughthe hot filament. The diamond film was then deposited onto a non-diamond substrate located 10 mmaway from the filament. Graphite was etched simultaneously by atomic hydrogen during depositionwhich rendered the cycling of deposition and etching unnecessary and therefore led to a highergrowth rate (�1 mm/h).

Since then, various activating methods for diamond CVD such as DC-plasma, RF-plasma,microwave plasma, electron cyclotron resonance-microwave plasma CVD (ECR-MPCVD), and theirmodifications were developed. The role of atomic hydrogen in diamond growth has gradually beenrecognized, and the growth rate approached the rate acceptable by industrial standards. In the late1980s, synthesis of diamond under low pressures attracted the involvement of many scientists andstimulated a `diamond fervor'. Currently, the DC-plasma jet diamond deposition method has earnedextensive attention in industry owing to its high growth rate. However, the apparatus of the DC-plasma jet is highly expensive.

Worthy of mention is the method of fluorocarbon pyrolysis. OH radicals, O2, O, F2, and F asgraphite etchants are even better than atomic hydrogen (refer to Fig. 3). Based on these results,Rudder et al. [41,42] predicted that the pyrolysis of fluorocarbons such as CF4 could produceepitaxial diamond growth. In their experiment, a mixture of CF4 and F2 diluted in He was blown ontoa diamond substrate heated to 8758C. The deposited film was verified to be diamond by Ramanspectroscopy, and the deposition of graphite was not detected. The pyrolysis process happened nearlyat thermodynamic equilibrium. However, a growth rate of only �0.6 mm/h was achieved. Thistechnique has the potential to be energetically more efficient than the CVD methods.

Fig. 3. Carbon atom removal probabilities for the attack of an isotropic polycrystalline graphite(� = 1.73 g/cm3) by O, O2,OH, F, F2 and H (Reprint with permission J. Phys. Chem. Copyright 1973, American Chemical Society). (Ref. [4] p. 152).


Beside CVD, physical vapor deposition (PVD) methods were also attempted in parallel [43,44]and were expected to deposit diamond at low temperatures. Recent work by S.T. Lee et al. [44]indeed has demonstrated that diamond nanocrystals in the matrix of amorphous carbon could beproduced by direct low-energy ion bombardment using a mixture of CH4/H2/Ar ions.

3.2. Hot filament-assisted CVD

Hot filament CVD method is the earliest method used for the growth of diamond under lowpressures, and is also the most popular method.

In 1982, Matsumoto et al. [40] exploited a refractory metal filament (such as W) and heated it toa temperature above 20008C, at which atomic hydrogen could be easily produced as H2 passed overthe hot filament. The simultaneous production of atomic hydrogen during hydrocarbon pyrolysiscould enhance the deposition of diamond. Diamond was deposited preferentially as graphiteformation was suppressed. As a result, the deposition rate of diamond increased to about 1 mm/h,which proved valuable for industrial manufacturing. A schematic diagram of hot filament CVDreactor was referred to [4], (p. 155). The simplicity and comparatively low capital and operating costof hot filament-assisted CVD have made the method popular in industry where it is imperative tominimize the price of synthetic diamond.

The company Diamonex has used HF-CVD to grow diamond films with diameters of up to30 cm. A wide variety of refractory materials have been used as filaments including W, Ta, and Re.Carbide-forming refractory metals must be first carburized before starting the deposition of diamondfilms. HF-CVD possesses the ability to adjust to a wide variety of carbon sources such as methane,propane, ethane, and other hydrocarbons. Even oxygen-containing hydrocarbons including acetone,ethanol, and methanol can be applied. The addition of oxygen-containing species may widen thetemperature range within which diamond deposition can take place.

In addition to the typical design of HF-CVD, some modifications have been developed. Themost popular is a combination of DC-plasma with HF-CVD where a bias voltage is applied to thesubstrate platen and filament (or accessory electrode) [45±47]. The application of a moderatelypositive voltage to the substrate platen and negative voltage to the filament (or accessory electrode)results in electron bombardment of substrate, which induces desorption of the surface hydrogen. Thelatter effect in turn increases the growth rate (up to about 10 mm/h). This technique is calledelectron-assisted HF-CVD. When the bias is strong enough to establish a stable plasma discharge,the decomposition of H2 and hydrocarbon is greatly enhanced, leading to a remarkable increase inthe growth rate (about 20 mm/h). Some laboratories even claim that the rate of diamond depositioncan exceed 30 mm/h on a 4 in. substrate. When the polarity of the bias is reversed, e.g. the substrateplaten is connected to a negative voltage, ion bombardment of the substrate surface will occur, andenhanced nucleation of diamond on the non-diamond substrate will result. Another modification is toreplace the single hot filament with multiple filaments or a filament net for uniform film depositionover large areas. For matching different requirements, further modifications are still underdevelopment. The disadvantage of HF-CVD is the contamination of the diamond films with elementsfrom the refractory metal filaments due to the evaporation of the hot filaments.

3.3. Microwave plasma-assisted CVD

In the early 1970s, scientists found that the concentration of atomic hydrogen could be increasedby use of a DC-plasma established by an electrical discharge [39,48]. The plasma became thereforeanother method to dissociate molecular hydrogen into atomic hydrogen and activate hydrocarbon


radicals into promoting diamond formation. Besides DC plasma two other kinds of plasma withdifferent frequency schemes are used. The excitation frequency for microwave plasma CVD istypically 2.45 GHz, while that for radio frequency (RF) plasma is 13.56 MHz. Microwave plasma isunique in that microwave frequency can oscillate electrons. High ionization fractions are generatedas electrons collide with gas atoms and molecules. Microwave plasma is often said to have `hot'electrons and `cool' ions and neutrals. A typical microwave reactor is referred to Ref. [4], p. 159.However, the drawback of this setup is the small substrate size. A substrate about 2±3 cm in diameteris the largest that can be introduced into a silica tube compatible with a WR284 waveguide. Aschematic representation of a microwave plasma system with a larger chamber reactor marketed byAstex is shown in Fig. 4. Microwaves enter into the reaction chamber from a proprietary antennawhich converts a rectangular WR284 microwave signal into a circular mode. The microwaveproceeds through a silica window into the plasma enhanced CVD process chamber. The size of theluminous plasma ball will increase with increasing microwave power. Diamond films have beengrown with the edge of the luminous plasma located about 2 cm higher than the substrate. Thesubstrate does not have to be in immediate contact with the luminous glow for diamond to grow viamicrowave plasma. Uniform diamond films with diameters of up to 4 in. can be deposited using thissystem.

3.4. RF plasma-assisted CVD

Radio frequency can be exploited to generate a plasma in two electrode configurations:capacitively coupled parallel plates and by induction. RF plasma-assisted CVD utilizes a frequencyof 13.56 MHz. The schematic of an inductively coupled RF plasma-enhanced CVD reactor and aparallel plate, capacitively-coupled RF reactor can be referred to Ref. [4], p. 164. RF plasma is goodin that it can be dispersed over more widespread areas than microwave plasmas. However, RFcapacitive plasma is limited in that the frequency of the plasma is optimal for sputtering, especially ifthe plasma contains argon. Because ion bombardment from the plasma results in serious damages todiamond, capacitively coupled RF plasma is not suitable for growth of high-quality diamond.

Fig. 4. Schematic of a microwave plasma-enhanced CVD reactor manufactured by ASTeX.


Polycrystalline diamond films have been grown by the RF induction method using depositionparameters similar to those of microwave plasma CVD [49]. Homoepitaxial diamond films have alsobeen deposited by RF induction PECVD [45].

3.5. DC plasma-assisted CVD

DC plasma is another method used to activate a gas source (typically a mixture of hydrogen andhydrocarbon) for diamond growth. The schematic of a DC plasma CVD reactor is shown in Ref. [4],p. 166. DC plasma assisted CVD possesses the ability to coat large areas, which are limited only bythe size of the electrodes and DC power supply.

As mentioned before, DC plasma has been combined with HF-CVD to obtain an increasedgrowth rate. Fujimori et al. [45] synthesized diamond films using a hybrid HF-CVD plus DC-plasmamethod. By applying ÿ120 V to tungsten filament heated to 22008C, he achieved a deposition ratethree times as fast while maintaining the integrity of the diamond films.

Another advance in DC plasma-assisted CVD is the DC plasma jet method. Researchers inJapan have devised a DC arc plasma-assisted CVD method which can deposit diamond films at ratesexceeding 20 mm/h. Kurihara et al. have designed a DC plasma jet facility termed DIA-JET [50,51].DIA-JET employs a gas injection nozzle consisting of a cathode rod surrounded by an anode tube.The typical growth rate is 80 mm/h. Norton company of USA has developed a very impressive DC-plasma jet system which utilizes a DC power of about 100 kW. Chinese scientists at BeijingUniversity of Science and Technology have also built a similar system. Because various DC arcmethods can synthesize high-quality diamond on non-diamond substrates at fast growth rates, theyprovide a marketable means for diamond film synthesis.

3.6. Electron cyclotron resonance microwave plasma-assisted CVD

As we discussed above, DC plasma, RF plasma, and microwave plasma all ionize anddecompose hydrogen and hydrocarbon species into hydrogen atoms and hydrocarbon radicals, andthus promote the formation of diamond. So, we can expect electron cyclotron resonance microwaveplasma-CVD (ECR-MP-CVD) to be a method especially capable of synthesizing diamond film sinceECR-MP generates high density plasma (greater than 1 � 1011 cmÿ3) which is favorable fordiamond growth. In fact, Hiraki et al. [52] used ECR-MP-CVD to fabricate diamond in 1990. Thegrowth temperature could be reduced to 5008C. Later, Yara et al. [53] and Mantei et al. [54]succeeded in diamond deposition using the ECR-MP-CVD technique. They obtained uniform filmsat substrate temperatures as low as 3008C. The experimental arrangement for a typical ECR plasmaCVD is shown in Fig. 5.

However, due to the extremely low pressure of the ECR process (10ÿ4±10ÿ2 Torr), diamondgrowth proceeds at a very low rate. Therefore, this method is only used in laboratories.

3.7. Combustion flame-assisted CVD

Combustion flame-assisted CVD method was first used by Hirose et al. [55,4]. The smolderingtip of a welder's torch oxidizes a mixture of C2H2 and O2 gas (ratio 1 : 1). Diamond crystals formwhere the tip of the bright interior section of the flame touches the substrate (temperature 800±10508C). Some advantages of the combustion method over the conventional CVD methods includethe simplicity and cost-effectiveness of the equipment, lack of power supply, high growth rate, andthe ability to deposit diamond over large areas and on curved substrate surfaces. The disadvantages


of this method are also obvious. Because deposition is difficult to control, the deposited diamondfilms are inhomogeneous both in microstructure and in composition. The welder's torch producesthermal gradients over the surface of substrate, thereby causing the substrate to warp or fractureduring the process of coating large substrate areas. Hirose et al. [55] and Cappelli et al. [56] havefurthered the work and achieved great progress in enhancing the area and film quality of diamondproduced by combustion flame-assisted CVD. In the future, this technique may be applied tofabricate diamond coatings used for tribological applications.

4. Nucleation mechanism of CVD diamond films

4.1. Practical significance of nucleation

Nucleation is the first and critical step of CVD diamond growth. The control of nucleation isessential for optimizing the diamond properties such as grain size, orientation, transparency,adhesion, and roughness that are necessary for targeted applications.

The investigation of diamond nucleation not only can lead to the controlled growth of diamondfilms suitable for various applications, but also it can provide insight into the mechanism of diamondgrowth. To date, the understanding of diamond nucleation is very limited. Carbon atoms can formdifferent types of chemical bonds via sp1, sp2, and sp3 hybridization. Diamond consists solely of sp3

bonds and is thermodynamically meta-stable compared to graphite, which is composed of sp2 bonds,under the experimental conditions used in CVD. It is an interesting and yet intriguing problem whymeta-stable diamond can be grown on diamond or non-diamond substrates under CVD conditions.

4.2. Methods of nucleation enhancement

During the early development of CVD diamond deposition, diamond single crystals were usedas substrates [57±60]. Later, diamond seeds were used [61,62]. Most early efforts were limited to

Fig. 5. Experimental arrangement for ECR plasma CVD of diamond.


homogeneous growth or homoepitaxy of diamond. In 1982 Matsumoto made a breakthrough ingrowing diamond on non-diamond substrates without using diamond seeds. Apart from having avery low nucleation density, a continuous film could not be formed. In 1987 Mitsuda et al. [63] foundthat scratching of the substrate surface with diamond powder could greatly enhance the nucleationdensity. Since then, substrate scratching has become the most common and powerful method forachieving nucleation that can form diamond with a high nucleation density and fine uniform grainsize. For silicon substrates, which have been studied intensively, a nucleation density of 107±108 cmÿ2 can be routinely obtained after scratching with diamond powder. In contrast, with non-scratched substrates, the nucleation density only reaches 104 cmÿ2. Besides diamond powder otherabrasive powders such as c-BN, TaC, SiC, and even iron can be used to scratch the substrate surfaceto enhance nucleation density. Nevertheless, diamond powder is regarded as the most effectiveamong the hard materials.

Later scientists have revealed that coating the substrate surface with graphite [64,65],amorphous carbon [64,66], diamond-like carbon [67±70], C60, and even mechanical oil [66] canenhance greatly nucleation density. Even so, these methods, including the scratching methodmentioned above, cannot lead to oriented nucleation or epitaxial growth on non-diamond substrates.

In 1991, Yugo et al. reported the bias-enhanced nucleation method by which they obtained ahigh density of nucleation on a mirror-polished substrate (without scratching) using the MWCVDsystem [71]. They applied a negative bias to the substrate during nucleation and obtained diamondnucleation with a density as high as 109±1010 cmÿ2 on Si. Subsequent developments of bias-enhanced nucleation by Jiang et al. [72,73] and Stoner et al. [74] have led to the heteroepitaxialgrowth of diamond on silicon and silicon carbide substrates, respectively.

For the popular HF-CVD method, since the gas reactants consisting of atomic hydrogen andhydrocarbon radicals are neutral species, a negative substrate bias cannot induce enhancednucleation. However, when a plasma is generated by the proper choice of biases, an enhancement ofdiamond nucleation similar to that for MPCVD can be achieved for the HFCVD process [46,47,75].Recently, other methods for enhancing diamond nucleation have been advocated. One method is thenucleation enhancement under very low gas pressures (0.1±1 Torr), while the other is by Si+

implantation into mirror-smooth Si substrate prior to the introduction of methane into depositionchamber. In the following various nucleation schemes will be discussed.

4.2.1. Mechanical abrasion of substrateIn 1987 Mitsuda et al. found that scratching of the surface substrate with diamond powder can

enhance greatly the nucleation density [63]. Since then, the effect of scratching has been extensivelystudied. It has been demonstrated that the scratching technique can be applied to most substratesused for diamond growth. Scratching pre-treatments with SiC [76], c-BN [77], Cu or stainless steel[78], ZrB2 [79] and Al2O3 [80] were also shown to enhance nucleation, although their effects werenot as strong as that of diamond powder. Heavy scratching or abrading of the Si substrate surfacewith diamond grit enhances the nucleation density by roughly three orders of magnitude comparedwith non-scratched Si (up to 107 cmÿ2). The nucleation density is proportional to the scratchingtime, and the morphology changes from large isolated crystals for short scratching times to smaller,high number density crystals with increasing scratching time [80,81]. The grit size of the diamondpowder used for scratching also influences the nucleation density; a 0.25 m grit is the most effective[82] for scratching by hand, and a 40±50 m powder is the best for scratching in an ultrasonic bathusing a grit suspension [83]. Why can scratching enhance diamond nucleation? One point of view isthat, during scratching with diamond, c-BN, or a-SiC powder, the residual powder or fragments areunavoidably left in the scratched groove and act as seeds for diamond growth. Although c-BN and a-


SiC are not diamond, their structures are close enough to that of diamond. Thus, diamond growseasily on them. Indeed, Iijima et al. [84] observed that fragments of diamond existed in the scratchedgrooves of Si, upon which the growth of diamond did occur.

Another opinion is that scratching with powder merely creates a change in the surfacemorphology, such as edges, steps, dislocations, and other surface defects. These kind of defects arelabeled chemical active sites, which prefer to adsorb diamond precursors together due to enhancedbonding at high energy intersecting surfaces with a high density of unsaturated bonds and lowcoordination numbers [80,85,86].

To identify this point experimentally, Dennig et al. [87] observed diamond nucleation on theridges lithographically etched on non-scratched Si by SEM. Singh [80] also reported enhancednucleation on Si etched by HF/HNO3. Recently, Si+ implantation on non-scratched Si has also beenfound to enhance nucleation density [88]. In these experiments, neither diamond seeds nor a carbonrich layer existed; only the surface morphology or surface structure was changed. However, otherattempts to enhance nucleation by creating etch pits on non-scratched Si using acids, H+ [89], orother reactive gases [90] proved unsuccessful. In the same way, attempts to generate large numbersof nucleating defect sites via implantation with surface saturation levels of C+ [91] or Ar+ [89]resulted in no nucleation enhancement on the unscratched substrates. The arrays of holes of 0.1±0.3 mm in diameter created by a focused Ga+ ion beam on a Si substrate resulted only in thedeposition of non-diamond carbon in the depressions [92]. The reason for the discrepancy amongthese experiments is still unclear. To resolve the problem, the substrate surface must be characterizedthoroughly since nucleation is a surface phenomenon. Results obtained on ill-defined surfacesinvariably lead to uncertainty and controversy.

4.2.2. Biased-enhanced nucleation in microwave plasma CVD and HFCVD

It is difficult for diamond to nucleate on mirror-polished Si and silicon carbide because of theirsurface free energies and lattice constants are much different from those of diamond. In 1991, Yugoet al. [71] obtained diamond nucleation with a density of about 109±1010 cmÿ2 on a mirror-polishedSi substrate by applying a negative substrate bias voltage in a MPCVD system. With this technique,the highest nucleation density on mirror-polished Si has reached the level of 1010±1011 cmÿ2 [93].The nucleation mechanism has been widely studied and different models have been proposed. Yugoet al. [94] and Gerber et al. [95] suggested a shallow ion implantation model in which the sp3 bondedcarbon clusters, formed by low energy ion implantation, function as the nucleation precursors. Thenegative bias causes the positively charged ions in the growth chamber to accelerate towards andbombard the substrate surface, thereby removing the contamination and facilitating cluster formationon the surface. These events in turn advance diamond nucleation. Stoner et al. [96] on the other handindicated the critical process should be the change in plasma chemistry, such as the increase in theconcentration of atomic hydrogen caused by substrate biasing and the formation of a carbide surfacelayer. Jiang et al. [97,98] found that the overall temporal evolution of the nucleation densitycorresponds well with a surface kinetic model involving immobile active nucleation sites, germs, andnuclei. They also suggested that, in addition to surface defects (point defects, steps and sp3-bondedcarbon clusters) serving as the nucleation sites, the enhanced surface diffusion and stickingprobability of carbon on silicon due to ion bombardment should be the decisive factors. Theenhancement of the surface diffusion of carbon species has been identified by investigation of thedistribution of the first nearest-neighbor distances [97].

Stubhan et al. [75] and Lin et al. [46,47] showed that, in HF-CVD system, diamond nucleationenhancement can also be realized when a negative bias is applied to the substrate. The highest


nucleation density can reach 109±1010 cmÿ2 on mirror-polished Si, which is similar to the resultsobtained using the MPCVD system.

To summarize the results obtained for the biased-enhanced nucleation, the decisive role of ionbombardment for the nucleation of diamond has been clearly demonstrated. The nucleation sequencecan be given as following: (1) Formation of nucleation sites, (2) Formation of carbon clusters due tothe enhanced surface diffusion, and (3) Formation of stable diamond nuclei.

4.2.3. Diamond nucleation at low gas pressureRecently, diamond nucleation has been conducted under very low pressures (0.1±1 Torr) and a

high density of diamond nucleation (109±1011 cmÿ2) has been achieved on mirror-polished Sisubstrates using either HF-CVD or ECR microwave plasma CVD [99,100] without applying surfacescratching or a substrate bias. Raman spectroscopy shows both diamond and graphite lines in thenucleated sample. Similar results have also been obtained on Ti substrates.

Using this method, diamond grains with a density >1010 cmÿ2 can be achieved, which is morethan 2 orders of magnitude higher than the highest density (107±108 cmÿ2) that can be obtained onscratched substrates under conventional pressure (10±50 Torr). This value matches the highestreported level to date, which was obtained using a negative substrate bias in a MPCVD system.

4.2.4. Ion implantation-enhanced nucleationIt has been demonstrated that in order to obtain a high nucleation density of diamond by CVD

methods, the substrate surface must be treated so that: (1) surface adsorption sites can be created forthe hydrocarbon radicals, and (2) the distributed region of the adsorption sites is large enough (largerthan the critical size for nucleation) for continuing growth of the diamond nuclei.

To modify the surface structure of the Si wafer, ion implantation can also be employed [88]. Si+

ions are implanted into a mirror-polished Si wafer. The implantation of Si+ ions changes only thesurface structure, and not the composition of Si. Therefore one can distinguish the surface structureeffect from others. After treatment with a Si+ energy of 25 keV (lower energy would be better) andan implantation dose of 2 � 1017 cmÿ2 diamond can easily nucleate and grow on a Si wafer andcontinuous diamond film can be synthesized.

Si+ implantation-enhanced nucleation is assumed to create nano-scale surface defects on the Sisubstrate. These defects serve as the active sites for the adsorption of hydrocarbon radicals necessaryfor initial diamond nucleation. A similar effect can also be found from the growth of diamond onporous silicon [101], the surface of which has a rich nano-scale microstructure.

5. Growth mechanism of CVD diamond films

5.1. Overview

The mechanism of CVD diamond growth has attracted increasing attention in recent years,mainly due to the fact that further technological advancement requires a more detailed understandingof the fundamental phenomena responsible for diamond synthesis. Questions such as how togrow diamond film more efficiently and economically, how to minimize the density of defects in thefilms, and which sources are most effective, all require a thorough understanding of the growthmechanism.

The first attempt to account for diamond {1 1 1} growth on the atomic scale was given byTsuda, Nakajima, and Oikawa [102,103], who proposed that diamond growth involved CH3

+ cations


or a positively charged surface. However, this mechanism is incompatible with HF-CVD, since CH3+

are scarce and the substrate surface is uncharged under HF-CVD conditions.Later, Chu et al. [104] proposed that the methyl radical is the dominant growth precursor under

HF-CVD conditions for all {1 1 1}, {1 0 0}, and {1 1 0} facets. They utilized isotope13 C todistinguish the growth precursor of diamond, and concluded that the methyl radical is the mainprecursor in diamond deposition. However, in a high temperature environment, CH4 and C2H2 willdecompose into various products and one cannot distinguish from which original source the productscome. Martin and coworkers [105,106] showed that methyl or methane is much more effective thanacetylene for growing diamond films.

Harris [107] proposed a growth mechanism involving only neutral CH3 and hydrogen atomsusing a nine-carbon model compound [bicyclononane (BCN)]. For the growth of diamond {1 0 0}surfaces, which is terminated by CH2 radicals, the intermolecular H±H spacing is only about 0.77 AÊ ,nearly the same as that in the H2 molecule (0.74 AÊ ). Since the interaction is nonbonding, very strongsteric repulsion between neighboring hydrogen atoms is expected. Such a repulsion will greatlyaffect the growth of the diamond {1 0 0} surface. Similar situations hold true for some other lowindex surfaces as well.

Frenklach [108±111] and co-workers proposed that acetylene, which is present in greaterquantities than CH3 under typical HFCVD conditions [112±114], is the primary growth precursor ondiamond {1 1 1} and other low index surfaces.

Recently, nanocrystalline diamond attracted people's attention due to its applications andfundamental significance. Studies on this kind of diamond revealed that they were constructed via C2

dimer [115,116] but not CH3 or C2H2.The answers as to whether the aforementioned models are correct and to what extent they

concur with the experimental findings are unclear due to experimental difficulties in studying thedynamic process of growth in-situ and on the atomic scale. As diamond nucleation and growth areboth surface phenomena, surface science techniques are suitable and powerful for investigating thegrowth process and for understanding its mechanism. Among the surface techniques, high-resolutionelectron energy loss spectroscopy (HREELS) is used to study the vibration of atoms or molecules toprovide information on the species, configurations, and adsorption sites at the surface. In the processof diamond growth, the precursors are always hydrocarbon radicals or modifications of them, whichadsorb on the growing surface. Therefore, HREELS is well suited for the study of the types ofprecursors attached on the surface and their evolution during nucleation and growth. Later we shalldiscuss recent mechanistic studies of diamond growth which have utilized various surface techniquesincluding HREELS.

5.2. Growth mechanism of diamond (1 1 1)

Many efforts have been devoted to deciphering the growth mechanism of diamond (1 1 1), andvarious growth models have been presented. To judge which model is correct requires experimentalsupport. To this end, HREELS measurements were conducted in-situ on as-grown homoepitaxialdiamond (1 1 1) surfaces and highly oriented (1 1 1) diamond films grown on Si (0 0 1) [117]. Fig. 6shows the HREELS spectra for different growth temperatures. These spectra are very simple. Theloss peak at 365 meV corresponds to the C±H stretching vibration, 155 meV to the C±H bendingvibration, 110 meV to the C±C stretching vibration, and 310 and 460 meV to the first and secondovertones of the C±H bending vibration. Judging from the vibration modes and comparing with thecharacteristic frequencies of the molecular subgroups of CHx in the handbook [118,119], the results


show that the surface is terminated not by CH3, but by CH radicals. Thus, diamond growth on the(1 1 1) surface proceeds in a two-by-two layers mode.

Notably, the C±H stretching vibration is perpendicular to the (1 1 1) surface, whereas the C±Hbending vibration is parallel to the (1 1 1) surface. According to the selection rule [120], thevibration parallel to the (1 1 1) surface cannot be activated and thus should be absent in the HREELSspectra. However, in Fig. 6, C±H bending vibration loss peak at 155 meV is indeed present. In orderto reconcile this apparent contradiction, the H on the diamond (1 1 1) surface was replaced by itsisotope D [117], and two types of adsorption sites were found to exist; one is on the (1 1 1) surface,and the other on another facet. If we speculate that the growing diamond (1 1 1) surface consists of(1 1 1) faces and (1 1 0) steps, it is not difficult to understand that among the two kinds of sitesmentioned above, one is located on the (1 1 1) faces and the other is on the (1 1 0) faces stepsbecause the bending vibration of C±H on the (1 1 0) faces steps is perpendicular to the (1 1 1)surface. Therefore, the C±H bending vibration should be active. As Frenklach pointed out, growth onthe (1 1 1) surface is carried out in two stages. First is kernel formation and second is the propagationstage. In the kernel formation stage, the appearance of island and the (1 1 0) faces steps are possible.Thus, the existence of the C±H bending vibration is understandable.

The second stage of the Frenklach model, which details the connection of C2H2 in a two- layersby two-layers manner, is supported by experimental results [117]. In addition, it has been upheldexperimentally that during kernel formation, CH3 is the possible precursor [121].

In summary, growth on the (1 1 1) surface has been proposed to be completed in two stages. Inthe first stage, the activation of a surface carbon by H abstraction, adsorption, and catenation of CH3

at the active site results in kernel formation. In the second stage, as predicted by Frenklach, the(1 1 1) surface grows along the (0 1 1) direction with acetylene only.

5.3. Growth mechanism on (1 0 0) surface

There also is a lot of debate pertaining to the growth mechanism of diamond on the (1 0 0)surface. Harris [107] proposed a growth model involving only neutral CH3 and hydrogen atomsusing a nine-carbon model compound [bicyclononane (BCN)], and the predicted growth rate on the(1 0 0) surface agreed well with experiment. However, the steric repulsion for the H±H site on BCNis different from that on the diamond (1 0 0) surface. On the diamond (1 0 0) surface terminated byCH2 radicals, a very strong steric repulsion should exist between the neighboring hydrogen atoms.

Fig. 6. High-resolution electron-energy-loss spectra of grown diamond (1 1 1) facets with 0.2% CH4 remote feed and at (a)8008C, (b) 9008C, and (c) 10008C.


Hamza et al. found a 1 � 1 LEED pattern on the diamond (1 0 0) surface at temperatures from 500 to750 K [122]. This was attributed to the saturation of the dangling bonds of a surface carbon atom bytwo hydrogen atoms, and thus surface reconstruction did not occur. This surface was a nominallydihydrogenated surface. Upon heating to over 1300 K, the surface structure showed a 2 � 1reconstruction due to the desorption of one H atom from a surface carbon atom. This surface was anominally monohydrogenated surface with an elongated C±C dimer bond [123]. A 2 � 1 reconstructionof the (1 0 0) diamond surface grown at 10008C was also observed by scanning tunneling microscopy(STM) and atomic force microscopy (AFM), respectively [124,125]. Using HREELS, the diamond(1 0 0) surface grown at a temperature of 8008C was investigated [126]. The spectra are shown in Fig. 7.Intensity losses occur at 156, 180, and 372 meV, with three smaller losses at 110, 310, and 530 meV.Compared with the characteristic frequencies of the molecular subgroups CHx [118,119], the spectrum isconsistent with that of the CH2 radical, which has its stretching vibration at 370 meV, scissors vibration at179 meV, wagging vibration at 157 meV, twisting vibration at 150 meV, and rocking vibration at108 meV. The 372 meV is assigned to CH2 stretching, 180 meV to scissors, 110 meV to rocking, and156 meV to the overlapping of wagging and twisting. The 310 meV is the overtone of the loss at156 meV, and the loss at 530 meV is the combination of CH2 wagging and stretching. The film grown atthis temperature exhibits good crystallinity.

If the growth temperature is increased to 10008C, the film shows a `cauliflower'-likemorphology, and the Raman spectrum exhibits a broad peak at 1580 cmÿ1, which is characteristic ofgraphite, while its HREELS is similar to that at 8008C. However, a prominent peak appears at140 meV. This peak corresponds to the bending vibration of the monohydrogenated dimer. Due tothe appearance of the monohydrogenated surface, hydrogen atoms that are bonded to the (1 0 0)surface become further separated. As a result, it releases the strong steric repulsion between the Hatoms. Nevertheless, the C±H bond of the monohydrogenated dimer possesses to some extent � bondcharacter. As the hydrocarbon radicals attach to the C±H bond, they also take in some � bondcharacter as well. Their tendency to form graphite is why we observe a broad peak at 1580 cmÿ1 inthe Raman spectrum. Noteworthy is that if the (1 0 0) surface of the as-grown sample is exposed toatomic hydrogen (no hydrocarbon involved), the 140 meV loss peak also appears in the HREELSspectrum. This can be explained by the abstraction of one of the two hydrogen atoms bonded to asurface carbon by gas phase atomic hydrogen.

If the abstraction is strong enough, monohydrogenated dimers appear in some local regionswhere the surface is similar to the growth surface at 10008C. If the amount of atomic hydrogen is

Fig. 7. High-resolution electron-energy-loss spectra of (a) diamond (1 0 0) facets grown at 8008C and 1.0% CH4, (b)(1 0 0) facets grown at 10008C, (c) the sample of (a) dosed with atomic hydrogen, and (d) dosed with oxygen.


small, or if the abstraction proceeds at a moderate temperature (about 8008C), or if hydrocarbon isinvolved in the gas source, the abstraction of H will not be strong. Thus, dihydrogenated carbonatoms remain in the neighborhood of C with one abstracted H, and monohydrogenated dimers cannotbe formed. The carbon atom keeps then only one H and one vacant site, which may bond to thehydrocarbon radical. As a result, diamond growth proceeds steadily.

The quantity of carbon with one hydrogen is determined by the growth temperature, the amountof atomic hydrogen near the surface, and the concentration of activated hydrocarbon. At growthtemperatures around 10008C, the (1 0 0) surface consists of monohydrogenated H±C±C±H dimers aswell. For even higher temperatures, large amounts of hydrogen desorb from the surface and thesurface carbons become C±C dimers.

The growth rate for the CH2-terminated surface is determined by the quantity of CH3 in the gasphase. It has been shown that atomic hydrogen can easily abstract one hydrogen atom from CH3,therefore, CH2 radicals become the precursor of (1 0 0) diamond growth [127].

5.4. Appearance regularity of the diamond facets

The control of the appearance of diamond grain facets becomes significant not only in practicalusage, but also in the testing of established diamond growth mechanisms. The morphology of CVDdiamond films is correlated with the growth parameters. Systematic studies on the relationshipbetween the appearances of the (1 1 1) and (1 0 0) facets and the growth parameters such as the ratioof CH4/H2, O2 content, and distance from the hot filament to the substrate have been performed[128,129]. A satisfactory explanation has however not yet been presented.

It is well known that, for a kinetically controlled growth system, the crystal morphology isdetermined by the appearance of facets which have the slowest growth rate in their normal directionand by the corresponding relative growth rates [130]. Because the (1 1 0) surface is a S (stepped)face and encounters no repulsion between adjacent hydrogen atoms, it should have the highestgrowth rate and appear as diamond facets. In fact, the growth rate of the (1 1 0) surface via the CVDmethod is the highest among the low-index surfaces, and either methyl-radical- or acetylene-species-based mechanisms can be postulated. Thus in most cases, the surface of CVD diamond crystalsappear with {1 0 0} faces and {1 1 1} faces. According to the established growth model, the growthrate of the (1 0 0) face depends on the concentration of CH2 or CH3 while the growth rate of the(1 1 1) facet relies on both CH3 and C2H2 for kernel formation and growth.

If the CH3 concentration near the substrate is much higher than the C2H2 concentration, the{1 0 0} growth rate will be high, and hence the {1 0 0} faces will be absent. The crystal appears ashaving just {1 1 1} faces. In contrast, if the C2H2 concentration near the growing surface isdominant, the {1 1 1} face grows fast [131], which consequently results in the appearance of {1 0 0}facets. In most cases the concentrations of CH3 and C2H2 are comparable, so both {1 0 0) and{1 1 1} facets are present. Harris and Weiner [132,133] have measured the dependence of the CH3

and C2H2 concentrations upon the CH4/H2 and O2/H2 ratios, as well as upon the spacing between thehot filament and substrate. Based on their experimental results, Sun et al. [126] explained theregularity of the appearance of the diamond facets.

5.5. The role of atomic hydrogen in CVD diamond growth

In the chemical vapor deposition of diamond crystals, the most crucial aspect is that thehydrocarbon gas must be diluted in hydrogen to as low as about 1% and the hydrogen must be


dissociated into atomic hydrogen. The role of atomic hydrogen for CVD diamond growth has beenclearly demonstrated in the HF-CVD method [40].

It has been recognized that atomic hydrogen plays a few specific roles. First, atomic hydrogenreacts with, or etches graphite about 20±30 times faster than diamond, so graphite and other non-diamond phases can be removed rapidly from the substrate and only clusters with diamond structureremain and continue to grow [126]. Second, atomic hydrogen stabilizes the diamond surface andmaintains the sp3 hybridization configuration [134]. Third, atomic hydrogen converts hydrocarbonsinto radicals, a necessary precursor for diamond formation. Fourth, atomic hydrogen abstractshydrogen from the hydrocarbons attached on the surface [135] and thus creates active sites foradsorption of the diamond precursor. However, too much atomic hydrogen causes unnecessarilystrong abstraction. The formation of monohydrogenated dimers will increase and the graphite phasewill readily appear, which results in the deterioration of diamond film quality.

5.6. The role of ionic hydrogen in diamond growth

It is generally accepted that diamond growth is a combined process of deposition and carbonetching which take place concurrently. The growth of diamond occurs if the deposition rate is largerthan the etching rate. During CVD growth of diamond films, both atomic H and H+ ions in theplasma cause etching, and H+ ions etch even faster than atomic H [136,137]. Recently it was foundthat a [0 0 1]-textured top layer can be prepared on polycrystalline or [1 1 1]-textured diamond filmsby the application of a negative substrate bias potential during diamond growth [137±139]. A noveletching effect of hydrogen ions on the growth of diamond films was observed and confirmed to playa dominant role for the [0 0 1]-textured growth. The H+ ion bombardment was performed byapplying a negative substrate bias during a microwave plasma CVD process, using only hydrogen asthe reactant gas. It was discovered that the etching efficiency of H+ ions on non-[0 0 1] orientedgrains is higher than that on grains with their (0 0 1) faces parallel to the substrate. Lateral growth ofthe (0 0 1) faces can occur during the bombardment process. As a result, the size of the (0 0 1) facesincreases after H+ etching while grains oriented in other directions are etched off. This effectprovides a way to improve the orientation degree of [0 0 1] oriented diamond films and may behelpful for obtaining very thin [0 0 1] oriented diamond films.

5.7. The role of atomic oxygen in diamond growth

The addition of oxygen to the reaction gases, as either CO, O2 or alcohol, has a beneficial effecton the growth rate and quality of the CVD diamond films and allows for diamond growth to occur atlow temperatures [140]. It was found [141] that although hydrocarbon species are somewhat reducedby oxygen addition, oxygen has only a relatively small effect on the mole fractions of radical speciessuch as H and CH3. Also, OH is formed at concentrations sufficient for the removal of non-diamondcarbon at a rate comparable to that of diamond growth. It was found that [142], at low temperatures(<8008C), the addition of oxygen not only enhances the growth rate but also extends the region ofdiamond formation. At temperatures lower than 5008C, oxygen was explained to have strongerpreferential etching behaviors than hydrogen. Another possible explanation about the role of atomicoxygen is the effective abstraction of surface hydrogen [143], which balances strongly theabstraction of surface hydrogen. These occurrences are helpful for diamond growth. On the otherhand, the addition of too much oxygen will cause too strong of a hydrogen abstraction and evenoxidation of the surface, which in turn will deteriorate the diamond film.


6. Heteroepitaxy of diamond

6.1. A great challenge

As mentioned in Section 2.2 pertaining to the electronic properties of diamond, diamond is animportant material for high temperature, high speed, high power, and high radiation-tolerantelectronic devices. However, to manufacture such devices, single crystals or epitaxial mono-crystalline films are required. Diamond can be grown epitaxially on either natural or high-pressure-synthesized diamond [144±152]. However, natural diamond is rare and expensive, while high-pressure-synthesized diamond crystals have only been obtained in small sizes (<0.5 mm). Thus, itappears the only practical way to grow diamond films is heteroepitaxially.

To date, c-BN is the best substrate on which diamond heteroepitaxy can be easily achieved[153±158], because of its close similarity to diamond in lattice parameter (1.3% misfit) and surfaceenergy (�4. J/m2 for the (1 1 1) surface of c-BN [159], 6 J/m2 for low-index planes of diamond[160]). Unfortunately, large c-BN single crystals (>2 mm) cannot be grown at present, althoughpolycrystalline c-BN films can be grown with an extremely small grain size (several nanometers).The second best candidate for diamond heteroepitaxy is b-SiC. In spite of a large lattice misfit ofabout 22%, Stoner and Glass [74] obtained highly oriented diamond crystallites using the bias-enhanced nucleation method. They obtained diamond crystallites with the following crystallographicrelationships: diamond (1 0 0)//b-SiC (1 0 0), diamond [011]//b-SiC [011].

In 1992, Jiang and Klages reported [72,73] that [0 0 1]-oriented diamond films can beepitaxially grown on the (0 0 1) plane of single crystal silicon, also by applying a negative electricalpotential to the substrate without the intentional formation of an intermediate layer. They found byhigh-resolution electron microscopy that diamond (0 0 1) grew directly on silicon (0 0 1) [161,162](Fig. 8). They stated that b-SiC is not a necessary interfacial layer required for diamondheteroepitaxy on Si. Later, Lin et al. [163±166] and Stuhban at al. [75] also realized heteroepitaxialgrowth of diamond (0 0 1) on silicon (0 0 1) in a HF-CVD system.

The structural quality of diamond films has been gradually improved during the recent years. Ithas been possible to reduce the orientation deviation on Si (1 0 0) substrates from a best value ofabout 98 in 1992 [167], as determined from the full width at half maximum (FWHM) of X-ray

Fig. 8. High resolution lattice image along the [1 1 0] direction of the diamond±silicon interface (Ref. [161]).


rocking curves, to a best value of about 28 today [168]. The misorientation of diamond crystallitegrown on b-SiC (0 0 1) substrates has also been strongly reduced. Diamond films with FWHMvalues smaller than 18 have been prepared on b-SiC (0 0 1) [169]. The application of a two-stepprocess (a [0 0 1] fast growth and a [1 1 1] fast growth) produced diamond particles misoriented by 2to 38 at the initial stage of deposition. However, as the diamond deposition proceeded, themisoriented particles disappeared and eventually smooth diamond films were formed.

Heteroepitaxial growth of diamond was also attempted on other types of substrates. Using HF-CVD, oriented diamond crystallites were deposited on a BeO single crystal [170]. Diamond filmswere also grown heteroepitaxially on Ni [171], Pt [172] and Co [173] substrates by treating thesubstrate surface with diamond powder. However, due to the complicated processing or the toxicityof the substrate material (BeO), it is questionable whether these approaches can have practicalapplications.

Recently, superior-quality heteroepitaxial growth of diamond on iridium layers has beenreported. The Ir layers have been deposited on the (0 0 1) cleavage planes [174±176] of MgO or onthe mechanically polished (1 0 0) planes of SrTiO3 [177]. Iridium has a lattice constant of 3.840 AÊ ,which is close to that of diamond 3.567 AÊ , and it (not a catalyst) does not form carbide under theconditions for diamond deposition. Thus Ir would seem to be an excellent substrate for diamondheteroepitaxy. By using bias-enhanced nucleation and subsequent growth by a MPCVD process,highly oriented diamond crystallites have been grown on Ir. The size of the crystallites obtained isaround 1 mm, while the XRD polar and azimuthal spread for the crystal orientation are less than 18.Therefore, this approach may offer a promising route for realization of large-area diamondheteroepitaxy.

In terms of practicality, the realization of diamond heteroepitaxy directly on silicon isparticularly attractive, because Si wafers are easily available and extensively used in electronicindustry. The epitaxial growth of diamond on Si is highly desirable, particularly in view of itsconvenient integration of diamond electronics with Si technology. Therefore the realization ofdiamond heteroepitaxy on Si has become a great challenge of immense technological and scientificinterest. In the following we will concentrate on this issue.

6.2. Oriented heterogeneous nucleation of diamond

Generally, diamond heteroepitaxy on silicon proceeds in two stages, namely, nucleation andgrowth. It is well known that it is difficult for diamond to nucleate on mirror-polished siliconsubstrates. Under conventional nucleation without substrate pretreatment, the nucleation density isonly about 104 cmÿ2, and only the growth of isolated individual diamond crystals is obtained.Because substrate scratching causes serious damage to the arrangement and periodic structures of thesurface atoms, epitaxial growth cannot occur via this technique. Bias-enhanced nucleation wasapplied to a mirror-polished Si substrate in a MPCVD system to avoid the substrate scratching withdiamond powder. Fig. 9 shows a schematic diagram of the process parameters. Prior to the bias-enhanced nucleation, in-situ hydrogen plasma etching was performed in order to remove the nativesurface oxide layer. By carefully controlling the nucleation process, diamond nucleation with morethan 90% [0 0 1]-oriented nuclei was achieved [98]. Later, high-density nucleation and orienteddiamond films were also achieved by applying a negative bias to the substrate in a HF-CVD system[75,163].

The investigation of the nucleation process reveals a narrow parameter window for epitaxialnucleation. The crucial parameters are the substrate temperature, methane concentration, appliedbias voltage to the substrate during nucleation, and the nucleation time. A critical bias voltage exists,


although experimental results show different values for various reaction chambers. Typically, thecritical bias voltage for a process pressure of 20±40 mbar is approximately ÿ80 to ÿ120 V (Fig. 10)[178].

Atomic force microscopy (AFM) and reflection-high-energy-electron diffraction (RHEED)were employed to investigate the early stage of diamond nucleation [179,97]. The RHEED patterns(Fig. 11) obtained for different nucleation times contained spots due to diffraction by diamond afterseveral minutes of deposition, which suggested a certain period of time is necessary for the initiationof nucleation. It is demonstrated that the diamond nuclei formed at the start of the nucleation areoriented. As the nucleation time increases, the nuclei increases dramatically in density to5 � 1010 cmÿ2 (Fig. 12) [97] and are also randomly oriented. These results confirm the importanceof controlling the nucleation process for epitaxial growth [98]. Recently, Schreck et al. studied this

Fig. 9. Schematic diagram of the microwave plasma CVD process for preparing heteroepitaxial diamond films (Ref. [98]).

Fig. 10. Diamond nucleation density as a function of substrate bias voltage (Ref. [178]).


processing space for heteroepitaxial nucleation of diamond on Si (0 0 1) using the bias process by X-ray diffraction texture measurements [180,181]. The temporal development of the azimuthal poledensity distribution is shown in Fig. 13. It can be found that the bias time lies within a distinct time

Fig. 11. RHEED patterns obtained for different nucleation time (a) 7.5 min, (b) 10 min, (c) 12.5 min, (d) 15 min (Ref.[97]).

Fig. 12. Nucleation density (islands/cm2) vs. deposition time. In the abscissa an induction time of 6.5 min is subtracted.The curve is obtained by computer modeling. The open and full circles in the plots represent data calculated from crystalsize distribution and obtained by direct particle counting, respectively (Ref. [97]).


interval. The width of the time window and the bias time for optimal crystal alignment decreasesharply as the absolute value of the bias voltage decreases (Fig. 14).

A possible mechanism of the loss of diamond epitaxy was suggested by Jiang et al. [182] andSchreck et al. [180] after investigating diamond growth under bias conditions. The authors reportedon slightly misoriented crystallites grown homoepitaxially on {1 0 0} facets under bias conditions.According to their results, the crystal misorientation can be traced back to either the formation ofdefects during the homoepitaxial growth of the crystallites induced by ion bombardment or theincreased re-nucleation of strongly misoriented grains.

6.3. Heteroepitaxal growth of diamond films

To date, heteroepitaxial growth on Si is limited to microwave plasma CVD [73,98] hot-filamentCVD [75,163]. Significant progresses have been achieved in obtaining heterogeneous nucleation oforiented diamond crystallites via bias-enhanced nucleation.

It has been noted that if the growth temperature is lower and the nucleation period (incubationtime required before the appearance of nuclei) is shorter, diamond can nucleate directly on aSi substrate without forming a SiC intermediate layer at the interface. The maximum areaof the epitaxial layer (single crystalline layer) of diamond grown directly on Si can reach20 � 20 mm2.

Fig. 13. Azimuthal {2 2 0} pole density distribution for diamond films. Results were obtained from a variation of thebiasing time. (Ref. [186]).


So far, the epitaxial crystallite of diamond grown on Si (1 0 0) has the followingcrystallographic relationships to the substrate: diamond(1 0 0)//silicon(1 0 0) and diamond[1 1 0]//silicon[1 1 0].

Textured growth utilizes an `evolutionary' selection (Van der Drift growth) if the growth processproceeds to a large thickness, and only permits crystallites which are oriented in the fastest growthdirection (approximately parallel to the general growth direction) to grow further. The fastest growthdirection can be varied by changing the process parameters, the most important of which are thesubstrate temperature and methane concentration [183]). Recently, it was found that the applicationof a proper bias potential to the substrate also leads to a strong selection of the growth direction[182]. The selection effect is attributed to the direction-dependent etching of the diamond crystal byH+ ions [184]. Therefore, a textured film texture can be achieved for thin films. Both of theaforementioned approaches have been used to prepare epitaxially oriented films [183,139].

Heteroepitaxial [0 0 1]-oriented diamond films with considerably increased lateral grain sizeand strongly improved orientational perfection could be prepared by microwave plasma-assistedchemical vapor deposition using a [0 0 1]-textured growth process on Si (0 0 1) substrates followedby a [1 1 0] step-flow growth process [168,185]. The diamond films were characterized by atomicforce microscopy, scanning electron microscopy, and transmission electron microscopy. The resultsindicate that diamond crystals increase their lateral dimensions at the (0 0 1) film surface either bycoalescence of grains combined with a termination of the propagation of grain boundaries (Fig. 15)or by changing the grain boundary plane orientations from preferentially vertical to preferentiallyparallel directions with respect to the (0 0 1) growth faces. In the second case, the grains withrelatively large angle deviation from the ideal epitaxial orientation are overgrown by those withrelatively small angle deviation. As a result, the degree of orientational perfection of the filmsimproves considerably in comparison to that of films prepared by the established process of [0 0 1]-textured growth. The presence of boron in the gas phase was found to strongly enhance the step-flowlateral grain growth. It was possible to achieve deposition of a thin boron-doped diamond filmscharacterized by a full width at half maximum value of the measured tilt angle distribution of only2.18.

Fig. 14. Variation of the process time window for oriented nucleation with the bias voltage (Ref. [180]).


6.4. Interface between diamond and silicon

Early studies revealed that the growth of diamond on Si occurred through a b-SiC transfer layer[186,187]. This perhaps is expected considering that the lattice misfit between diamond and b-SiC(�22%) is much smaller than that between diamond and silicon (52%). However, recently, a directlattice observation of the interface between silicon and epitaxially grown diamond crystal has beensuccessfully performed [161,162]. It is clearly demonstrated that the diamond nuclei are grown indirect contact with the silicon substrate (Fig. 8). Because the ratio of the lattice constants of siliconand diamond is close to 1.5, a nearly perfect 3 to 2 correspondence (1.5% mismatch) of latticespacing is seen at the interface, e.g. every three Si (1 1 1) fringes are matched well with fourdiamond (1 1 1) fringes (the mismatch is about 1.5%). Individual 608 interface misfit dislocations forevery third (1 1 1) atomic plane, with a spacing between two such dislocations of about 7 AÊ , can beclearly identified (Fig. 8). The presented results are reproducible for many of the epitaxially orienteddiamond grains observed and they provide strong evidence that diamond crystal can be epitaxiallygrown directly on Si. From the experimental results, it can be concluded that if the growthtemperature is not too high and the incubation time is sufficiently short, diamond can grow directlyon the Si substrate, avoiding the formation of the b-SiC interlayer. A thin epitaxial SiC intermediatelayer is unnecessary for diamond heteroepitaxy. Recent theoretical investigations of the interfacestructure between diamond and silicon have also provided support for the direct growth of diamondon silicon [188,189].

6.5. Possible improvements of the heteroepitaxy of diamond on silicon

Although great progresses have been made in the heteroepitaxy of diamond on silicon, largearea, single crystalline diamond layers have still not been obtained. In order to realize large-area andhigh-quality epitaxy, further efforts are needed. For the heteroepitaxy of diamond directly on silicon,the following approaches may be taken into consideration.

Fig. 15. High-resolution lattice image of a grain boundary region, showing disappearance of a 28 low-angle grain boundary(Ref. [161]).


6.5.1. Improvement of the surface treatmentThis is the first and essential requirement. The substrate surface must be clean, without surface

contamination and oxidation, and the dangling bonds of the Si surface atoms must be saturated byhydrogen. During the process of bias-enhanced nucleation, surface damage due to ion bombardmentand the formation of amorphous carbon must be avoided.

6.5.2. Control of nucleation must be as exact as possible

The bias voltage, biasing time, and pressure strongly influence crystal orientations. Thebombardment of energetic ions above a critical energy is necessary for the formation of nuclei. Anegative influence by ion bombardment on the alignment of the diamond grains has also beendemonstrated. An improved epitaxy requires a compromise of the positive and negative influences.On the other hand, at present the vacuum for CVD diamond nucleation and growth is low and theresidual gas in the chamber would contaminate or even oxidize the Si surface and change the surfacestatus. Therefore improvement of the base vacuum in the growth chamber and the purity of gassource would be helpful in achieving better epitaxy.

7. Surface science studies of CVD diamond growth

The nucleation, growth, and epitaxy of CVD diamond all proceed on the surface. Theseessential phenomena belong to the fields of surface science, and thus surface techniques are usefuland powerful tools in studying CVD diamond growth. Especially, the growth of diamond under lowpressure reveals a very complicated process, which to a certain extent is similar to the growth processof some functional organic materials, i.e. the growth process is mainly completed by the tailoringand interconnecting of functional radicals. Diamond growth under low pressure is not a simpledeposition of carbon atoms, for in this way only graphite is obtained; diamond growth is completedby the surface adsorption, migration, and interconnection of hydrocarbon radicals as well as thereconstruction (reformation) of carbon atoms in the lattice. Such a complicated process can beclarified only by surface techniques.

The relationship between CVD diamond studies and surface techniques has been extensivelydiscussed in the review articles entitled Characterization of diamond films written by Zhu et al. [190]and Hydrogen-terminated diamond surface and interface by Kawarada [191]. Zhu's paper alsointroduced the principle of various surface techniques and their application in diamond surfacecharacterization, which we will not repeat in this paper. Here, we will just highlight the usefulness ofsurface techniques in the studies of diamond nucleation, growth, and epitaxy.

7.1. In preparation of substrate for nucleation and growth

Prior to diamond nucleation and growth, the substrate must first be cleaned completely in orderto obtain an uncontaminated surface; second, the clean surface then undergoes surface pre-treatmentin some cases. Of the many surface cleaning and pre-treatment methods that are available, which oneis effective? How can a surface be cleaned and still maintain the integrity of the surface structure? Toevaluate the results from the different cleaning methods, surface techniques such as Auger electronspectroscopy (AES), X-ray photoelectron spectroscopy (XPS), low energy electron diffraction(LEED), reflected high-energy electron diffraction (RHEED), high resolution electron energy lossspectroscopy (HREELS), scanning tunneling microscope (STM), and atomic force microscope


(AFM) are often used. They identify surface composition, surface structure, surface bonding, as wellas surface electronic structure after the cleaning and pre-treatment stages.

7.2. In nucleation studies

CVD diamond hardly grows on the surface of other materials, however, under certain conditionsit can be grown on some materials. Why? In order to answer this question, it is essential to study theproperties of the substrate surface including its composition, structure, reconstruction, morphology,and electronic structure by using surface techniques. The reason for this is that nucleation dependsheavily on the status of the substrate surface. Why are the data of diamond nucleation obtained fromdifferent laboratories frequently in disagreement? The main explanation is that the substrate surfacesused for nucleation and growth are not well-defined, therefore, due to differences in the surfacestatus, it is not surprising that the nucleation results vary and are even controversial. In order toobtain `clean results' of nucleation process, a well-defined surface of the substrate must be used as astarting point. Thus, surface techniques are needed to identify the substrate surface. Furthermore,surface techniques are also imperative for investigating factors pertaining to the nucleation process,such as the adsorption sites, configuration of adsorbates, and their evolution.

7.3. In diamond growth studies

Many debates still remain over the mechanism of CVD diamond growth. Which species (CH,CH2, CH3, C2H2, etc.) is the precursor for diamond growth? Where are the adsorption sites of theseprecursors? How do they interconnect with each other to form a diamond crystal? The answers tothese questions are still unclear. Although many models have been proposed, for the lack ofexperiment data at the atomic level, a satisfactory mechanism has not yet been developed. Here,surface techniques can be utilized to study the issues mentioned above and provide needy data forunderstanding diamond growth.

LEED (RHEED) and STM (AFM) offer information relating to the surface and reconstructionbefore and after adsorption of hydrocarbon radicals and hydrogen atoms. Both techniquescomplement each other. Here, we just introduce high-resolution-electron energy loss spectroscopy(HREELS), which, similar to infrared (IR) spectroscopy, has become a powerful tool for studyingsurface vibration. Given the information of the vibration mode (number of loss peak) and vibrationfrequency (position of loss peaks) obtained from the HREELS spectra, we can deduce theconfiguration of the adsorbates (CH, CH2, CH3, etc.) and the adsorption sites on the surface as wellas the evolution of the adsorbates during growth. Therefore, HREELS provides the data necessaryfor understanding the diamond growth mechanism. Details of HREELS can be found in Refs. [191±194]. When HREELS is combined with other techniques, like AES, LEED/STM in an UHV system,more complete and reliable information about diamond growth can be obtained.

8. Summary

Diamond possesses numerous unmatchable properties in mechanical, thermal, optical, andelectronic aspects, therefore it demonstrates extensive application prospects. Unfortunately, naturaldiamond is rare and expensive, which prohibits extensive usage. At the same time, high-pressure-high-temperature-diamond can only be obtained in small crystallite amounts, which also cause manylimitations to their applications. The deposition of diamond films by chemical vapor deposition has


served as a great breakthrough in the last decade of the 20th century and has made many applicationsof diamond possible. In order to increase the growth rate and improve the quality of diamond films,many methods have been developed in succession, including HF-CVD, MP-CVD, RF-P-CVD, DC-P-CVD, ECR-MP-CVD, and combustion flame-CVD. Due to the intensive efforts expended on thepart of research scientists and industrialists worldwide, CVD methods have reached a rather highlevel at present. Diamond is opening the door to the fields of coating, cutting tool, optical, andthermal management applications. Currently, several products with respectable quantities ofdiamond have entered the market. This is a great advance in diamond development. However, despitethe many successes achieved in CVD diamond, serious difficulties still lie ahead of us. So far, westill have not succeeded in growing mono-crystalline diamond film with a large area, which limitsgreatly the important application of diamond in high-temperature electronic devices. Naturally, tosolve the problem of diamond heteroepitaxy is a big challenge, and it has become a hot topic indiamond research and attracted many researchers to this study.

In order to solve such complicated and difficult problems, first of all, we must identify themechanism of diamond nucleation, growth, and epitaxy, because only on the basis of mechanisticunderstanding can we succeed in large-area heteroepitaxial growth of diamond.

In the studies of diamond nucleation, growth, and epitaxy, surface techniques serve as veryuseful and powerful tools, and of course, the combination of surface techniques with other analyticalmethods will be even more powerful. At present, it is clear that diamond facets with differentorientations have their own growth modes, atomic hydrogen and oxygen have essential roles in thegrowth process, and the appearance of facets with different orientation can be regulated, etc. In thenucleation aspect, the effect of substrate surface conditions on nucleation has been studied, the roleof atomic hydrogen and the concentration of hydrocarbon radicals on the nucleation has beenemphasized, and several methods for high density nucleation of diamond on mirror-polished Si orSiC have been developed. As for heteroepitaxy, excellent and reproducible results have beenobtained mainly on Si and SiC. Highly oriented diamond films grown epitaxially on Si or SiC can beachieved in many laboratories without difficulty. The diameter of the crystallites can reach 20 mm.Although many achievements have been made in diamond nucleation, growth, and epitaxy, they stillfall short of a thorough understanding of the mechanism. Therefore, a lot of studies are still required.Nevertheless, we will anticipate success in large-area diamond heteroepitaxy in the near future. Thisis the main line of the paper.

References

[1] J.C. Angus, C.C. Hayman, Science 214 (1988) 913.[2] W. Yarbrough, R. Messier, Science 247 (1990) 688.[3] J.E. Field, The Properties of Diamond, Academic Press, Oxford, 1979.[4] R.F. Davis (Ed.), Diamond Films and Coatings, Noyes Publications, New Jersey, 1992[5] L.S. Pan, D.R. Kania (Eds.), Diamond: Electronic Properties and Applications, Kluwer Academic Publishers,

Boston, 1995.[6] J. Wilks, E. Wilks (Eds.), Properties and Application of Diamond, Butterworth Heinemann Ltd.[7] Max N. Yoder, Diamond properties and application, in: Robert F. Davis (Ed.), Diamond Films and Coatings,

Development, Properties, and Applications, Noyes Publications, Park Ridge, New Jersey, p. 5.[8] R. Seitz, The end of the age of sand, presented at the Gorham Advanced Materials Institute, Marco Island, FL 15±17

Oct., 1989.[9] A. Backon, A. Szymanski, Practical uses of Diamond, in polish, translated into English by P. Daniel, pp. 29±31.

[10] A.P. Rudenko et al., DAN SSSR, Ser, mat. Fiz. 163(5) (1965) 1169.[11] R.C. Eden, Diamond Rel. Mater. 25±7 (1993) 1051.


[12] A. Altshuler, Diamond Technology in the USSR, Delphic Associates Press, 1989.[13] K. Baba, Y. Aikawa, N. Shohata, H. Yoneda, K.-I. Ueda, NEC Research Development 363 (1995) 369.[14] H. Yoneda, K.-I. Ueda, Y. Aikwaa, K. Baba, N. Shohata, Appl. Phys. Lett. 664 (1995) 460.[15] A.T. Collins, E.C. Lightowlers, in: J.E. Field (Ed.), The Properties of Diamond, Academic Press, San Diego, 1979.[16] V.K. Bazhenov, I.M. Vikulin, A.G. Gontar, Sov. Phys. Semicond. 19 (1985) 829.[17] A.T. Collins, Semiconductor Sci. Technol. 4 (1989) 605.[18] K. Shenai, R.S. Scott, B.J. Baliga, IEEE Trans. Electron Dev. 36 (1989) 1811.[19] M.W. Geis, Proc. IEEE 79 (1991) 669.[20] G.S. Gildenblat, S.A. Grot, C.W. Hatfield, A.R. Badzian, IEEE Electron Dev. Lett. 12 (1991) 37.[21] H. Shiomi, Y. Nishibayashi, N. Fujimori, in: R. Messier, J.T. Glass, J.E. Butter, R. Roy (Eds.), Proc. 2nd Intl. Conf.

New Diamond Sci. Technol. Mat. Res. Soc., Pittsburgh, 1991, p. 975.[22] K. Nishimura, K. Das, J.T. Glass, J. Appl. Phys. 69 (1991) 3142 .[23] N. Fujimori, T. Imai, A. Doi, Vacuum 36 (1986) 99.[24] M.W. Geis, in J.T. Glass, R. Messier, N. Fujimori (Eds.), Diamond, Silicon Carbide and Related Wide Bandgap

Semiconductors, 162, Mat. Res. Soc. Symp. Proc., 1990, p. 15.[25] V.S. Vavilov, M.I. Guseva, E.A. Konorova, V.F. Sergienko, Sov. Phys. Semicond. 4 (1970) 6.[26] G. Braunstein, R. Kalish, J. Appl. Phys. 54 (1983) 2106.[27] J.F. Prins, Phys. Rev. B. 38 (1988) 5576.[28] G.S. Sandhu, M.L. Swanson, W.K. Chu, Appl. Phys. Lett. 55 (1989) 1397.[29] P.R. de la Houssaye, C.M. Penchina, C.A. Hewett, R.G. Wilson, J.R. Zeidlers, in: R. Messier, J.T. Glass, J.E. Butler,

R. Roy (Eds.), Proc. 2nd Intl. Conf. New Diamond Sci. Techn., Mat. Res. Soc., Pittsburgh, 1991, p. 113.[30] G. Popovici, M.A. Prelas, Diamond Rel. Mater. 4 (1995) 1305.[31] G.S. Sandhu, Ph.D. Thesis, Univ. of North Carolina at Chapel Hill, 1989.[32] J.F. Prins, J. Phys. D: Appl. Phys. 22 (1989) 1562.[33] T.P. Humphreys, J.V. LaBrasca, R.J. Nemanich, K. Das, J.B. Posthill, Electron. Lett. 27 (1991) 1515.[34] S.R. Sahaida, D.G. Thompson, in: T.D. Moustakos, J.I. Pankove, Y. Hamakawa (Eds.), Mater. Res. Soc. Symp. Proc.

1992.[35] G.S. Gildenblat, S.A. Grot, C.W. Hatfield, A.R. Badzian, IEEE Electron Dev. Lett. 12 (1991) 37.[36] C.R. Zeisse, C.A. Hewett, R. Nguyen, J.R. Zeidler, R.G. Wilson, IEEE Electron Dev. Lett. 12 (1991) 602.[37] W.G. Eversole, U.S. Patent 3,030,188, April 17, 1962.[38] D.J. Poferl, N.C. Gardner, J.C. Angus, J. Appl. Phys. 444 (1973) 1428.[39] B.V. Spitsyn, L.L. Bovilov, B.V. Derjaguin, J. Crys. Growth 52 (1981) 219.[40] S. Matsumoto, Y. Sato, M. Tsutsimi, N. Setaka, J. Mat. Sci. 17 (1982) 3106.[41] R.A. Rudder, J.B. Posthill, R.J. Markunas, Semiannual Report on Semiconductor Diamond Technology, SDIO/ONR

contract N00014-86-C-0460, November 1989[42] R.A. Rudder, J. B Posthill, R.J. Markunas, Electron. Lett. 25 (1989) 1220.[43] W.M. Lau, S.-T. Lee, F. Qin, Surface Modification Technologies V, 1992, p. 263.[44] X.S. Sun, N. Wang, H.K. Woo, W.J. Zhang, G. Yu, I. Bello, C.S. Lee, S.T. Lee, Diamond & Relat. Mater., in press.[45] N. Fujimori, A. Ikegaya, T. Imai, K. Fukushima, N. Ota, in: J. Dismukes, et al. (Eds.), Diamond and Diamond-like

Films, Electrochem. Soc. Proc., Vol PV 89-12, Pennington, NJ, 1989, p. 465.[46] Qijin Chen, Zhangda Lin, Appl. Phys. Lett. 68(17) (1996) 2450.[47] Qijin Chen, Jie Yang, Zhangda Lin, Appl. Phys. Lett. 67(13) (1995) 1853.[48] A.R. Badzian, R.C. De Vries, Mater. Res. Bull. 23 (1988) 385.[49] D.E. Meyer, T.B. Kustka, R.O, Dillon, J.A. Woollam, in: A. Feldman, S. Holly (Eds.), Diamond Optics II, SPIE

Proc., 1146:21, Bellingham, WA, 1990.[50] K. Kurihara, K. Sasaki, M. Kawarada, N. Koshino, Appl. Phys. Lett. 526 (1988) 437.[51] M. Kawarada, K. Kurihara, K. Sasaki, A. Teshima, K. Koshino, in: A. Feldman, S. Holly (Eds.), Diamond Optics II,

SPIE Proc. 1146, Bellingham, WA, 1990, p. 28.[52] Akio Hiraki, Hiroshi Kawarada, Jin Wei, Jun-Ichi Suzuki, Surface and Coating Technology 43/44 (1990) 10.[53] Takuya Yara, Motokazu Yuasa, Manabu Shimizu, Hiroshi Makita, Akimitsu Hatta, Jun-ichi Suzuki, Toshimichi Ito,

Akio Hiraki, Jpn J. Appl. Phys. 33(1) (1994) 4405.[54] D. Mantei, Zoltan Ring, Mark Schweizer, Spirit Tlali, Howard E. Jackson, Jpn. J of Appl. Phys. 35 (1996) 2516.[55] Y. Hirose, N. Kondo, Program and Book of Abstracts, Japan Applied Physics 1988 spring meeting, March 29, 1988,

p. 434.[56] M.A. Cappelli, P.H. Paul, J. Appl. Phys. 675 (1990) 2596.[57] J.J. Lander, J. Morrison, Surf. Sci. 2 (1964) 553.[58] B.V. Spitsyn, L.L. Bouilov, B.V. Deryagin, J. Cryst. Growth 52 (1981) 219.[59] S. Matsumoto, Y. Sato, M. Kamo, J. Tanaka, N. Setaka, in Proceedings of the 7th International Conference on

Vacuum Metallurgy, Tokyo, Japan, 1982. Iron and Steel Institute of Japan, Tokyo, Japan, 1982, pp. 386±391.[60] N. Fujimori, T. Imai, H. Nakahata, Symposium N Plasma Assisted Deposition of New Materials, MRS Fall 1987

Meeting. Boston, MA. Nov. 30±Dec. 5, 1987.


[61] J.C. Angus, Y. Wang, M. Sunkara, Ann. Rev. Mater. Sci. 21 (1991) 221.[62] W.G. Eversole, 1962, U S Patent No. 3, 030, 187; 3, 030, 188.[63] K. Mitsuda, Y. Kojima, T. Yoshida, K. Akashi, J. Mater. Sci. 22 (1987) 1557.[64] J.C. Angus, Z. Li, M. Sunkara, R. Gat, A.B. Anderson, S.P. Mehandru, M.W. Geis, in: A.J. Purdes, J.C. Angus, R.F.

Davis, B.M. Meyerson, K.E. Spear, M. Yoder (Eds.), Proceedings of the 2nd International Symposium on DiamondMaterials, The Electrochemical Society, Pennington, N J, 1991, p. 125.

[65] Z. Feng, K. Komvopoulos, I.G. Brown, D.C. Bogy, J. Appl. Phys. 744 (1993) 2841.[66] A.A. Morrish, P.E. Pehesson, Appl. Phys. Lett. 59 (1991) 417.[67] J. Singh, M. Vellaikal, J. Appl. Phys. 73 (1993) 2831.[68] K.V. Ravi, M. Peters, L. Plano, S. Yokota, M. Pinneco, in First International Conference on the New Diamond

Science and Technology, Program and Abstr. 24±26, Oct. 1988. (JNDF, Tokyo, 1988), pp. 96±97.[69] K.V. Ravi, C.A. Koch, Appl. Phys. Lett. 57 (1992) 348.[70] K.V. Ravi, C.A. Koch, H.S. Hu, A. Joshi, J. Mater. Res. 5 (1990) 2356.[71] S. Yugo, T. Kanai, T. Kimura, T. Muto, Appl. Phys. Lett. 58 (1991) 1036.[72] X. Jiang, C.-P Klages, Diamond Relat. Mater. 2 (1993) 1112.[73] X. Jiang, C.-P. Klages, R. Zachai, M. Hartweg, H.-J. FuÈsser, Appl. Phys. Lett. 62 (1993) 3438.[74] B.R. Stoner, J.T. Glass, Appl. Phys. Lett. 60 (1992) 698.[75] F. Stubhan, M. Ferguson, H.-J. FuÈsser, R.J. Behm, Appl. Phys. Lett. 66 (1995) 1900.[76] A. Sawabe, T. Inuzuka, Thin Solid Films 137 (1986) 89.[77] K. Suzuki, A. Swabe, H. Yasuda, T. Inuzuka, Appl. Phys. Lett. 50 (1987) 728.[78] C.P. Chang, D.L. Flamm, D.E. Ibbotson, J.A. Mucha, J. Appl. Phys. 63 (1988) 1744.[79] P.A. Denning, D.A. Stevenson, Proc. of the 2nd Int. Conf. on New Diamond Sci. and Techn. Pittsburg, PA, MRS,

1991, pp. 403±408.[80] B. Singh, O. Mesker, A.W. Levine, Y. Arie, Growth of Polycrystalline Diamond Particles and Films by Hot-

Filament Chemical Vapor Deposition, Report CN-5300, David Sarnoff Research Center, SRI International,Princeton, NJ, 1988.

[81] B. Singh, Y. Arie, A.W. Levine, O.R. Mesker, Appl. Phys. Lett. 526 (1988) 451.[82] C.P. Beetz Jr, T.A. Perry, Hot and Cold Wall Filament-Assisted Growth of Diamond Films, General Motors

Research Publication GMR-6093, November 24, 1987, pp. 1±21.[83] S. Yugo, T. Kimura, H. Kanai, in: S. Saito, O. Fukunaga, M. Yoshihara (Eds.), Science and Technology of New

Diamond, KTK Scientific Pub., Tokyo, 1990, pp. 119±128.[84] S. Iijima, Y. Aikawa, K. Baba, J. Mater. Res. 6 (1991) 1491.[85] S.P. Murarka, M.C. Peckerar, Electronic Materials Science and Technology, Academic Press Inc. San Diego, CA,

1989.[86] S. Yugo, A. Izumi, T. Kanai, T. Muto, T. Kimura, Proc. of the 2nd Int. Conf. on New Diamond Science and

Technology, Pittsburg, PA, MRS, 1991, pp. 385±390.[87] P.A. Dennig, D.A. Stevenson, Proc. of the 2nd Int. Conf. on New Diamond Science and Technology, Pittsburg, PA,

MRS, 1991, pp. 403±408.[88] Jie Yang, Xiaowei Su, Qijin Chen, Zhangda Lin, Appl. Phys. Lett. 66(24) (1995) 3284.[89] P.K. Bachmann, R. Weimer, W. Drawl, Y. Liou, R. Messier, Diamond Technology Initiative Symposium, Paper T20,

SDIC/IST ONR, Crystal City, VA, July 12±14, 1988.[90] C.P. Chang, D.L. Flamm, D.E. Ibbotson, J.A. Mucha, J. Appl. Phys. 63 (1988) 1744.[91] M. Marchywka, P.E. Pehrsson, D.J. Vestyck, D. Moses, Appl. Phys Lett. 63 (1993) 3521.[92] A.R. Kirkpatrick, B.W. Ward, N.P. Sconomou, J. Vac. Sci. Tech. B7 (1989) 19471949.[93] B.R. Stoner, G.-H.M. Ma, S.D. Wolter, J.T. Glass, Phys. Rev. B45 (1992) 11067.[94] S. Yugo, T. Kimura, T. Kanai, Diamond and Relat. Mater. 2 (1993) 328.[95] J. Gerber, S. Sattel, K. Jung, H. Ehrhard, J. Robertson, Diamond and Relat. Mater. 4 (1995) 559.[96] B.R. Stoner, G.H. Ma, S.D. Wolter, W. Zhou, Y.-C. Wang, R.F. Davis, J.T. Glass, Diamond and Relat. Mater. 2

(1993) 142.[97] X. Jiang, K. Schiffmann, C.-P. Klages, Phys. Rev. B50 (1994) 8402.[98] X. Jiang, C.-P. Klages, Phys. Stat. Sol. 154 (1996) 175.[99] S.T. Lee, Y.W. Lam, Zhangda Lin, Yan Chen, Qijin Chen, Phys. Rev. B. 5524 (1997) 15937.

[100] C. Sun, W.J. Zhang, J. Yang, I. Bello, C.S. Lee, S.T. Lee, to be published.[101] Zhaohui Liu, B.Q. Zong, Zhangda Lin, Thin Solid Films 254 (1995) 3.[102] M. Tsuda, M. Nakajima, S. Oikawa, J. Am. Chem. Soc. 108 (1986) 5780.[103] M. Tsuda, M. Nakajima, S. Oikawa, Jpn. J. Appl. Phys. 26 (1987) L527.[104] C.J. Chu, M.P. D'Evelyn, R.H. Hauge, J.L. Margrave, J. Appl. Phys. 70 (1991) 1659.[105] L.R. Martin, M.W. Hill, J. Mater. Sci. Lett. 9 (1990) 621.[106] S.J. Harris, L.R. Martin, J. Mater. Res. 5 (1990) 2313.[107] S.J. Harris, Appl. Phys. Lett. 56 (1990) 2298.[108] M. Frenklach, K.E. Spear, J. Mater. Res. 3 (1988) 133.


[109] D. Huang, M. Frenklach, M. Maroncelli, J. Phys. Chem. 92 (1988) 6379.[110] M. Frenklach, H. Wang, Phys. Rev. B 43 (1991) 1520.[111] D. Huang, M. Frenklach, J. Phys. Chem. 95 (1991) 3692.[112] F.G. Celii, J.E. Butler, Appl. Phys. Lett. 54 (1989) 1031.[113] M. Frenklach, J. Appl. Phys 65 (1989) 5142.[114] D.G. Goodwin, G.G. Gavillet, J. Appl. Phys. 68 (1990) 6393.[115] D.M. Gruen, MRS Bull. 239 (1998) 32.[116] P.C. Redfern, D.A. Horner, L.A. Curtiss, D.M. Green, J. Phys. Chem. 100 (1996) 11654.[117] Biwu Sun, Xiaopin Zhang, Qinzhe Zhang, Zhangda Lin, Appl. Phys. Lett. 62(1) (1993) 31.[118] T. Shimanouchi, Table of Vibrational Frequencies, Natl. Bur. Stand. Ref. Data Ser., Natl. Bur. Stand. (U.S.) Circ.

No. 39, U.S. GPO, Washington, DC, 1968, Consolidated Vols. I and II[119] J. Phys. Chem. Ref. Data 6 (1977) 993.[120] H. Ibach, D.L. Mills, Electron Energy Loss Spectroscopy and Surface Vibrations, Academic New York, 1982.[121] L.B. Zhao, K.A. Feng, Z. D Lin, Diamond Rel. Mater. 3 (1993) 155.[122] A.V. Hamza, G.D. Kubiak, R.H. Stulen, Surf. Sci. 237 (1990) 35.[123] S.P. Mehandru, A.B. Anderson, Surf. Sci. 248 (1991) 369.[124] T. Tsuno, T. Imai, Y. Nishibayasi, K. Hamada, N. Fujimori, Jpn. J. Appl. Phys. 30 (1991) 1063.[125] L.F. Sutcu, Appl. Phys. Lett. 60 (1992) 1685.[126] Biwu Sun, Xiaopin Zhang, Zhangda Lin, Phys. Rev. B. 47(15) (1993) 9816.[127] K.A. Feng, Z.D. Lin, J. Phys Chem. Solids 556 (1994) 525.[128] K. Kobashi, K. Nishimura, Y. Kawate, T. Horiuchi, Phys. Rev. B. 38 (1988) 4067.[129] Y.J. Baik, K.Y. Eun, Applications of Diamond Films and Related Materials, Elsevier, New York, 1991, p. 521.[130] Ch. Wild, N. Herres, P. Koidl, J. Appl. Phys. 68 (1990) 973.[131] Zhangda Lin, Biwu Sun, J. Vacuum Society Japan 37(7) (1994) 579.[132] S.J. Harris, A.M. Weiner, J. Appl. Phys. 67 (1990) 6520.[133] S.J. Harris, D.N. Betton, A.M. Weiner, S.J. Schmieg, J. Appl. Phys. 66 (1989) 5353.[134] B.J. Waclawski, D.T. Pierce, N. Swanson, R.J. Celotta, J. Vac. Sci. Technol. 21 (1982) 368.[135] T.R. Anthony, Vacuum 41 (1990) 1356.[136] K. Ishibashi, S. Furukawa, Jpn. J. Appl. Phys. 24 (1985) 912.[137] X. Jiang, W.J. Zhang, M. Paul, C.-P. Klages, Appl. Phys. Lett. 68 (1996) 1927.[138] W.J. Zhang, X. Jiang, Appl. Phys. Lett. 68 (1996) 2195.[139] W.J. Zhang, X. Jiang, Y.B. Xia, J. Appl. Phys. 82 (1997) 1896.[140] Y. Muranaka, H. Yamashita, H. Miyadera, J. Vac. Sci. Technol. A. 9 (1991) 76.[141] S.J. Harris, A.M. Weiner, Appl. Phys. Lett. 55 (1989) 2179.[142] Y. Liou, A. Inspektor, R. Weimer, D. Knight, R. Messier, J. Mater. Res. 5 (1990) 2305.[143] Biwu Sun, Xiaopin Zhang, Qinzhe Zhang, Zhangda Lin, J. Appl. Phys. 73(9) (1993) 4614.[144] S.-T. Lee, Y.W. Lam, Zhangda Lin, Yan Chen, Zhiqiang Gao, Phys. Rev. B. 54 (1996) 14185.[145] A. Badzian, T. Badzian, X.H. Wang, Carbon 28 (1990) 804.[146] M. Kamo, H. Yurimoto, Y. Sato, Appl. Surface Sci. 33/34 (1988) 553.[147] B.V. Spitsyn, L.L. Bouilov, B.V. Derjaguin, J. Crystal. Growth 42 (1981) 219.[148] B.V. Derjaguin, B.V. Spitsyn, A.E. Gorodetsky, A.P. Zakharov, L.I. Bouilov, A.E. Aleksenko, J. Crystal. Growth 31

(1975) 44.[149] H. Nakazawa, Y. Kanayama, M. Kamo, K. Osumi, Thin Solid Films 151 (1987) 199.[150] K. Kobashi, K. Nishizawa, Y. Kawate, T. Horiuchi, Phys. Rev. B. 38 (1988) 4067.[151] N. Fujimori, T. Imai, H. Nakahata, H. Shiomi, Y. Nishibayashi, Mater. Res. Soc. Symp. Proc. 162 (1990) 109.[152] M. Yoshikawa, H. Ishida, A. Ishitani, T. Murakami, S. Koizumi, T. Inuzuka, Appl. Phys. Lett. 57 (1990) 428.[153] S. Koizumi, T. Murakami, K. Suzuki, T. Inuzuka, Appl. Phys. Lett. 57 (1990) 563.[154] S. Koizumi, T. Inuzuka, Jpn. J. Appl. Phys. 32 (1993) 3920.[155] A. Argoitia, J.S. Ma, J.C. Angus, Appl. Phys. Lett. 63 (1993) 1336.[156] A. Argoitia, J.C. Angus, J.S. Ma, L. Wang, P. Pirouz, W. Lambrecht, J. Mater. Res. 9 (1994) 1849.[157] R.O. Jones, O. Gunnarsson, Rev. Mod. Phys. 61 (1989) 689.[158] T. Tomizuka, S. Shikata, Jpn. J. Appl. Phys. 32 (1993) 3938.[159] R.C. De Vries, Cubic Boron Nitride: Handbook of Properties, G.E. Representative. No. 72 CRD, 1972, p. 178.[160] T. Suzuki, A. Argoitia, Phys. Stat. Sol. 154 (1996) 239.[161] X. Jiang, C.L. Jia, Appl. Phys. Lett. 679 (1995) 1197.[162] C.L. Jia, K. Urban, X. Jiang, Phys. Rev. B 52 (1995) 5164.[163] Jie Yang, Zhangda Lin, Li-Xin Wang, Sing Jin, Ze Zhang, Appl. Phys. Lett. 65(25) (1994) 3203.[164] Ke-an Feng, Jie Yang, Zhangda Lin, Phys. Rev. B. 51(4) (1995) 2264.[165] Jie Yang, Zhangda Lin, Lixin Wang, Sing Jin, Ze Zhang, J. Phys. D: Appl. Phys. 28 (1995) 1153.[166] Qijin Chen, Li-Xin Wang, Ze Zhang, Jie Yang, Zhangda Lin, Appl. Phys. Lett. 68(2) (1996) 176.[167] X. Jiang, C.-P. Klages, M. RoÈsler, R. Zachai, M. Hartweg, H.-J. FuÈsser, Appl. Phys. A 57 (1993) 483.


[168] X. Jiang, K. Schiffmann, C.-P. Klages, D. Wittorf, C.L. Jia, K. Urban, W. Jaeger, J. Appl. Phys. 83 (1998) 2511.[169] H. Kawarada, Ch. Wild, N. Herres, R. Locher, P. Koidl, J. Appl. Phys. 81 (1997) 3490.[170] A. Argoitia, J.C. Angus, L. Wang, X.I. Ning, P. Pirouz, J. Appl. Phys. 73 (1993) 4305.[171] W. Zhu, P.C. Yang, J.T. Glass, Appl. Phys. Lett. 63 (1993) 1640.[172] T. Tachibana, Y. Yokota, K. Miyata, K. Kobashi, Y. Shintani, Diamond Rel. Mater. 5 (1996) 197.[173] W. Liu, D.A. Tucker, P. Yang, J.T. Glass, J. Appl. Phys. 78 (1995) 1291.[174] K. Ohtsuka, K. Suzuki, A. Sawabe, T. Inuzuka, Jpn. J. Appl. Phys. Part 2 35 (1996) L1072.[175] K. Ohtsuka, H. Fukuda, K. Suzuki, A. Sawabe, Jpn. J. Appl. Phys. Part 2. 36 (1997) L1214.[176] T. Saito, S. Tsuruga, N. Ohya, K. Kusakabe, S. Morooka, H. Maeda, A. Sawabe, K. Suzuki, Diamond Rel. Mater. 7

(1998) 1381.[177] M. Schreck, H. Roll, B. Stritzker, Appl. Phys. Lett. 74 (1999) 650.[178] X. Jiang, C.-P. Klages, R. Zachai, M. Hartweg, H.-J. Fuesser, Diamond Rel. Mater. 2 (1993) 407.[179] X. Jiang, K. Schiffimann, A. Westphal, C.-P. Klages, Appl. Phys. Lett. 639 (1993) 1203.[180] M. Schreck, K.H. Thuerer, R. Klarmann, B. Stritzker, J. Appl. Phys. 81 (1997) 3096.[181] M. Schreck, K.H. Thuerer, B. Stritzker, J. Appl. Phys. 81 (1997) 3092.[182] X. Jiang, W.J. Zhang, C.-P. Klages, Phys. Rev. B58 (1998) 7064.[183] C. Wild, P. Koidl, W. Mueller-Sebert, H. Walcher, R. Kohl, N. Herres, R. Locher, R. Samlenski, R. Bren, Diamond

Rel. Mater. 2 (1993) 158.[184] W.J. Zhang, X. Jiang, C.-P. Klages, J. Cryst. Growth 171 (1997) 485.[185] X. Jiang, C.L. Jia, Appl. Phys. Lett. 69 (1996) 3902.[186] S.D. Wolter, B.R. Stoner, J.T. Glass, P.J. Ellis, D.S. Buhaenko, C.E. Jenkins, P. Southworth, Appl. Phys. Lett. 6211

(1993) 1215.[187] W. Zhu, X.H. Wang, B.R. Stoner, G.H.M. Ma, H.S. Kong, M.W.H. Braun, J.T. Glass, Phys. Rev. B. 47 (1993) 6529.[188] R.Q. Zhang, W.L. Wang, J. Esteva, E. Bertran, Thin Solid Films 3171-2 (1998) 6.[189] X. Jiang, R.Q. Zhang, G. Yu, S.T. Lee, Phys. Rev. B, to be published.[190] Wei Zhu, Hua-Shuangn Kong, Jeffrey T. Glass, Characterization of diamond films, in: Robert F. Davis (Ed.),

Diamond Films and Coatings, Noyes Publications, Park Ridge, NJ, USA, pp. 244±338.[191] Hiroshi Kawarada, Surface Science Reports 26 (1996) 205.[192] B.J. Waclawski, D.T. Pierce, N. Swanson, R.J. Celotta, J. Vac. Sci. Technol. 21 (1982) 368.[193] S.-T. Lee, G. Apai, Phys. Rev. B. 48 (1993) 2684.[194] T. Aizawa, T. Ando, M. Kamo, Y. Sato, Phys. Rev. B 48 (1993) 18348.


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