Nanocrystalline Diamond

Encyclopedia of Nanoscience and Nanotechnology www.aspbs.com/enn Nanocrystalline Diamond Narendra B. Dahotre The Unive...

Author: Dahotre N.B. | Kichambare P.D.

29 downloads 782 Views 617KB Size Report

This content was uploaded by our users and we assume good faith they have the permission to share this book. If you own the copyright to this book and it is wrongfully on our website, we offer a simple DMCA procedure to remove your content from our site. Start by pressing the button below!

Report copyright / DMCA form

DOWNLOAD PDF

Encyclopedia of Nanoscience and Nanotechnology

www.aspbs.com/enn

Nanocrystalline Diamond Narendra B. Dahotre The University of Tennessee, Knoxville, Tennessee, USA

Padmakar D. Kichambare The University of Kentucky, Lexington, Kentucky, USA

CONTENTS 1. Introduction 2. Methods of Preparation 3. Nucleation and Growth Kinetics 4. Spectroscopic Characterization 5. Physical Properties 6. Applications 7. Summary Glossary References

1. INTRODUCTION Polycrystalline diamond (PCD) ﬁlm has excellent mechanical and tribological properties such as extreme hardness, chemical inertness, and high electric resistance [1–3]. Hence there is considerable interest in the synthesis of diamond ﬁlm by various chemical vapor deposition techniques for wide range of applications from hard coating to electron emitting surfaces for ﬂat panel display devices. However, the surface of these ﬁlms is often very rough due to the polycrystalline nature with crystallite sizes of the order of micrometers. Such a rough surface causes excessive light scattering, reducing the optical transparency, causing high friction and wear losses on mating surface, and thereby limiting its use for various applications. Hence attempts were made to develop effective chemical etching and mechanical polishing techniques to get smooth diamond ﬁlms. Because of chemical inertness and extreme hardness of diamond ﬁlms, these techniques are not efﬁcient and cost effective. Since one can change neither the microstructure nor the composition of diamond, the best way to get these desired properties will be to control the microstructure of diamond ﬁlms [4].

ISBN: 1-58883-062-4/$35.00 Copyright © 2004 by American Scientiﬁc Publishers All rights of reproduction in any form reserved.

Recently the nanocrystalline diamond (NCD) ﬁlms have been recognized as an important class of structure in the family of diamond materials [4, 5]. It has been investigated that the microstructure of diamond ﬁlms could be controlled so that crystallite size ranges from micrometer to nanometer. The size can be controlled continuously over this range by changing the gas phase chemistry of the plasma enhanced chemical vapor deposition technique. The enhancement in mechanical, electrical, and optical properties of NCD ﬁlms over the PCD ﬁlm has been achieved. This is due to the reduction in crystallite size of diamond particles. The decrease in crystallite size increases vastly the number of grain boundaries and the entire ﬁlm becomes electrically conducting. Second, the reduction in crystallite size increases the smoothness of the NCD ﬁlms. Such ultrasmooth ﬁlms not only provide an opportunity for tribological applications but also their wide bandgap makes them suitable to act as ideal transparent ﬁlm for optical components. The lower friction coefﬁcient and higher hardness makes NCD a better material for coating mechanical tools. Thus controlling the microstructure of diamond ﬁlms could bring many remarkable properties for various industrial applications [4]. There are a few articles related to the preparation of NCD ﬁlms and several reporting the unique properties of NCD. A number of different deposition techniques and conditions have been used to get smooth ﬁlms that are summarized below. Gruen and co-workers reported the nucleation and growth of NCD ﬁlm on silicon from hydrogen-poor argon plasma using fullerene as growth precursors [6, 7]. Deposition of NCD ﬁlm was performed by introducing C60 into an argon microwave discharge in contact with a silicon wafer at 750 C. The growth of NCD ﬁlm was carried out on the substrates that were previously mechanically scratched with 0.1 m diamond powder. The deposition of NCD ﬁlms was carried out with total argon pressure of 98 Torr, C60 partial pressure of 10−2 Torr, ﬂow rate of 100 sccm, and

Encyclopedia of Nanoscience and Nanotechnology Edited by H. S. Nalwa Volume 6: Pages (435–463)

436 microwave power of 800 W. It is difﬁcult to get stable argon plasma in absence of ∼1% added hydrogen. The presence of 1 to 3% hydrogen in the carbon containing argon plasma does not appreciably affect the properties of the diamond ﬁlms. The phase pure diamond ﬁlms have been grown from C60 /Ar microwave plasmas and these ﬁlms have nanocrystalline microstructure. Zhou and co-workers prepared NCD ﬁlms in an Ar–CH4 microwave discharge without addition of molecular hydrogen [8]. N -type single crystal silicon wafers with (100) orientation were used as the substrate, and mechanical polishing with ﬁne diamond powder (0.1 m) was employed to enhance the nucleation density. The NCD ﬁlms consist of a pure NCD phase with grain sizes ranging from 3 to 20 nm. The ﬁlms have very smooth surfaces, which do not depend on the thickness of the ﬁlm at least in the range of 1–5 m. Erz et al. grew 800 nm thick NCD ﬁlms with up to 10% CH4 in H2 [9]. Characterization was done with X-ray diffraction (XRD), Raman, and visible infrared spectroscopy. Incorporated diamondlike carbon (DLC) degrades the transmissivity, which in some samples reaches 93%. Meanwhile Konov et al. and Nistor et al. deposited 0.2– 1 m thick ﬁlms from CH4 /H2 /Ar mixtures that were characterized with XRD, high-resolution transmission electron microscopy (HRTEM), and Raman spectroscopy [10–13]. The percentage of Ar in the mixtures was always kept at 50%, while the CH4 /CH4 + H2 ratio was varied. Grain sizes were found to be in the range 30–50 nm. TEM did not reveal amorphous carbon in signiﬁcant quantities, but disordered sp3 and sp2 bonded amorphous carbon was detected, presumably located at grain boundaries. The diamond crystallites are highly defected with many twins and other planer defects. Gu and Jiang prepared NCD ﬁlms by continuous H+ ion bombardment of different energies induced by applying a negative substrate bias voltage in microwave plasma assisted chemical vapor deposition (CVD) [14]. The growth of NCD ﬁlms was carried out on mirror polished n-type Si (001) wafers with a diameter of 50 mm and thickness of 300 m. They used a three-step process. In the ﬁrst step, the Si wafer was etched in-situ in hydrogen plasma for 40 min at 860 C. The second step involves bias pretreatment for the ﬁlm nucleation at 845 C in 5% methane in hydrogen plasma at 20 mbar and 840 W power, while the negative bias voltage was kept at −150 V relative to the vacuum chamber, which was electrically grounded. In the last step, H+ ion bombardment assisted growth of diamond ﬁlms was performed at different negative dc bias voltages (0–140 V), substrate temperatures (730–900 C), and total pressures of 15–50 mbar. The nucleation density was about 1010 cm−2 after 17 min. A series of NCD ﬁlms was grown with grain size ranging from 4 nm to a few hundreds of nanometers using a microwave plasma enhanced chemical vapor deposition (MPCVD) technique by Chen et al. [5]. The deposition of NCD ﬁlms was carried out on ultrasonically polished quartz substrates in a 5 kW microwave reactor. A mixture of semiconductor grade CH4 , H2 , and O2 was used as source gas, wherein the fraction of methane was varied between 4 and 42%, while O2 was always kept constant at 0.1%. The total ﬂow rate was 200 sccm, the chamber pressure was 22 Torr,

Nanocrystalline Diamond

the microwave power was 1 kW, and the substrate temperature was maintained at 590–600 C. Effects of the substrate pretreatment and the methane fraction in the source gas on the microstructure, surface roughness, and optical transmittance of the NCD ﬁlms were studied. Speciﬁcally, comparison was made between two different sizes, 4 nm and 0.1 m, of the diamond powder used for substrate pretreatment. Interestingly, the ﬁlms grown on substrates scratched with coarser powder (0.1 m) are smoother and more transparent than those on substrates scratched with ﬁner powder (4 nm) despite the similarity in the grain size of these two types of ﬁlms prepared at high methane fractions. It is also observed that the major factor that controls the optical transparency is the surface roughness irrespective of the grain size as long as the sp2 -bonded carbon in the ﬁlm is avoided. Sharda et al. have grown the NCD ﬁlms by biased enhanced growth in a MPCVD system on mirror polished Si (100) [15, 16]. No diamond powder or any other ex-situ treatment was performed prior to the depositions. A mixture of 5% CH4 in H2 was used at a pressure of 30 Torr with a microwave power of 1000 W for deposition of NCD ﬁlms. The whole growth was performed for 1 hour in a single stage while applying a negative dc bias voltage of 260 V to the Si substrates with respect to the chamber that was grounded. Hong et al. prepared NCD ﬁlms on Si substrate from CH4 , H2 , and O2 gas mixture by microwave plasma enhanced CVD and performed real time spectroscopic ellipsometry to understand the growth process of NCD ﬁlms [17]. This study emphasizes the broad utility of the ellipsometric measurements to obtain accurate substrate surface temperature and characterization of the initial substrate damage that occurs upon seeding, the annealing of the substrate damage that occurs upon heating to the growth temperature, and the structure of the diamond ﬁlm during the nucleation and bulk ﬁlm growth regimes. The effect of methane pressure on the ﬁlm growth was studied by Fedoseev et al. [18]. The ﬁlms deposited by Zarrabian et al. from an electron cyclotron resonance (ECR) plasma were found to consist of 4–30 nm crystallites embedded in DLC [19]. Magnetron sputtering of vitreous carbon leads to the nanocrystallites of 55 to 75 nm in size that were embedded in an amorphous carbon matrix. The sputtering was carried in Ar with 0–10% added H2 . Amaratunga et al. reported mixed phases using ﬁlms containing 10–200 nm diamond crystallites embedded in a nondiamond carbon matrix [20]. In another report, ultrananocrystalline diamond ﬁlms with up to 0.2% total nitrogen content were deposited by a MPCVD method using a CH4 (1%)/Ar gas mixture and 1–20% nitrogen gas added. An enhancement in electrical conductivity of nitrogen doped NCD by ﬁve orders of magnitude was observed. The surface morphologies of as grown NCD ﬁlms have been studied by atomic force microscopy (AFM) [21]. Three-dimensional AFM images of the ﬁlms with different thicknesses of 1 and 5 m illustrate that the NCD ﬁlms produced from the Ar–CH4 plasma have very smooth surfaces. The surface roughnesses measured over an area of 5 m2 for these two ﬁlms are 36.5 and 38.6 nm, respectively, suggesting thereby that the surface roughness of NCD ﬁlms is largely insensitive to the ﬁlm thickness, which is in

437

Nanocrystalline Diamond

contrast to the conventional CVD PCD ﬁlms prepared by an atomic hydrogen-rich microwave plasma. It is observed that the NCD ﬁlm hardness depends on the methane content used for diamond deposition. The hardness of these ﬁlms falls in the narrow range from 73 to 85 GPa that is lower by ∼20% than the value reported for the hardest direction (100) in natural diamond stones. A combination of high hardness and low surface roughness of the NCD ﬁlms suggests the ﬁlms to be perspective coatings for tribological applications. The transparency of NCD ﬁlm is strongly dependent not only on the methane fraction but also on the pretreatment of the substrate. Optical transparency over 60% in the spectral range of 0.6–2.0 m is considered sufﬁciently high for most of optical applications. The NCD ﬁlms also exhibit remarkable electron emission behavior. A high current density of 4 × 10−4 amps/cm2 with turn-on ﬁeld of ∼1 V/m was reported for NCD ﬁlms grown from C60 (80%)–Ar(20%)–H2 microwave plasmas. Similarly these NCD ﬁlms were successfully deposited on sharp Si microtip emitters in order to enhance the electron emission properties of bare Si microtips. The Si tip emitter coated with 0.1 m thick NCD demonstrates a substantial reduction in the threshold ﬁeld and an increase in the maximum emission current in comparison to the uncoated Si tip. The NCD ﬁlms produced from C60 –Ar plasma have several interesting electrochemical properties. These ﬁlms have a wide working potential window, a low voltametric background current, and a high degree of electrochemical activity for several inorganic redox systems.

2. METHODS OF PREPARATION In the conventional diamond ﬁlms grown by CVD, the gas mixture of hydrocarbon–hydrogen is used. Typically NCD deposition requires a gas mixture of 1% CH4 in 99% H2 . It has been thought that atomic hydrogen is an absolutely essential ingredient of vapor from which the diamond ﬁlms are grown. Atomic hydrogen is also generated by thermal decomposition or by collision processes in the plasma. There are various methods of preparation of NCD ﬁlms among which plasma CVD and magnetoactive CVD are worth mentioning and most attractive. Various hydrocarbon precursors are available for the CVD processes of diamond ﬁlms, among which the precursor having methyl group is favored for the preparation of diamond ﬁlms. Methane (CH4 , which dissociates into CH3 in plasma and was preferred over other precursors such as CH4 with hydrogen, produced diamond easily compared to other hydrocarbon precursors.

the sublimation of graphite. It is produced in microwave plasmas under nonequilibrium conditions from hydrocarbon molecules, resulting in supersaturated carbon vapor. With CH4 as the precursor, the C2 stationary concentration in argon microwave plasmas is higher than that of saturated carbon vapor by many orders of magnitude at the neutral Ar gas temperature 1600 K. At this temperature, CH4 thermally decomposes to C2 H2 . The fraction of C2 H2 is then converted into C2 during plasma processes by nonequilibrium processes. The dimer, C2 , has a gas kinetic pseudoﬁrst-order rate reaction with both CH4 and C2 H2 and leads to likely formation of methylacetylene [24] and cyclobutyne, respectively, as product. This high probability of formation of these products of reaction and successively larger hydrocarbon molecules is regarded as a process of homogeneous nucleation. As a result, there is a deposition of graphitic carbon. On the other hand, if the goal of the experiment is to deposit the diamond, that could be achieved by suppressing a gas-phase reaction of the sort described for heterogeneous nucleation on the diamond nuclei initially placed on the substrate ex-situ. Hence, there is a low probability of reactions by C2 to form higher carbon clusters. Under these conditions, diamond could nucleate and grow using carbon dimer. The growth of NCD ﬁlms consisted of introducing C60 into an argon microwave plasma at 750 C on Si substrates [6, 7]. The Si substrates were mechanically treated with 0.1 m diamond powder. The deposition was performed for several hours with total argon pressure of 98 Torr, C60 partial pressure of 10−3 Torr, ﬂow rate of 100 sccm, and microwave power of 800 W. It is difﬁcult to operate and get the stable microwave plasma in the absence of 1% added hydrogen. The presence of 1–3% hydrogen in the carbon containing ﬁlm does not appreciably affect the properties of diamond ﬁlms, which are different from conventionally grown ﬁlm with 1% CH4 and 99% H2 .

2.2. MPCVD NCD ﬁlms have been deposited on the quartz substrates by the MPCVD method [5]. A schematic diagram of the system employed during the work is shown in Figure 1 and consists

2.1. From Fullerene-Argon Microwave Plasma The production in large quantities of the fullerenes [22] provides an opportunity to deposit diamond ﬁlms from noble gas plasmas with very low hydrogen contents using C60 as a precursor. The fragmentation mechanism of C60 and kinetics of the process have been studied in detail [23] and it is observed that fragmentation of C60 by photon, collisions with surfaces, or gas-phase collisions occurs primarily by the elimination of carbon dimer, C2 . This is a highly reactive chemical species and is one of the carbon species resulting from

Figure 1. Schematic diagram of microwave plasma enhanced chemical vapor deposition chamber.

438 of a stainless steel four-way deposition vacuum chamber, a graphite holder that is attached to an induction heater, 5 kW microwave power supply, various gas lines connected to a vacuum chamber, and a vacuum pumping system. A mixture of semiconductor grade CH4 , H2 , and O2 was used as a source gas, wherein the fraction of CH4 was varied between 4% and 42% and O2 was always kept constant at 0.1%. The total ﬂow rate was 200 sccm, the chamber pressure was 22 Torr, the microwave power was 1 kW, and the substrate temperature was maintained at 600 C during deposition. An independent radio-frequency (rf) heater was used to control the substrate temperature. During deposition, a single-color optical pyrometer was used to record the temperature from the front surface of the substrate. The smooth ﬁlms were deposited on the pretreated quartz substrates. The substrate pretreatment could be carried by ﬁrst chemically cleaning the quartz plates with acetone and doubly distilled water and then these plates could be used for the ultrasonic polishing with diamond powder having grain sizes 4 nm and 0.1 m. The ultrasonic polishing of the quartz substrates is simple. This involves the preparation of suspension of diamond powder in acetone, placing the substrates in this suspension, and subjecting the beaker to ultrasonic vibration for at least 8 h.

2.3. Biased Enhanced MPCVD The growth of diamond is considered to be taking place via surface processes of addition and subtraction of radicals from the gas phase [25, 26]. During the growth of diamond, ions enhance lattice disorders and also promote graphitic content in the deposit. On the other hand, in the growth of carbon ﬁlms by hyperthermal species, ions with energy in the range 40–200 eV produce a high concentration of sp3 carbon in the ﬁlm [27–30]. In this case, the growth is controlled by a subplantation mechanism, which relies on the shallow subsurface implantation of the carbon ions. Both of these routes of growth of carbon ﬁlms are totally different. In the bias enhanced route the NCD ﬁlms are deposited in a CVD diamond environment while continuously biasing the substrate to take advantage of both the processes in the growth. The CVD diamond conditions result in high etching rates of the nondiamond carbon from the surface processes while continuous biasing takes advantage of subsurface phenomena. This route is simply an extension of bias enhanced nucleation of CVD diamond and results in the growth of hard and smooth NCD ﬁlms. The NCD ﬁlms could be grown in a MPCVD reactor on mirror polished silicon substrate placed on an Mo holder [15]. No diamond powder or any other ex-situ treatment is required prior to the deposition. The substrate is then immersed in methane and hydrogen plasma. A mixture of 5% CH4 in H2 should be used at a pressure of 30 Torr, with a microwave power of 1 kW for all the ﬁlms. A negative dc bias voltage of 260 V should be applied to the substrate with respect to the chamber that is grounded. The whole growth of NCD ﬁlm could be achieved in a single stage and requires a 1 hour duration without breaking the bias to the substrate. While holding the bias voltage constant throughout the deposition, bias current could be varied with time in the same fashion as commonly observed in biased

Nanocrystalline Diamond

enhanced nucleation in the growth of CVD diamond; the current increases after some incubation period followed by saturation at longer deposition time. It is observed that there is an increase in the current attributed to enhancement in electron emission from the surface as highly emissive diamond is deposited on silicon.

2.4. Ar–CH4 MPCVD Atomic hydrogen plays a crucial role in the growth of PCD ﬁlms by CVD using hydrocarbon as a carbon source [31, 32]. It is well known that atomic hydrogen from a hydrogen-rich reactant gas can terminate the carbon dangling bonds with a tetrahedral sp3 conﬁguration and etch out nondiamond materials at the growth surface of diamond. The reduction in concentration of atomic hydrogen causes the growth of nondiamond phase or no diamond ﬁlm deposition. Hence, argon has been used in place of hydrogen. A mixture of Ar (99 sccm) and CH4 (1 sccm) was used as the reactant gas for growth of NCD ﬁlm. N -type single crystal silicon (100) wafers were used as substrates. These substrates were polished mechanically with 0.1 m diamond powder. During the entire ﬁlm growth, the substrate temperature and input microwave power were 800 C and 800 W, respectively. The ﬁlms prepared in this work were 1 or 5 m thick. As grown ﬁlms prepared from the Ar–CH4 plasma, they consist of phase-pure crystalline diamond grains ranging from 3 to 20 nm in size independent of the NCD ﬁlm thickness [21].

2.5. Hot Filament CVD It has been shown that the addition of Ar into the CH4 – H2 or C60 MPCVD can provide a route for controlling the microstructure of the diamond ﬁlm, leading to thick and smooth NCD ﬁlms at an Ar concentration of 90% [33, 34]. Due to the very high concentration of C2 species at higher percentage of Ar, it has been concluded that C2 dimer may be the growth precursor for NCD ﬁlms in the MPCVD system. While most of studies on growth of NCD are focused on using MPCVD, there are only few reports on the use of the hot-ﬁlament CVD (HFCVD) system for the growth of NCD ﬁlms. An interesting question is whether the chemistry observed in the MPCVD system for the growth of NCD ﬁlms can be extended to HFCVD as the mechanism for gas activation is different in both these methods. Similarly the fraction of argon ions and electron density is signiﬁcantly lower in the HFCVD system than the MPCVD system. Lin et al. employed a conventional HFCVD system using curved tantalum wire arrays for the growth of NCD ﬁlms [35]. A CH4 /H2 gas feed ratio that can yield the growth of high quality, well-faceted diamond ﬁlm was used. Ar was then introduced in increasing concentration to the system in order to investigate the changes in microcrystallinity of the ﬁlm as a function of ArH2 ratio. The H2 and CH4 ﬂow rates were ﬁxed at 0.05 and 0.0013 l/min, respectively. The Ar/H2 ratio was adjusted by increasing the Ar ﬂow rate from 0 to 0.43 l/min. In order to get a higher ArH2 ratio, the H2 ﬂow rate was reduced. The total pressure was kept at 60 Torr. The substrate temperature was held at 950 C for Ar concentration below 92%. When higher Ar concentration was used, the substrate temperature was reduced to 870 C to

439

Nanocrystalline Diamond

avoid the deposition of graphite. The Si (100) wafer used in this work was supersonically scratched with diamond powder prior to the deposition of diamond ﬁlms. The range of Ar concentration in the Ar–CH4 –H2 mixture that permits wellfaceted diamond growth is up to 90%. At a concentration of 95.5%, there is a marked transition into NCD phase without graphitic phases. Such compositional mapping is one of the effective routes to prepare thick and smooth NCD ﬁlms.

before entering the deposition chamber. NCD ﬁlms were deposited at 1.5 kV with 75 mA current at a chamber pressure of ∼0.3 mbar on glass, quartz, and silicon substrates. The substrate temperature was varied within 573–713 K depending on the substrate material. The value of Ts for glass substrate was within 573–623 K, while for quartz it was within 573–713 K. The substrates could be heated to the desired temperature inside the substrate holder with proper insulation.

2.6. CVD of Camphor It is well established that the methyl radical along with atomic hydrogen is key to growing the diamond ﬁlms. A wide variety of precursor materials are available for the CVD processes of diamond ﬁlms. Although any carbon containing precursor material may be used in the CVD technique, the preparation of diamond ﬁlm is favored by the presence of a methyl group [36, 37]. Hence, methane, which dissociates into CH3 in plasma, has been preferred over other precursor materials because CH4 with hydrogen produced the diamond easily. Similarly the organic compounds, which can easily generate methyl radicals, may also be suitable for the growth of diamond ﬁlms. Camphor (C10 H16 O) as the precursor material has been mentioned for the growth of NCD ﬁlms. In camphor, there are three carbon atoms attached to three methyl groups while the remaining seven carbon atoms are associated with a ring structure. NCD ﬁlms were deposited by CVD of camphor and hydrogen (∼75%) [38]. Camphor is a natural source which sublimes at room temperature and is soluble in alcohol, ether, acetone, and benzene. Figure 2 [38] depicts the schematic diagram of the apparatus used for CVD of camphor. The camphor vapor was passed through a heated stainless steel tube into a cylindrical deposition vacuum chamber. The heating tape was used over the stainless steel tube in order to prevent condensation of camphor vapor

Figure 2. Schematic diagram of the deposition system. Reprinted with permission from [38], K. Chakrabarti et al., Diamond Relat. Mater. 7, 845 (1998). © 1998, Elsevier Science.

3. NUCLEATION AND GROWTH KINETICS The preparation of pure and compact diamond thin ﬁlms by MPCVD from CH4 , H2 , and O2 gas mixture on nondiamond surfaces has a great potential for various commercial applications [39–41]. Such thin ﬁlms can endow solid surfaces with one or more of the superlative properties of diamond, including resistance to wear and chemical attack, optical as well as infrared transparency, and high thermal conductivity. The limitations in achieving these prospective properties include: (a) the high substrate temperatures required for diamond nucleation with reasonable growth rates, (b) the surface pretreatments like abrasion with diamond powder for high nucleation density, and (c) the defective nature of the resulting nano- and/or PCD ﬁlm. Diamond ﬁlm deposited on nondiamond substrates exhibits a variety of difﬁculties related to their nano/polycrystalline structure, mostly in the appearance of a second phase of sp2 bonded C atoms. The transmission electron microscopy and Raman spectroscopy studies performed on the underside of a 20 m diamond ﬁlm removed from its substrates revealed that diamond clusters are surrounded by sp2 -bonded C phase in the form of graphite crystals [42], while the growth side of these ﬁlms exhibited the pure diamond phase. The formation of a defective region near the substrate interface has been attributed to heteroepitaxy [42] or to impurities that arise from the interaction of the plasma with the substrate during the nucleation stage [43]. Previous studies have revealed a 0.7 to 0.8 eV optical gap for the graphitic component of thin (0.2 m) diamond ﬁlms [44] consistent with the existence of sp2 bonded C atoms in clustered units sufﬁciently small so that a gap has developed [45]. A dominant graphite phase has also been detected in the grain boundary regions for 30–40 m thick diamond ﬁlms prepared with excess CH4 in the CH4 /H2 mixture. On the other hand, in the ﬁlms prepared with an optimum CH4 /H2 mixture, this phase is no longer dominant and contains dilute sp2 C atoms within the diamond grains [46]. These pioneering studies of diamond thin ﬁlms demonstrate that the atom fraction, physical location (intragrain or grain boundary), and nature (clustered versus dilute) of the dominant sp2 bonded C are sensitive parameters for the evolutionary stage of the ﬁlm and the preparation conditions. The characteristics of the sp2 C phase and the void network are of greatest interest because such bonding and structural defects lead to degradation in the mechanical, optical, and electrical properties of diamond thin ﬁlms deposited on nondiamond substrates. In view of this, real time monitoring of the deposition process becomes increasingly important in

440

3.1. Real Time Spectroscopic Ellipsometry Spectroscopic ellipsometry (SE) and optical emission spectroscopy (OES) are powerful, real time optical probes for characterizing ﬁlm growth. Ellipsometry has an advantage over other optical techniques for characterizing the evolution of diamond thin ﬁlms because it can provide the volume fractions of diamond, sp2 C, and void with good sensitivity in the ﬁlms as thin as 40 Å [51, 52]. Hong et al. have applied real time SE to characterize the preparation of ∼2000 Å thick NCD ﬁlm on Si substrate by MPCVD [53]. Si wafers were seeded and placed on a graphite substrate platform. The platform is rf heated to a temperature that can be controlled independently of microwave power over a relatively wide range. In order to achieve a stable plasma conﬁguration, a pure H2 plasma was ignited with the substrate platform about 2 cm above the optical plane. The substrate platform was adjusted until optical alignment was achieved. CH4 and O2 were introduced, deﬁning time zero in the data collections. The spectroscopic ellipsometer employs a polarizer rotating at 15.6 Hz that modulates the polarization state of an incident, collimated, white light beam generated by an Xe source. A ﬁxed analyzer and a photodiode array based detection system should be used for the analysis of the polarization state change induced by the growing ﬁlm over a range of photon energies [54, 55]. The resulting spectra consist of ∼110 pairs of ellipsometry angles ( , ) from 1.5 to 4.5 eV and were obtained as an average of 40 rotations in an acquisition time of 2.6 s. In this time, ∼1 Å accumulates at the maximum bulk layer deposition rate of 0.4 Å/s. The repetition time for the collection of successive spectra is adjusted to achieve sufﬁcient thickness resolution for the growth rate measurements. The ( , ) spectra were analyzed by leastsquares regression [56]. The nuclei studied in the ﬁrst 5 Å of mass thickness dm are simulated using a simpliﬁed single layer model consisting of a composite of diamond and sp2 carbon with void. The optical properties of this layer are calculated with the Bruggeman effective medium approximation (EMA) using a ﬁxed nuclei composition established after dm ∼ 50 Å. The outcome of the analysis in the ﬁrst 5 Å is the physical thickness d, from which dm = fd d can be determined, where fd is the diamond volume fraction. The continuous ﬁlm is simulated as a two layer, roughness/bulk combination, where the surface roughness is modeled with the EMA as a 0.5/0.5 mixture of diamond/void, and the bulk layer is modeled as a composite of diamond, sp2 C, and void. The output of the analysis includes the bulk and roughness layer thickness, db and ds , and the bulk composition, from which dm = fab db + 05ds , where fab is the volume fraction of diamond in the bulk. The dm versus time in the ﬁrst 5 Å of the nucleation regime and in the bulk ﬁlm growth regime

shows a linear ﬁt that yields the deposition rates in terms of mass thickness. The mass thickness deposition rates are then plotted vs T −1 from 390 to 840 C in Figure 3 [57]. A well-deﬁned activation energy of 8 ± 05 kcal/mol has been noted for diamond growth for 390 C ≤ T ≤ 800 C in the bulk regime. Above 800 C, a deviation from activated behavior could be observed, characterized by a maximum in the rate near 800 C followed by a decrease with increasing T , and is attributed to a reduction in the incorporation probability. It should be noted that the 8 kcal/mol activation energy deduced in the bulk growth regime is close to the values obtained for diamond homoepitaxy on (100) and (111) surfaces [57]. This activation energy is attributed to sp3 C atom incorporation in speciﬁc atomic sites on these surfaces. Figure 3 shows larger scatter for individual depositions in the nucleation regime due to substrate seeding variation. During nucleation, the rates are 10–102 times lower than in the bulk regime with an apparent activation energy of 17 ± 2 kcal/mol. The lower rates arise from growth at discrete nucleation centers, in which case only Ccontaining precursors that arrive within ∼100 Å of a center can be incorporated. The higher activation energy is attributed to a decrease in the capture area as the nucleation density and precursor surface diffusion length decrease with decreasing T . Figure 3 gives a clear understanding of the limitation of diamond ﬁlm growth at low temperature. Lee et al. [57] suggested that the apparent activation energy of ∼17 kcal/mol in the nucleation regime is a characteristic of the seeding process. Higher activation energies are expected with less effective seeding, and this explains values as high as ∼30 kcal/mol [58, 59]. Figure 3 also demonstrates that the nuclei growth rate at T = 400 C is ∼103 times lower than the bulk rate at 800 C. Of this factor of 103 , one order is attributed to the decrease in the average incorporation 102 Bulk Ea = 8 kcal/mol 101

Rm (Å/min)

order to quantify these characteristics as a function of ﬁlm thickness and to exert greater control over the growth process for improved near-interface and bulk properties. Optical probes are beneﬁcial in this role as they are passive and require no instrumentation internal to the reactor [47–50] because of the adverse environment of diamond deposition, highly reactive gases, high gas pressures, and high temperatures.

Nanocrystalline Diamond

100

Nuclei Ea = 17 kcal/mol

10-1

10-2

IL-6281

0.8

1.0

1.2 1.4 103 /Tellips (K-1 )

1.6

Figure 3. Diamond growth rates in terms of mass thickness vs inverse temperature. The lines are the best linear ﬁts for T > 800 C in the bulk growth regime, and to the full data set in the nucleation regime. The slopes of these lines yield the activation energies as shown. The broken line indicates a reproducible deviation from activated behavior. Reprinted with permission from [17], B. Hong et al., Diamond Relat. Mater. 6, 55 (1997). ©1997, Elsevier Science.

441

Nanocrystalline Diamond

3.2. Optical Emission Spectroscopy The diamond ﬁlms were deposited on pretreated quartz substrates by the MPCVD technique as described in detail elsewhere [5]. In order to obtain information on plasma chemistry, OES was employed to monitor the H2 –CH4 plasma. A monochromater with a focal length of 500 mm and grating of 1200/mm, which gives a resolution of 0.05 nm and a precision of about ±02 nm/500 nm, was used for the OES measurements. Figure 4 [5] presents some typical optical emission spectra of the H2 –CH4 plasma used for the NCD ﬁlm deposition. The H2 spectrum consists mainly of emission from atomic H Balmer lines (H and H lines at 656.3 and 486.1 nm, respectively) and excited H2 lines at around 460 and 590 nm. When methane was introduced into the system, C3 emission lines at around 400 nm, CH+ emission at 417 nm, CH lines at around 431 nm, and the C2 Swan bands around 469 and 515 nm appeared in the emission spectrum. Using the H2 line at 580.6 nm as reference, the relative intensities of emissions from the latter four species as a function of the methane fractions in the source gas are plotted in Figure 5 [5]. All emissions except that of C2

4500 4000

Intensity (a.u.)

3500 3000 2500

37.5% CH4

2000 1500

4% CH4

1000 500 0

H2

300

400 500 600 Wavelength (nm)

IL-6301

700

Figure 4. Some typical optical emission spectra of the H2 /CH4 plasma. Reprinted with permission from [5], L. C. Chen et al., J. Appl. Phys. 89, 753 (2001). © 2001, American Institute of Physics.

1.3 1.2 1.1

Relative Intensity

probability from 800 to 400 C, another order is due to the fraction of the substrate area from which initial nuclei can collect impinging ﬂux, and factors of 4 and 2 are due to the reduction in the nucleus capture radius and nucleation density, respectively. Thus considerable gains can be achieved in preparing diamond at reasonable rates on nondiamond substrates at low T by enhancing the density and area of the nucleation centers and retaining these characteristics as T is reduced. It is most probable that disordered sp2 or sp3 C embedded within the silicon during seeding is converted to NCD nuclei by H exposure in the initial stages of nucleation and that this process becomes more effective as T increases. Although gains in diamond ﬁlm can be made at low T by increasing the precursor ﬂux, a primary limitation is the low capture cross-sectional area of nuclei at low density on the substrate.

CH*/H2 C2/H2 CH/H2 C3/H2

1.0 0.9 0.8 0.7 0.6 0.5 0.4

0

10

20 CH4 (%)

30

40

IL-6282

Figure 5. The relative intensity of the CH+ , C2 , CH •, and C3 + × emissions with respect to that of H2 as a function of the methane fraction in the source gas. Reprinted with permission from [5], L. C. Chen et al., J. Appl. Phys. 89, 753 (2001). © 2001, American Institute of Physics.

dimer increased with increasing methane content, reached a maximum at around 20%, and decreased at higher methane content. Emission from C2 dimer, however, showed a continuing increase above 20% methane content. It can be seen from Figure 4 that the 515 nm C2 emission line was very pronounced at high methane content. Due to the fact that the emission intensity is a function of many factors including Frank–Condon factor, the density of states, various quenching effects, etc., a quantitative description of the observation is very complicated and is beyond the scope of this chapter. The gas-phase analysis by OES or any other single technique is far from complete since it is not possible to detect some important growth species such as CH3 by this technique. It is observed from these studies that the gas phase experienced substantial change when the methane fraction went from below to above 20%, with which the optical transmittance and surface roughness data [5] also exhibited similar crossover behavior. Shown in Figure 6 [5] is the intensity ratio of two atomic hydrogen lines at 486.1 nm (H ) and 656.3 nm (H ) as a function of the methane fraction in the source gas. It is quite clear from the graph that the intensity ratio of H /H , which is related to the electronic temperature in the plasma, showed a maximum value at a methane fraction of 20% and declined rapidly with variation in methane content on either side. Once again, the electronic temperature exhibited a methane-fraction dependence similar to that of most emissions (except C2 as shown in Figure 5. From Figures 5 and 6, it is speculated that a change in major growth species might have occurred around a methane fraction of 20%. Most of the conventional PCD ﬁlms were grown in low methane fraction. The major growth species C2 has been identiﬁed for the growth of NCD ﬁlms using C60 as precursor under hydrogen-free or hydrogen-deﬁcient environments [60]. For the conditions during this work [5], growth involving CH3 was expected since all depositions were carried out

442

Nanocrystalline Diamond

1.0

0.50 0.45 0.40 0.35

0.8

0.8

0.6

0.6 C2

0.4

0.4 C

0.2

0.2 OH (a)

0

0.30

0

10

20 30 CH4 (%)

40

Figure 6. The relative intensity of H /H as a function of the methane fraction in the source gas. Reprinted with permission from [5], L. C. Chen et al., J. Appl. Phys. 89, 753 (2001). © 2001, American Institute of Physics.

in H2 –CH4 plasma. However, it is clear from OES studies that C2 dimer becomes the predominant growth species at methane fractions far above 20%.

3.3. Vacuum Ultraviolet Absorption Spectroscopy The deposition of NCD ﬁlm by using a low-pressure inductively coupled plasma (ICP) was employed by Teii et al. [61]. NCD ﬁlms were grown at a pressure of 80 mTorr using CO/CH4 /H2 and O2 /CH4 /H2 gas mixtures. The densities of C atoms in ground state are measured by vacuum ultraviolet absorption spectroscopy (VUVAS). CO is employed as one of the simple gases that yield C atoms, through the primary reaction: CO + e− → C + O + e− . The emissive species observed in the CO/CH4 /H2 plasmas in the wavelength range of 300–800 nm were OH, CH, H, C2 , and Ar. In addition to these species, pronounced emissions from O were observed in the O2 /CH4 /H2 plasmas. The normalized emission intensity of OH, CH, H , C2 , and O [62, 63] and the normalized absorption of C atoms as a function of XCO and XO2 for a ﬁxed XCH4 of 5% are depicted in Figure 7 [61], respectively. It was very difﬁcult to obtain sufﬁcient absorption by C atoms when an emission line at 296.7 nm was initially used in (UVAS). It therefore follows from this that the absolute C atom density was lower than 1 × 1012 cm−3 , the lower detection limit in this UVAS system [64]. The trend of the C atom density as determined by VUVAS was consistent with that by OES under experimental conditions, indicating thereby that the inﬂuence of collisional excitation processes was small. In Figure 7a, the H and CH intensity ratios remain constant independent of XCO , while the C atom absorption and OH and C2 intensity ratios increase in proportion to XCO . To ensure that an increase in C and C2 was not due to an overall increase in the total carbon concentration in the feeding gas, a comparative experiment was performed by varying XCO with the sum of XCO + XCH4 kept constant at 5%. This experiment gave similar results except for a slight decrease in the CH

1.0

Normalized emission intensity ratio

0

IL-6285

0 2 4 CO concentration (%)

IL-6283

1.0 Hα

0.8

0.6

0.8

CH C2

0.6

C 0.4

0.4 O ( ) 0.2

0.2 OH

Normalized absorption of C atoms

Relative Intensity Hβ/Hα

0.55

1.0 CH Hα

Normalized absorption of C atoms

Normalized emission intensity ratio

0.60

(b)

0

0 IL-6 28 6

0 2 4 O2 concentration (%)

Figure 7. Variations of the normalized emission intensity ratios of OH, CH, H , C2 , O, and the normalized absorption of C atoms (a) as a function of XCO for a XCH4 of 5%, and (b) as a function of XO2 for a XCH4 of 5%. Reprinted with permission from [61], K. Teii et al., J. Appl. Phys. 87, 4572 (2000). © 2000, American Institute of Physics.

intensity ratio probably due to a decrease in parent CH4 , informing thereby that the major origin of C and C2 is CO. In Figure 7b, the intensity ratios of OH, CH, H, C2 , and O as well as the C atom absorption increase with increasing XO2 . Emissions from O for XO2 = 0%, arising from the impurity from quartz windows, were as low as the noise level. An increase in C atom absorption is ascribed to the result of sequential reactions induced by O atoms. For the transition from upper state j to lower state i, the relationship between the emission line intensities Iji and excitation temperature Tex is given in detail elsewhere [65]. The relative excitational population of H atoms in the Balmer series as a function of excitation level for the three gas mixtures is shown in Figure 8 [61]. The emission line of H was not available because of its low signal-to-noise ratio. The distribution equilibrium of excitation levels of the H atom Balmer series is not satisﬁed because the three sets of plots do not produce straight lines, in accordance with the previous study in the same pressure range [66]. The energy distribution is almost equivalent in any gas mixture, suggesting a common excitation process of H atoms. Furthermore, the overall emission intensities are equivalent in any gas mixture, indicating that the H atom density was

443

Nanocrystalline Diamond

6

ln (λ/gA)

5 4 3

5%CH4-95%H2 5%CO-5%CH4-90%H2 5%CO-95%H2

2

19

20 Energy (0-19 J)

21

IL-62 87

Figure 8. Energy distribution of the H atom Balmer series as a function of excitation level for the three sets of data in the 5%CH4 –95%H2 , 5%CO–5%CH4 –90%H2 , and 5%CO–95%H2 plasmas. Reprinted with permission from [61], K. Teii et al., J. Appl. Phys. 87, 4572 (2000). © 2000, American Institute of Physics.

almost equivalent and the H atoms were produced mainly by dissociation of hydrogen molecules. These similarities for H atoms regardless of the gas mixtures imply that the H atoms were not responsible for the difference in the grown phases. Thus VUVAS analysis suggests that the C atoms resulting mainly from CO by electron impact dissociation increased the precursors to nondiamond phase by promoting gas-phase polymerization, while the OH radicals resulting predominantly by loss reactions of O atoms with H2 and CH4 were responsible for the enhancement of diamond growth. An excess amount of O atoms resulting from O2 interrupt the formation of nuclei at the initial stage of the growth. Besides, the production path of C atoms induced by O atoms is also possible.

4. SPECTROSCOPIC CHARACTERIZATION NCD ﬁlms are materials of a great potential use in various ﬁelds such as tools, optical windows, microelectromechanical systems (MEMS), etc., due to their remarkable properties. In recent years, attempts have been made to develop NCD ﬁlms with a smooth surface and special attention has been paid to NCD ﬁlms with characteristics of a small crystallite size. In order to understand and improve the properties of NCD ﬁlms, it is desirable to relate the remarkable properties of NCD with their microstructure. The various techniques used for characterization of NCD are discussed next.

4.1. Transmission Electron Microscopy Transmission electron microscopy has been the most powerful technique for characterizing the microstructure of NCD and is widely used to obtain bonding information. Figure 9 [5] depicts a typical TEM bright-ﬁeld image of a diamond ﬁlm deposited on a nanometer-scratched substrate with a methane fraction of 4% [5]. A typical selective-area diffraction pattern (not shown here) of a nanometer-scratched sample prepared with a methane fraction of 41% indicates

Figure 9. Bright-ﬁeld TEM image of a NCD ﬁlm grown on nanometerscratched quartz substrate with a methane fraction of 4%. Reprinted with permission from [5], L. C. Chen et al., J. Appl. Phys. 89, 753 (2001). © 2001, American Institute of Physics.

high degree of diamond phase. All the sharp diffraction rings can be indexed to the diamond structure. It is observed that diamond ﬁlms consist of randomly oriented grains and that the substrate pretreatment has a strong effect on the average grain size of the resulting ﬁlms. With a methane fraction of 4%, the average grain size of the micrometerscratched sample was about 200 nm, whereas the corresponding value of the nanometer-scratched sample was only about 4 nm. As the methane fraction was increased from 4% to 41%, the grains of the micrometer-scratched samples became coarser and coarser. In fact, it is noted that the average grain size of the nanometer-scratched sample prepared at a methane fraction of 41% is quite similar to that of its micrometer-scratched counterpart. High-resolution transmission electron microscopy has been utilized to understand the initial stages of diamond ﬁlm growth [12, 67–69]. The detailed electron microscopic study [70] of NCD ﬁlms grown from C60 fullerene precursors showed that various planar defects such as twinning boundaries and stacking faults were observed in large crystallites found in the ﬁlms. Similarly, measurements of the diamond crystallite size distribution using diffraction contrast TEM followed a gamma distribution with an average grain size of 15 nm [71]. Since the resolution of the TEM technique is limited by the small objective aperture, HRTEM has a great edge for reliable grain size measurements [72]. In addition, the TEM imaging techniques give additional information that can be used to determine the phase purity of NCD ﬁlms. The selected-area electron diffraction is a powerful method to examine the phase composition of these ﬁlms as short-range order in either sp2 -bonded glassy carbon or predominantly sp3 -bonded diamondlike amorphous carbon as they provide different diffraction intensities. Gruen et al. have grown NCD ﬁlms on (100) surfaces of n-doped silicon wafers using electrical biasing of the silicon substrate to initiate nucleation [4]. Growth was carried out at a rate of about 0.5–1 m/h. Although NCD ﬁlms have uniform morphology, two additional factors mandate the use of the HRTEM technique. When the packing density is high, the diffraction contrast due to orientational differences of the diffracting crystallites appears weak for

444

Nanocrystalline Diamond

bright-ﬁeld images. Also, the resolution of the diffractioncontrast method is quite low, and, therefore, it is difﬁcult to resolve nanosized crystallite boundaries. For all the reasons mentioned above, reliable information on grain size could be determined from lattice images of nanodiamond crystallites. Mostly the (111) lattice planes are resolvable in the microscopes. Hence, the lattice images from those crystallites are oriented such that at least one set of their (111) lattice planes is parallel to the incident imaging electron beam. The crystallite size with high accuracy can be measured because of the low contrast from surrounding structures is close to random white noise. Although lattice images reveal individual crystallites at atomic resolution, the presence of amorphous carbons cannot be determined. When the crystallite size is reduced to the nanometer scale, the percentage of atoms located at the grain boundaries increases drastically with decrease of grain size. It is estimated that for NCD ﬁlms, the fraction of atoms residing at grain boundaries can be as high as 10% when the average crystallite size becomes 3 to 5 nm. On the other hand, electron microscopy observations show that there is no graphitic phase between the grain boundaries. The presence of sp2 bonds suggests that carbons are -bonded across sharply delineated grain boundaries limited to widths of 0.2–0.5 nm [73].

4.2. Electron Energy Loss Spectroscopy Shown in Figure 10 [74] are high-resolution electron energy loss spectroscopy (EELS) spectra of the plasmon for C60 fullerene, amorphous, nanocrystalline, and polycrystalline diamond [74]. C60 fullerene shows peaks around 6.5 eV which are -plasmons due to the – ∗ band transition

[75, 76]; such peaks are not observed in the amorphous diamond, indicating thereby the absence of -plasmons. Among the three diamond materials, -plasmon peaks that appear around 30–34 eV [77] are different in peak position and shape. Plasmon excitation of valence band electrons of EELS spectra is related to the dielectric function by an equation [78]. Applying the Kramers–Kronig equation, the real and imaginary parts of the dielectric function, 1 and 2 , can be estimated from the plasmon of EELS. The calculated

2 for the diamond materials is shown in Figure 11 [74]. The rise-up position of 2 corresponds to the bandgap. For the natural crystalline diamond, the rise-up position is 5.5 eV, indicating excellent agreement with the established bandgap of diamond (5.5 eV). Two peaks appear at 8.2 and 12.7 eV, and the former is assigned to direct excitation of an electron at the point, the latter being assigned to X and L positions, respectively [79]. For the amorphous diamond, the rise in peak is gradual and the rise-up position changes with the sample from 3.5 to 4.5 eV. In any case, the magnitude of the bandgap is smaller than that of crystalline diamond. A large peak at 7.2 eV, corresponding to point , is observed. For the nanocrystalline powder, the rise-up position is at 3.0 eV and occurs very gradually. This implies that the present method of estimating 2 from the EELS plasmon is reliable. For the amorphous diamond, the magnitude of the bandgap is 3.5 to 4.5 eV, smaller than that of diamond. The variation in magnitude of the bandgap may be caused by the temperature distribution during the synthesis. In addition, the amorphous diamond exhibited the characteristic distribution in 2 and only excitation at the point

7.2

C60 fullerene

π-π*

30

Intensity (arb. units)

nanocrystalline powder

3.3

7.2

ε2

20 amorphous diamond

amorphous diamond

4.5 nanocrystalline powder

12.7

10 8.2

cystalline diamond

5.5 IL-6298

0

5

10

15

20

25

30

crystalline diamond

35

40

45

50

Energy Loss (eV)

Figure 10. High-resolution EELS plasmon spectra for C60 fullerene, amorphous diamond, nanocrystalline powder, and crystalline diamond. Reprinted with permission from [74], H. Hirai et al., Diamond Relat. Mater. 8, 1703 (1999). © 1999, Elsevier Science.

0

IL-6288

0

5

10 Energy Loss (eV)

15

20

Figure 11. Imaginary part of the dielectric function obtained from EELS spectra for diamond materials, showing different rise-up manners. Reprinted with permission from [74], H. Hirai et al., Diamond Relat. Mater. 8, 1703 (1999). © 1999, Elsevier Science.

445

Nanocrystalline Diamond

appeared, while excitations at X and L points were not noticed or considerably weakened. X and L points, where the wave number vector has a certain value, are deﬁned when the direction and periodicity exit in the material, whereas the point is independent of direction and periodicity. The observed characteristics might have originated from the short-range structure of the amorphous diamond. For the nanocrystalline powder the small bandgap of 3.3 eV and the gradual increase in 2 are probably generated from the disordered carbon layer surrounding individual nanometer-sized crystallites observed under high resolution TEM. It is quite likely to tailor the bandgap of diamond ﬁlms by controlling the synthesis parameters.

4.3. X-Ray Photoelectron Spectroscopy The NCD ﬁlms with a thickness of 800 nm were grown by MPCVD and these ﬁlms were implanted with nitrogen ions at energy of 110 keV [80]. During the implantation process, the temperature of the samples was controlled below 80 C using a water-cooling device. X-ray photoelectron spectroscopy (XPS) was used to analyze the changes in surface structure and chemical state of the ﬁlms before and after implantation. Figure 12 [80] presents the XPS spectra for C 1s and O 1s states of the as-deposited ﬁlm. The C 1s envelope has been deconvoluted into three components (Fig. 12a). The main peak at 284.5 eV can be attributed to contribution from C–C and C–H bonds [81]. The peak signals at 285.8 and 286.9 eV are indications of the presence of C–O

and C O [82], respectively. Figure 12b is the O 1s spectra of the as-deposited ﬁlm. A single-peak signal centered at 532.8 eV is evident, which may be assigned to the C–O bond [83]. The surfaces of as-grown CVD diamond ﬁlms are typically terminated with hydrogen atoms [84]. However, the ﬁlm grown as previously stated indicates that the O concentration on the surface reaches 9.83%, which implies that the terminals of dangling bonds of sp3 C–C atoms on the surface are occupied not only by hydrogen atoms but also by some oxygen atoms and absorption of hydrogen and oxygen may also exit on the surface [85]. Shown in Figure 13 [80] are the XPS spectra for C 1s and O 1s states of the ﬁlm implanted with 1 × 1017 ions cm−2 . As compared with the unimplanted sample, distinct changes in the C 1s and O 1s spectra are observed. The C 1s core-level spectrum shifts to a higher binding energy, and the intensity of peaks for C–O and C O shows an obvious increase (Fig. 13a). The O 1s lines are asymmetric, and more than one binding state is presented. Absorbed oxygen (529.6 eV), C O (531.5 eV), C–O (532.8 eV), and absorbed H2 O (534.9 eV) [86] are shown in the O 1s spectrum of Figure 13b. These studies show that nitrogen-ion implantation changes the surface chemical state of the NCD ﬁlm. The C–H and C–O bonds on the surface of as-deposited ﬁlms were broken by highly energetic nitrogen ions. Hydrogen and oxygen desorb from the diamond surface and yield unoccupied surface states, which cause the change in the surface state of the as-deposited ﬁlm. These chemically unsaturated surface carbon atoms have free valences that may be saturated by chemisorption of foreign elements or may react with foreign elements to form bonded surface functional groups [87]. When the implanted samples are exposed in ambient atmosphere, the dangling bonds can be quickly

10 C-H C-C

(a)

10 (a)

6

C-O

8

4 2

N(E)/E

N(E)/E

8

C-O C=O

0 294

290

288

286

284

C-C, C-H

4 C=O

2

IL-6289

292

6

282

Binding Energy (ev)

π-π*

0

IL-6291

294

290

288

286

284

282

10

10

(b)

(b)

C-O

8

N(E)/E

8

N(E)/E

292

6

C-O

6

C=O

4

4 2 H2 O

2 0

0 540

IL-6 29 0

538

536

534 532 530 Binding Energy (ev)

528

526

Figure 12. XPS spectra for: (a) C 1s and (b) O 1s states for the as-deposited ﬁlm. Reprinted with permission from [80], T. Xu et al., Diamond Relat. Mater. 10, 1441 (2001). © 2001, Elsevier Science.

IL-6292

538

536

534

532 530 Binding Energy (ev)

528

526

Figure 13. XPS spectra for (a) C 1s and (b) O 1s states for the ﬁlm implanted with 1 × 1017 ions cm−2 . Reprinted with permission from [80], T. Xu et al., Diamond Relat. Mater. 10, 1441 (2001). © 2001, Elsevier Science.

446 occupied by OH, CO2 , or H2 O molecules and ions in air. In this way the surfaces of NCD ﬁlms are covered by a layer of oxygen-containing groups after high ﬂuence implantation, thereby resulting into the change in chemical state of oxygen on the surface.

4.4. Raman Spectroscopy It is well known that Raman spectra, that are sensitive to changes in translational symmetry, are very powerful for the study of disorder and crystalline formation in thin carbon ﬁlms [88]. In Raman spectroscopy, a visible laser is used for excitation, and a sharp peak at 1333 cm−1 appears in the spectrum. For NCD ﬁlms, with decreasing the grain size, the grain surface area and grain boundary that mainly consist of amorphous sp2 C phases increase. In the Raman spectrum, a broad D peak (disorder-induced mode) at 1350 cm−1 will appear and possibly overlap the diamond peak at 1333 cm−1 with increasing sp2 C bonding. The high frequency stretch modes of sp2 C atoms are overemphasized due to the – ∗ transition resonance effects [89], and the sp2 C network exhibits resonance enhancement in the Raman cross-section since the local sp2 C energy gap of approximately 2 eV is comparable to the energy of the incident photons. The sp3 C atoms do not exhibit such a resonance effect because of the higher local gap of approximately 5.5 eV. Hence the Raman spectra obtained with visible excitation are completely dominated by the sp2 C atoms [90]. Raman scattering in the UV region is promising for vibrational studies of sp3 bonded C phase [91]. Advantages of using UV over visible photons include the suppression of the dominant resonance Raman scattering from sp2 C atoms and increase in the intensity from sp3 C bonding [92]. Figure 14 [93] is the Raman spectra of NCD ﬁlms grown by the MPCVD technique at different temperatures [93]. The Raman spectrum of the NCD ﬁlms mainly exhibits three features near 1150, 1350, and 1580 cm−1 . The bands in the Raman spectra of these ﬁlms near 1350 and 1580 cm−1 , popularly known as D and G bands, are related to graphitic

Nanocrystalline Diamond

islands [90]. The D band appears due to small domain size in graphite. The band near 1150 cm−1 is shown to be related to the calculated photon density of states of diamond and has been assigned to the presence of a nanocrystalline phase of diamond [94, 95]. It is observed from Figure 14 that the NCD feature is almost missing in the ﬁlm grown at 400 C, and this feature increases with temperature in the ﬁlm grown at 500 C and becomes quite pronounced in the ﬁlms grown at 600 C followed by a drastic decrease in its intensity at higher temperature. It has also been observed as a weak band in the microcrystalline diamond ﬁlms along with a sharp peak near 1332 cm−1 giving thereby the signature of crystalline cubic diamond [21, 96]. It is interesting to note that though the ﬁlm grown at 600 C shows an intense band related to NCD, it does not show any peak near 1332 cm−1 . This is mostly due to uniformly distributed short-range sp3 crystallites in the ﬁlms [96, 97]. However, the higher intensities of the graphitic bands in the ﬁlms grown at 600 C, as compared to the intensity of NCD features, do not represent high amounts of sp2 carbon in the ﬁlm. This is because the cross-section of Raman scattering is 50–60 times higher for sp2 -bonded carbon in comparison to sp3 -bonded carbon [98]. This small amount of graphitic carbon in the ﬁlm may exit between the nanodiamond grains [99]. The variation in the intensity of NCD peak can be taken as a qualitative measure of concentration of NCD grain in the ﬁlm. It appears that the growth of NCD starts at 500 C, becomes a maximum at 600 C, and decreases at higher temperatures.

4.5. Infrared NCD ﬁlms deposited by CVD of camphor on Si substrate were analyzed by infrared (IR) spectroscopy [38]. Figure 15 [38] shows the variation of transmittance with wave number of thermal CVD of diamond (curve a) and by rf + dc plasma CVD of CH4 and hydrogen (curve b). It is observed from these ﬁgures that over the IR region, the transmittance of NCD is comparable to that of PCD ﬁlm [100]. The transmittance of bulk diamond (curve c) is also shown

2.86

3.3

λ (µm)

5

6.6 8 10

1.0 0.9

Transmittance

0.8 0.7 0.6

diamond bulk (c)

0.5 0.4

(b)

0.3

(a)

0.2

CO2

SiO2

0.1 0.0 I L-6327 4000 3500 3000 2500 2000 1500 1000 Wave number (cm-1) Figure 14. Raman spectra of the ﬁlms deposited at different temperatures: (a) 400 C, (b) 500 C, (c) 600 C and (d) 700 C. The samples were excited by a 488-mm Ar+ laser. The three positions 1150, 1350 and 1580 cm−1 correspond to nanocrystalline diamond, graphitic D and G bands, respectively. Reprinted with permission from [93], T. Sharda et al., Diamond Relat. Mater. 10, 1592 (2001). © 2001, Elsevier Science.

Figure 15. IR spectra of two representative ﬁlms deposited on Si substrates: (a) deposited by CVD of camphor and hydrogen, (b) deposited by rf + dc plasma CVD of CH4 + H2 , (c) the spectrum of bulk diamond for comparison. Reprinted with permission from [38], K. Chakrabarti et al., Diamond Relat. Mater. 7, 845 (1998). © 1998, Elsevier Science.

Nanocrystalline Diamond

in Figure 15. Clearly, the transmittance of bulk diamond is higher than that of NCD and PCD ﬁlm. The reduction in transmittance of the CVD diamond is related to the formation of the initial layer of silicon carbide on Si substrate, while for bulk diamond the transmittance is higher as no such substrate effect is observed in this case. In addition, the two-phonon absorption is prominent in bulk diamond. The IR spectrum of CVD ﬁlm exhibits two small peaks at wave numbers ∼2325 and 1050 cm−1 that are due to CO2 and SiO2 absorptions. The CVD diamond ﬁlms deposited by CVD of camphor + H2 have much lower average grain size of 0.1 m and smoother surface whereas the ﬁlms deposited by rf + dc plasma CVD have large polycrystals with average grain size of ∼3.2 m. No prominent C–H absorption peak around 2900 cm−1 was noticed in IR spectra indicating thereby a good quality of NCD ﬁlms. On the other hand, the IR spectra of diamond ﬁlms grown on Si substrates by MPCVD exhibit predominant C–H bands [101]. In particular, sp3 -CH2 symmetric stretch at 2850 cm−1 and sp3 -CH2 asymmetric stretch at 2925 cm−1 were identiﬁed.

5. PHYSICAL PROPERTIES Diamond has unique and remarkable properties, for instance high hardness, low friction coefﬁcient [102], high thermal conductivity [103], high chemical stability, and low or negative electron afﬁnity (NEA) [104]. NEA enables electrons at the conduction band minimum to escape from the diamond without an energy barrier at the surface. However, diamond is a good insulator and therefore it is difﬁcult to deliver electrons to the surface. Similarly diamond is a potential material for wear-resistant applications due to its excellent physical and chemical properties. However, conventional CVD diamond coatings that are deposited at high temperatures possess rough surfaces. The high surface roughness is a major problem when these diamond ﬁlms are used for machining and wear applications. The successful growth of thin ﬁlms of NCD at low deposition pressures opens up their large scale uses in tribological applications and ﬁeld emission display devices.

5.1. Mechanical Properties 5.1.1. Surface Roughness The surface of PCD ﬁlms is often very rough and hinders their use for optical and tribological applications. The consensus is that efﬁcient ways must be developed to produce smooth diamond ﬁlms wherein high sp3 content can be maintained. The NCD ﬁlms appear promising toward this goal. It is well known that the key parameter to obtain smooth NCD ﬁlms is the primary nucleation density of diamond nuclei on the substrate. Nucleation densities exceeding 1010 cm−2 can be achieved by ultrasonically scratching the substrate surface with ﬁne-grained diamond powder [9, 105, 106]. A direct relationship between the grain size of diamond powder and primary nucleation density on the substrates, and consequently the optical transparency, has been observed. Other parameters such as the methane fraction in the source gas and the substrate temperature can also affect the morphology and the size of the crystallites. However, previous studies

447 have only focused on the variations in grain size of the diamond powder, keeping the methane fraction in the source gas either constant or varying it over only a small range. The optimized value of the optical transparency at 700 nm and beyond was about 60% and the surface roughness was about 20 nm [9, 105, 106]. The effects of both the grain size of diamond powder and methane fraction in the source gas over a wide range on the microstructure and the optical transmittance of the ﬁlms have been investigated by these researchers. The substrate pretreatment involved ultrasonic polishing of the substrates with diamond powder of two different grain sizes, 4 nm and 0.1 m, for 8 h, followed by standard chemical cleaning with acetone and de-ionized water. For convenience, hereafter the substrates treated with 4 nm diamond powder will be designated as nanometerscratched and those treated with 0.1 m diamond powder as micrometer-scratched. Both types of substrates were placed together in the deposition chamber. Notably, for no ultrasonic polishing at all or polishing for less than 8 h with either of the powders, the resulting ﬁlms were not continuous, indicating that the substrates were not uniformly and entirely scratched for less than 8 h of treatment. The substrates were also cleaned in pure hydrogen plasma at 1.5 kW for 30 min before the depositions. Typical deposition time was about 3–4 h to give a ﬁlm thickness of about 0.5 m. Typical surface proﬁles of NCD ﬁlms grown at two methane fractions on nanometer- and micrometer-scratched substrates are presented in Figure 16 [5]. At low methane fraction, the micrometer-scratched sample had a rough surface whereas the nanometer-scratched sample was smooth. Under this speciﬁc condition, the average surface roughness is comparable to the respective average grain size, which in turn is directly related with the powder size used for substrate pretreatment. Surprisingly, this correlation between the average grain size and the surface roughness was not observed in the samples prepared at high methane fraction. For ﬁlms deposited at 31% methane fraction, despite the similarity in the grain size distribution and the average grain size (∼30 nm), the nanometer-scratched sample has a much rougher surface [root mean square (rms) of roughness ∼100 nm] than its micrometer-scratched counterpart with rms of roughness ∼8 nm. Shown in Figure 17 [5] are some typical transmission spectra of these NCD ﬁlms. It is observed that the transparency of NCD ﬁlm is strongly dependent not only on the methane fraction but also on the pretreatment of the substrate. The optical absorption edge for most of the NCD ﬁlms was quite similar to that for type IIa diamond. However, detail tailing near the absorption edge varied from sample to sample, presumably due to the structural imperfections of the ﬁlms, as well as the internal light scattering at the grain boundaries. Optical transparency over 60% in the spectral range of 0.6–2.0 m is considered sufﬁciently high for most practical applications [105]. For comparison, a variation in the transmittance at 700 nm as a function of the methane content of the source gas is depicted in Figure 18 [5]. Since most of the optical transmittance spectrum showed the inﬂuence of interference of the light in the ﬁlm, the value of a smooth curve that represented the average behavior has been used. At lower methane fractions, the ﬁlms grown on nanometer-scratched

Nanocrystalline Diamond 100 100 0

-100 0

4

8

12 length (µm)

16

20

(b)

100

Transparency (%)

(a)

(a)

80

4%

60

13%

40 37.5% 20

0 -100

0 0

4

8

12 length (µm)

16

20

(c)

100

200

300

400 500 600 700 Wavelength (nm)

800

100

0 -100 0

4

8

12 length (µm)

16

20

(d)

100 0

Transparency (%)

Surface profile (nm) Surface profile (nm) Surface profile (nm) Surface profile (nm)

448

(b)

80

37.5% 60 13%

40

4%

20

-100 IL-6293

4

8

12 length (µm)

16

20

Figure 16. Surface proﬁles of NCD ﬁlms grown at (a), (b) methane fraction of 4%, nanometer- and micrometer-scratched substrates, respectively; (c), (d) methane fraction of 31%, nanometer- and micrometer-scratched substrates, respectively. Both pairs of ﬁlms have similar thickness of about 0.55 (±0.05) m. Reprinted with permission from [5], L. C. Chen et al., J. Appl. Phys. 89, 753 (2001). © 2001, American Institute of Physics.

substrates were considerably more transparent than those on micrometer-scratched substrates. However, as the methane fraction increased, the difference between the transmittance of these two types of ﬁlms diminished and actually reversed at a methane fraction of 20%, such that beyond this methane fraction, the ﬁlms grown on micrometer-scratched substrates were more transparent than those on nanometer-scratched substrates. The ﬁlms grown on micrometer-scratched substrates remained highly transparent until a methane fraction of 42%, while those on nanometer-scratched substrates became fairly opaque for methane fractions beyond 25%. Figure 19 [5] shows the plot between the optical transmittance at 700 nm and the inverse surface roughness and clearly reveals the existence of a region showing a linear relationship between the two parameters. In a second region, saturation of the optical transmittance of the ﬁlms to 80% to 84% occurred despite continuing reduction in their surface roughness. Thus, the major factor that dictates the optical transmittance of NCD ﬁlms is the surface roughness, instead of the grain size, provided that the contents of the sp2 -bonded carbon and the structural disorder were negligible.

5.1.2. Hardness The indentation technique has been most widely used for the determination of hardness. This technique is suitable for ﬁlms with higher thickness since the method requires

0

IL-6 28 4

200

300

400 500 600 Wavelength (nm)

700

800

Figure 17. Optical transmittance spectra of the NCD ﬁlms grown on (a) nanometer-scratched and (b) micrometer-scratched and substrates at various methane fractions. Reprinted with permission from [5], L. C. Chen et al., J. Appl. Phys. 89, 753 (2001). © 2001, American Institute of Physics.

a penetration depth (h) so that h < 01d, where d is the ﬁlm thickness [107]. The stress/strain and hardness in thin ﬁlms may be evaluated indirectly by measuring the physical properties that are inﬂuenced by the mechanical properties. 90

Optical Transmittance at 700 nm (%)

0

80 70 60 50 40 30 IL-6280

0

10

20 30 Methane Fraction (%)

40

Figure 18. Variations in optical transmittance at 700 nm as a function of methane fraction for NCD ﬁlms grown on nanometer-scratched () and micrometer scratched (O) substrates. Reprinted with permission from [5], L. C. Chen et al., J. Appl. Phys. 89, 753 (2001). © 2001, American Institute of Physics.

449

Nanocrystalline Diamond

Optical Transmittance at 700 nm (%)

90

80

70

60

50

40

30 0

2

4

6

8

10

12

14

16

1/Ra (x10-3 )

18 IL-6279

Figure 19. Correlation between optical transmittance at 700 nm and average surface roughness for NCD ﬁlms grown on nanometerscratched () and micron-scratched (O) substrates. The asterisk (∗) denotes the data from Reference 9. Reprinted with permission from [5], L. C. Chen et al., J. Appl. Phys. 89, 753 (2001). © 2001, American Institute of Physics.

Chakrabarti et al. [38] studied the effect of mechanical stress on the optical absorption band tail to determine the strain (a/a) from the theoretical curve-ﬁtting method. This being a nondestructive technique is often very useful since the ﬁlm may be utilized for other applications. If and 0 are the absorption coefﬁcients at any wavelength () and at the band edge, respectively, then a plot of /0 vs (Eg − h, as shown in Figure 20 [38], may be used to obtain the strain (a/a) in the ﬁlm by using the curveﬁtting procedure [108]. The stresses (S) in the ﬁlms were obtained from the relation

Figure 20. Plot of /0 vs (Eg − h) for four representative ﬁlms deposited on quartz substrates, 0 being the absorption coefﬁcient at the band edge Eg. Reprinted with permission from [38], K. Chakrabarti et al., Diamond Relat. Mater. 7, 845 (1998). © 1998, Elsevier Science.

mechanical properties of as-grown and annealed NCD ﬁlms were evaluated. Shown in Figure 21 [110] are the Raman spectra of the as-grown and annealed samples of NCD ﬁlms. The spectra consist of bands near 1150, 1350, 1500, and 1580 cm−1 . The bands near 1350 and 1580 cm−1 are D and G bands, respectively, and are related to graphitic islands. Another band near 1150 cm−1 is assigned mainly to the presence of the nanocrystalline phase of diamond [15, 88, 90]. The band intensity near 1500 cm−1 varies proportionally with

S = Y a/a/1 − with Young’s modulus Y = 600 GPa and Poisson’s ratio = 011. The hardness (Hv) in the ﬁlms were obtained from [109] Hv = 29Y /1 − 1000/x−nc a/a1−nc where nc is the strain-hardening coefﬁcient and x is the indentational strain. The values of S, (a/a), and Hv, for NCD ﬁlms grown from CVD of camphor at various temperatures, were obtained and are in the range 1.9–8.4 GPa, (2–12)×10−3 , and 47–62 GPa, respectively [38].

5.2. Thermal Properties 5.2.1. Thermal Stability To investigate the thermal stability of NCD ﬁlms grown on mirror-polished Si substrates by biased enhanced MPCVD, the NCD ﬁlms were annealed in an ambient Ar atmosphere at 200, 400, 600, and 800 C [110]. The structural and

Figure 21. Raman spectra of the (a) as-grown and annealed samples at (b) 200 C, (c) 400 C, (d) 600 C, and (e) 800 C. Reprinted with permission from [110], T. Sharda et al., J. Nanosci. Nanotech. 1, 287 (2001). © 2001, American Scientiﬁc Publishers.

450

Nanocrystalline Diamond

the band intensity near 1150 cm−1 and is related to the disordered sp3 carbon in the ﬁlms [12, 15]. The Raman results in this case indicate association of these Raman features closer to NCD. If these features were related to hydrogen, their intensity should have decreased with annealing temperature, because the hydrogen should have started to evolve from the samples after 200 C [111]. On the other hand, if these features are related to sp2 -bonded carbon, their intensities should have increased with annealing temperature as the concentration of sp2 carbon increases with annealing temperature [111, 112]. Figure 22 [110] depicts the plot of the intensity ratio of the Raman features near 1150 cm−1 to the graphitic G band as a function of annealing temperature. The ratio does not change signiﬁcantly as a function of annealing temperature and shows an increasing trend up to 400 C but saturates at higher temperatures to a lower value, which is identical to the value of the as-grown sample. Figure 23 [110] depicts the cubic crystalline diamond features in the XRD patterns of the NCD ﬁlms, which also support to some extent the conjecture that the Raman feature near 1150 cm−1 is related to NCD. In addition, the calculated interplanar spacings corresponding to the peaks at 2 ∼ 4405 and 7525 ± 020 in XRD patterns of the ﬁlms closely match the interplanar d-values of (111) and (220) planes of cubic diamond, respectively. The full width at half maximum of diamond peaks in the ﬁlms is high as compared to the CVD grown microcrystalline ﬁlms and these observations are well correlated with that of diamond nanocrystallites [21]. It should be noted that no XRD peaks associated with graphite or features related to amorphous carbon could be identiﬁed either in the as-grown ﬁlms or in the annealed ﬁlms. These results support the assignment of the Raman feature near 1150 cm−1 to the presence of NCD. Similarly, the intensity ratio of the Raman NCD to the graphitic G band does not change much with annealing temperature, indicating thereby the structural stability of NCD ﬁlms to temperatures as high as 800 C.

Figure 23. XRD patterns of the (a) as-grown and annealed samples at (b) 200 C, (c) 400 C, (d) 600 C, and (e) 800 C. Reprinted with permission from [110], T. Sharda et al., J. Nanosci. Nanotech. 1, 287 (2001). © 2001, American Scientiﬁc Publishers.

5.2.2. Speciﬁc Heat Capacity The speciﬁc heat capacity bears important implications on various parameters like thermal conductivity, diffusivity, and thermal expansion. The independent measurement of this quantity is important to gain an understanding of the thermodynamic behavior of NCD. The speciﬁc heat capacity of NCD samples has been measured by differential scanning calorimetry [113]. Shown in Figure 24 [113] is the plot of the speciﬁc heat at constant pressure Cp for the type IIb diamond stone, coarse-grained diamond, NCD, amorphous carbon, and reference data (solid line [114]) vs temperature. The Cp values of the coarse-grained sample are found to be lower than the reference data by less than 10% but are consistent with the measured Cp of the diamond single 2.0

1.2

reference data single crystal coarse grained nanocrystalline amorphous carbon

1.5

0.8 Cp [J/g K]

Intensity Ratio (ln /lg)

1.0

0.6 0.4

1.0

as grown

0.2

0.5

0.0

IL-6330

0

200

400

600

800 IL-6331

Annealing Temperature (oC)

Figure 22. Plot of the Raman intensity ratio of NCD (In) to the graphitic G band (Ig) as a function of annealing temperature. The error bar is the standard error in ﬁtting the individual Raman curves. Reprinted with permission from [110], T. Sharda et al., J. Nanosci. Nanotech. 1, 287 (2001). © 2001, American Scientiﬁc Publishers.

0

100

200 Temperature [ oC]

300

Figure 24. Speciﬁc heat Cp of poly- and nanocrystalline CVD diamond, amorphous carbon and single-crystalline diamond (solid line: [114]) versus temperature. Reprinted with permission from [113], C. Moelle et al., Diamond Relat. Mater. 7, 499 (1998). © 1998, Elsevier Science.

451

Nanocrystalline Diamond 100

75

Tr (%)

Ts = 713 K (13θ) Ts = 673 K (13θ) Ts = 573 K (13θ) 25

0 200

400

600 800 λ (nm)

I L-6332

1000

Figure 25. Transmittance (Tr vs wavelength ( of three representative ﬁlms deposited on fused silica at different substrate temperature (Ts . Reprinted with permission from [38], K. Chakrabarti et al., Diamond Relat. Mater. 7, 845 (1998). © 1998, Elsevier Science.

10

5.3. Optical Properties

400

9 8 300 7

(α hν)2 x 10-9 (eV/cm)2

Figure 25 [38] presents the optical transmittance (Tr ) versus wavelength () traces of three NCD ﬁlms deposited at different substrate temperature (Ts ) on quartz. It may be observed that the transmittance (Tr ) had a high value (∼93%) for Ts > 673 K. The direct and indirect bandgaps (Egd and Egi ), estimated from the corresponding plots of (h2 vs h and (h1/2 vs h is presented in Figure 26 [38], where Egd = 484 eV and Egi = 218 eV. It may be noted that both the direct and indirect bandgaps were high and their values increased with increasing Ts . The surfaces of the as-deposited ﬁlms were generally sp2 rich. The sp2 phases from the surfaces of the ﬁlms could be removed either by etching in oxygen plasma or by using a chemical process [116, 117]. After etching, the absorption edge of the optical spectra became sharper than that of the asdeposited ﬁlm as presented in Figure 27 [38]. There was also a substantial increase in the direct bandgap (Egd ) due to the removal of the sp2 phase from the ﬁlm surface by etching, although the change of indirect bandgap (Egi ) was insignificant. The variation of bandgap energy due to removal of the sp2 phase etching was similar to that reported earlier for NCD ﬁlms deposited by high pressure sputtering of vitreous carbon target [118] and for diamond ﬁlms deposited by CVD of freon-22 and hydrogen [119]. In fact, the etching characteristics of diamond ﬁlms may provide the information of sp2 graphitic and sp3 diamond structure in the ﬁlm because the trigonal sp2 phase can be more easily etched than the tetragonal sp3 [120, 121]. The etch rate will be higher in ﬁlms

50

Sample = 12θ Ts = 673 K

6 5

200

4

α hεν(eV/cm) 1/2

crystal. The NCD sample exhibits a higher Cp at 50 C and slightly lower values for T > 100 C. The accuracy of the experimental values for the NCD and amorphous sample is limited by the very small sample amount and is estimated to be about 15% based on repeated test runs. The Cp values of the amorphous carbon ﬁlms are found to be increased by 0.5 J g−1 K−1 compared to the reference data and are due to the different atomic short-range order and the sp2 :sp3 bonding composition in the amorphous structure. Additionally, the large amount of hydrogen homogeneously incorporated in the amorphous structure contributes to the increase of Cp . The Cp data of the coarse-grained diamond and NCD sample deviate from the reference data by a maximum of 15%. Since the microstructure and hydrogen concentration considerably differ between them, it can be concluded that these properties have a negligible impact on the speciﬁc heat. These results are in good agreement with those of large-grained thick CVD diamond ﬁlms [115]. The impact of nondiamond phases on the speciﬁc heat appears to be small. A residual volume fraction of amorphous carbon in the nanocrystalline sample should make a small contribution to Cp , with Cp of the second amorphous phase being increased by a factor of about 2. The fact that the speciﬁc heat data of the CVD diamond samples fall within a 15% interval around the reference data up to 300 C implies that the temperature dependence of the thermal expansion coefﬁcient should be the same as that of single-crystalline diamond.

3 100 2 Egi = 2.10 eV

1

Egd = 4.80 eV 0

2.0

4.0

6.0

I L-6334

0

hν (eV) Figure 26. Plots of (h2 vs h and (h1/2 vs h for a representative ﬁlm deposited on quartz. Reprinted with permission from [38], K. Chakrabarti et al., Diamond Relat. Mater. 7, 845 (1998). © 1998, Elsevier Science.

452

Nanocrystalline Diamond

CAM-4Q etched

75

CAM-4Q

Tr (%)

unetched

2.0 x 1020

100

(a)

10 1.5 x 1020 1 1.0 x

1020 0.1

5.0 x 1019

50

Conductivity (Ω cm)-1

3

Nitrogen concentration (atoms/cm )

100

0.01

0

5 10 15 20 N2 in gas phase (%)

(b)

20% N2 NCD

25

0

I L-6328

200

400

600 λ (nm)

800

1000

Figure 27. Variation of transmittance (Tr with wavelength ( for a representative ﬁlm before and after removal of sp2 -rich surface layer by chemical etching. Reprinted with permission from [38], K. Chakrabarti et al., Diamond Relat. Mater. 7, 845 (1998). © 1998, Elsevier Science.

having low sp3 /sp2 ratio and also the effect of the etching on the optical absorption edge will be higher in sp2 -rich ﬁlm. The refractive index (n) of the ﬁlms was estimated from the transmittance spectra. n varied within 1.3 to 1.5 in most of the as-deposited nanodiamond ﬁlms showing practically no signiﬁcant change with the variation of wavelength. The low value of n may be due to the presence of a signiﬁcant amount of sp2 phases in the surface of the as-deposited ﬁlm. Removal of sp2 by surface etching resulted in higher n values (i.e., within 1.8–2.2).

5.4. Electrical Conductivity Nitrogen doped ultrananocrystalline diamond (UNCD) ﬁlms with 0.2% of total nitrogen content were synthesized by MPCVD, and temperature dependences of electrical conductivity of these ﬁlms are reported [122]. These results are shown in Figure 28 [122]. In addition, Figure 28a shows secondary ion mass spectroscopy (SIMS) data for the total nitrogen content in the ﬁlms as a function of the percentage of N2 gas added to the plasma. Along with these data is a plot of the room temperature conductivities for the same. The graph shows that the nitrogen content in the ﬁlms initially increases but then saturates at ∼2 × 1020 atoms/cm−3 for 5% N2 in the plasma, which is ∼0.2% total nitrogen content. The increase in room temperature conductivity is dramatic and represents an increase by roughly ﬁve orders of magnitude over undoped UNCD ﬁlms. The value of conductivity increases from 0.016 −1 cm−1 for 1% N2 to 143 −1 cm−1 for 20% N2 . The latter value is much higher than n-type diamond [123, 124] and is comparable to heavily boron-doped p-type diamond [125].

Conductivity (Ω cm)-1

100.00

20% N2

10% N2

0.10

5% N2 1% N2

0.01 0

40

80

120

160

1000/T (K-1 )

200

240 IL-6341

Figure 28. (a) Total nitrogen content (left axis) and room temperature conductivity (right axis) as a function of nitrogen in the plasma. (b) Arrhenius plot of conductivity data obtained in the temperature range 300–4.2 K for a series of ﬁlms synthesized using different nitrogen concentrations in the plasma as shown. Reprinted with permission from [122], S. Bhattacharya et al., Appl. Phys. Lett. 79, 1441 (2001). © 2001, American Institute of Physics.

Figure 28b depicts the temperature dependent conductivity of nitrogen doped UNCD ﬁlm over the temperature range of 300 to 4.2 K. This graph clearly exhibits ﬁnite conduction for temperatures even as low as 4.2 K. This behavior is also seen in heavily boron-doped diamond thin ﬁlms. These graphs are indicative of multiple, thermally activated conduction mechanisms with different activation energies. Hall measurements have been carried on two samples. The carrier concentrations for the 10% and 20% N2 sample were found to be 20 × 1019 and 15 × 1020 cm−3 , respectively, and the carrier mobility of 5 and 10 cm2 /V s for 10% and 20% N2 samples, respectively. The negative value of the Hall coefﬁcients indicates that electrons are the majority carriers in these ﬁlms. It is proposed that conduction occurs via the grain boundaries. Nitrogen in microcrystalline diamond thin ﬁlms usually forms a deep donor level with an activation energy of 1.7 eV [123]. It is therefore likely that the enhanced conductivity in UNCD is due to nitrogen doping

453

Nanocrystalline Diamond

6. APPLICATIONS The high optical transmittance, mechanical, and electrical properties of NCD ﬁlms offer ample opportunities for applications in electron emitting cold cathodes, tribology, MEMS, electrochemical electrodes, surface acoustic wave (SAW) devices, and NCD coatings. A brief discussion of the results of research in these areas is provided.

6.1. Electron Emission Field electron emission, which has important applications in the ﬁeld of ﬂat panel displays and high performance electron guns, has been a subject of extensive studies for many years. Early studies used sharp geometry [126] to attain the local ﬁeld enhancement necessary to extract electron from materials with high work functions, such as Mo or W. This requires complex and expensive patterning and fabricating process. Another approach is to explore new materials with lower emission thresholds. In the past few years, diamond [125, 127] has attracted much attention. It is widely recognized that the property of NEA that can be displayed by diamond does not alone make this material ideal for low ﬁeld electron emission applications. Few electrons exit in the conduction band of diamond, which is difﬁcult to dope n-type, making transport of electrons from metallic contacts and emission of electrons from the diamond surface difﬁcult. However, ﬁne grain, highly defective diamond ﬁlms have shown more promise, presumably due to the presence of a network of grain boundaries and a higher level of nondiamond carbon, which can increase the material’s conductivity. These ﬁlms, often called NCD ﬁlms, are being investigated for ﬁeld emission display devices [4]. The development of ﬁeld emission ﬂat panel displays (FEDs) has triggered very intense research especially in various kinds of carbon thin ﬁlms because one of their properties is to emit electrons at relatively low applied electric ﬁeld [127, 128]. A wide range of carbon ﬁlms, such as amorphous and diamondlike carbon [129, 130], CVD diamond [125, 131], and single and multiwalled carbon nanotubes [132, 133], have shown ﬁeld emission current in the mA range for applied ﬁelds below 10 V/m. Cold cathode ﬁeld emission has been demonstrated in CVD PCD ﬁlms. In the undoped CVD diamond ﬁlms, the space-charge-limited current limits the conductance of the bulk. As a result, grain boundaries, which are highly disordered and may contain codeposited graphitic impurities, have been suggested as the main conducting pathway. Because reduction in diamond grain size may increase the conducting pathways, it is possible to improve diamond ﬁeld emission by depositing size controlled diamond ﬁlms. The ﬁeld emission of CVD grown diamond ﬁlms shows very analogous behavior to the emission of diamondlike carbon ﬁlms. CVD diamond ﬁlms of good quality characterized by a sharp and intense 1332 cm−1 Raman peak seem to exhibit poor ﬁeld emission properties. Zhu et al. [134] reported a relation between the quality of the diamond ﬁlms measured by the full width at half maximum (FWHM) of

the Raman 1332 cm−1 line and the ﬁeld emission properties. By decreasing the crystalline quality and monitoring it by the FWHM of the 1332 cm−1 Raman line, the threshold ﬁeld to get an emission current of 1 nA decreases [134]. The NCD ﬁlms show a very weak Raman line or even its absence and exhibit threshold ﬁeld below 5 V/m. Numerous models such as classical tip emission [135], emission from the conduction band due to a negative electron afﬁnity [136], defect band emission [134], and ﬁeld enhancement at conducting channels in an insulating matrix [137] have been suggested to explain the low ﬁeld electron emission of CVD diamond ﬁlms. The current–voltage characteristics of NCD ﬁlms were measured at a base pressure of 10−8 Torr. The ﬁeld emission measurements were performed using a Keithley 237 instrument with an incorporated high voltage. The emission current is collected by a graphite counter electrode placed 30 m above the anode. The experimental procedures for current density measurements were described in detail elsewhere [138]. Figure 29 [139] shows the typical electron emission characteristics of a NCD ﬁlm deposited on nanometer-scratched (4 nm) substrate with methane of 20%. These ﬁlms were grown for 4 h at 900 C and demonstrate an emission current density of 15 mA/cm2 at an applied ﬁeld of 17 V/m with turn-on ﬁeld of 12 V/m. Here the turn-on ﬁeld expresses a value of ﬁeld at emission current density of 0.01 mA/cm2 . The ﬁeld emission characteristics of the NCD ﬁlms are further analyzed by the Fowler–Nordheim (FN) plot, a plot of log I/V 2 vs 1/V , as depicted in Figure 30 [139]. The plot yields a straight line according to the FN equation [140], −B3/2 A V 2 exp J = d Vd where A and B are constants, current density J is in A/m2 , voltage V is in volts, anode-to-cathode distance d is in meters, is the effective barrier height for electrons in eV, and is the ﬁeld enhancement factor that depends on the 0.02

Current Density (A/cm2)

of the grains. The nitrogen in these ﬁlms is present predominantly in the grain boundaries and not within the grains.

0.015 0.01 0.005 0 -0.005

IL-6303

0

5

10

15

20

Electric Field (V/µm) Figure 29. Electron emission characteristics of a NCD ﬁlm deposited on nanometer-scratched substrate with methane of 20%. Reprinted with permission from [139], P. D. Kichambare, Center for Applied Energy Research, University of Kentucky, Lexington, KY, unpublished work, 2001.

454

Nanocrystalline Diamond

will make it insufﬁcient to form conducting channels; thus the emission properties will be rapidly degraded. Recently, ﬁeld emission from nitrogen-incorporated NCD ﬁlms has also been reported [143]. All these studies indicate that NCD will be an excellent candidate for cold cathode materials for ﬁeld emission devices.

-8.8

Log (I/V2 ) [(A/V2)]

-9.0 -9.2 -9.4 -9.6 -9.8 -10.0

6.2. Electrochemical Electrodes

-10.2

The use of diamond in electrochemistry is a relatively new ﬁeld of research that has begun to blossom in recent years [144, 145]. Conductive diamond possesses several properties that distinguish it from conventional sp2 carbon electrode, like glassy carbon, and makes it most promising for electroanalysis [146–148]. These distinguishing properties are (i) background current densities an order of magnitude lower than freshly polished glassy carbon, (ii) a working potential window of 3 to 4 V in aqueous media, (iii) an extremely stable surface structure resulting in better response precision and long-term response stability, (iv) a high degree of response activity for several aqueous-based analytes, and (v) weak adsorption of polar adsorbates such that the electrode material resists fouling and passivation. It has been demonstrated that untreated diamond outperforms freshly polished glassy carbon in terms of limit of quantitation, response precision, and response stability [146–148]. It is therefore necessary to understand the factors that inﬂuence electron transfer at conducting diamond thin ﬁlm electrode. These electrode materials are challenging to investigate because these are typically polycrystalline with multiple crystallographic orientations, and because they contain extended and point defects, grain boundaries, and low levels of nondiamond carbon impurity. The multiple crystallites, defects, and possible nondiamond carbon phases could provide discrete sites for heterogeneous electron transfer. Within this context, NCD ﬁlms are challenging as it is nearly pure diamond. Recently, the electrochemical properties of NCD thin ﬁlms deposited from C60 /Ar and methane/nitrogen gas mixture were investigated [149]. Figure 31 [149] depicts the cyclic voltammetric i–E curve for a nanocrystalline diamond ﬁlm deposited from a C60 /Ar mixture and a boron-doped, microcrystalline diamond ﬁlm deposited from a CH4 /H2 mixture. The voltammograms were obtained in 1 M H2 SO4 at 50 mV/s. The open circuit potential of the untreated nanocrystalline ﬁlm was +174 mV (vs saturated calornel electrode (SCE)). Figure 31a shows voltammograms for nanocrystalline and microcrystalline ﬁlm between −05 and 1.0 V (vs SCE). The curve for the nanocrystalline ﬁlm is basically featureless within this potential range, and after the initial scan, the curves are reproducible in shape. The anodic current at 0.1 V is 2.5 A/cm2 , and this is signiﬁcantly smaller than the cathodic current, 20 A/cm2 observed for polished glassy carbon. The open circuit potential for the untreated microcrystalline ﬁlm was +205 mV (vs SCE). The voltammogram for this ﬁlm between −05 and 1.0 V is not quite as featureless, as there are two small anodic and cathodic peaks at 0.75 and 0 V, respectively. These features are associated with a quasi-irreversible but as yet unknown surface redox process, probably at the nondiamond carbon impurities in the grain boundaries. The peak current ratio is near 1, so the

-10.4 0.8

0.9

1.0 1000/V (1/V)

1.1

IL-6294

1.2

Figure 30. The Fowler–Nordheim plot depicting ﬁeld emission characteristics of NCD ﬁlm. Reprinted with permission from [139], P. D. Kichambare, Center for Applied Energy Research, University of Kentucky, Lexington, KY, unpublished work, 2001.

emitter geometry. Consequently, a plot of logJ /V 2 vs 1/V yields a straight line with slope B3/2 d/ indicating that the ﬁeld emission property can be explained by a tunneling mechanism. The slope of the FN plot can be used to determine the work function of the emitter. These studies indicate that the NCD ﬁlms would be suitable for application in electron emission devices as these devices operate at 10 mA/cm2 . The ﬁeld emission properties of the NCD ﬁlms [138] can be explained by a conducting–tunneling mechanism. A model based on the graphite/nanodiamond mix-phase structure reported earlier [141] is considered. The NCD ﬁlms used for ﬁeld emission measurements were grown on silicon substrates. In such a structure, graphite plays the role of conducting channels from the silicon substrate to the ﬁlm surface. While on the surface, it is assumed that the nanodiamond has a relatively low or even NEA as that of bulk. Thus electron will ﬁrst tunnel through the nearby diamond edges and then emit from the diamond surface [127, 142]. There are two main parameters deciding the ﬁeld emission. The ﬁrst is the diamond grain size and the other is graphite content in the ﬁlm. As electrons from the diamond/graphite interface must tunnel through a barrier between them, the diamond grain size is a critical factor determining the tunneling probability. For large-size diamond grains, electrons can only be emitted from regions close to crystal edges, which are thin enough for electrons to tunnel through. In contrast, for small-size diamond grains, electrons can be effectively emitted from a larger surface area, or even the whole diamond surface, thus greatly increasing the emission site density. The graphite content is another critical factor as the emission starts from the diamond/graphite interface. There is an optimum graphite content for maximizing the graphite/diamond interface area on the ﬁlm surface. The optimal value of the graphite content is just enough to ﬁll the gaps between the densely agglomerated diamond grains. Above this value, a decrease in graphite content will increase the diamond/graphite interface area and thus enhance ﬁeld emission. But when the graphite content is below this value, a decrease in graphite content will decrease the interface area at the surface, and the ﬁeld emission properties will drop. Moreover, a further decrease of the graphite content

Nanocrystalline Diamond

Figure 31. Cyclic voltammetric i–E curves for a nanocrystalline diamond ﬁlm deposited from a 1% C60 /Ar gas mixture and a boron-doped microcrystalline diamond ﬁlm deposited from a 0.3% CH4 /H2 gas mixture over (a) a potential range from −05 to 1.0 V (vs. SCE) and (b) a potential range from −16 to 2.0 V (vs. SCE). The electrolyte was 1 M H2 SO4 and the potential sweep rate for (a) was 50 mV/s and for (b) was 25 mV/s. Reprinted with permission from [149], B. Fausett et al., Electroanalysis 12, 7 (2000). © 2000, Wiley-VCH.

surface sites that are being oxidized on the forward sweep are fully reduced on the reverse sweep. The anodic current at 0.1 V is 0.7 A/cm2 , which is a factor of 4 lower than the value observed for the NCD. It should be noted that the microcrystalline ﬁlms show no evidence for any surface redox processes. The slightly higher background current for this particular nanocrystalline ﬁlm has been mentioned due to the increased fraction of grain boundary carbon. Figure 31b shows voltammograms for the nanocrystalline and microcrystalline ﬁlms between −14 and 2.0 V (vs SCE). The electrolyte was 1 M H2 SO4 and the potential sweep rate was 25 mV/s. The working potential window in this medium for the NCD ﬁlm is ca. 3 V (±250 A/cm2 and is larger than the ca. 2.5 V window observed for freshly polished glassy carbon. The response shows a large anodic peak at 1.5 V and a smaller cathodic peak at 0.4 V. The anodic peak at 1.5 V is likely due to redox-active carbon in the grain boundaries. This indicates that some reduction of these surface carbon–oxygen functionalities must occur during the cathodic potential sweep in order to allow for their reoxidation on the subsequent forward sweep. The voltammogram for the boron-doped microcrystalline ﬁlm, between −14 and 2.0 V, reveals a working potential window of 3.3 V, which is slightly wider than observed for NCD ﬁlms. The voltammogram shows an anodic peak at 1.8 V just prior to oxygen evolution; however, the peak charge is signiﬁcantly less than that for the NCD ﬁlms. This reﬂects the reduced coverage of redox-active grain boundary carbon on the microcrystalline ﬁlms. These studies indicate that NCD ﬁlms prepared from C60 /Ar gas mixtures appear to have basic electrochemical properties similar to boron-doped microcrystalline diamond ﬁlms with a wide working potential window and a low voltammetric background current. The potential application of NCD in electroanalysis seems to be due its nanocrystalline structure.

6.3. Tribology The unique properties of diamond (i.e. the highest hardness, stiffness, and thermal conductivity as well as imper-

455 viousness to acidic and saline media) are exceptional and far exceed those of any other known material. The cleaved diamond surfaces exhibit one of the lowest friction coefﬁcients of any known material. The combination of these qualities in a material is ideal for highly demanding tribological applications. The prospects for large-scale uses of diamond in tribology increased sharply when it was discovered that diamond can be grown as thin ﬁlms at low deposition pressures by various methods [150–153]. High quality microcrystalline diamond ﬁlms exhibit most of the desired properties of the natural diamonds. They are made up of large columnar grains that are highly faceted and generally rough. They tend to grow continuously rougher as the thickness of the deposited ﬁlms increases. The generally rough surface ﬁnish of these ﬁlms precludes their immediate uses for most machining and wear applications. When used in sliding-wear applications, such rough ﬁlms cause high friction and very high wear losses on mating surfaces [154–156]. The rough surface can be polished by laser beams and ﬁne diamond powders or by rubbing against a hot iron plate [157]. Despite the high interest in using diamond ﬁlms for diverse tribological applications, their widespread utilization in the industrial world has not yet met expectations [158–160]. Recently, new methods have been developed for deposition of smooth NCD ﬁlms [161] and the ﬁlms afford ultralow friction and wear in sliding tribological applications. As elaborated above, conventional diamond ﬁlms are generally rough and consist of large grains. Depending on deposition conditions, these grains exhibit 111 or 100 crystallographic growth orientations. Figure 32 [162] depicts the friction coefﬁcients of microcrystalline diamond ﬁlms (surface roughness: 0.35 m, rms) and NCD (surface roughness: 30 nm, rms) ﬁlms against Si3 N4 balls in open air and dry N2 [162]. After an initial run-in period during which friction is relatively high, the friction coefﬁcient of the NCD ﬁlm decreases rapidly to ∼0.1 in open air and to ∼0.05 in dry N2 , whereas the friction coefﬁcient of the microcrystalline diamond ﬁlm remains high and unsteady in both test environments. Wear rates of Si3 N4 balls slid against the smooth NCD and rough microcrystalline diamond ﬁlms in air and dry N2 are shown in Figure 33 [162] which clearly shows that the wear rates of balls slid against the NCD ﬁlms are more than two orders of magnitude lower than the wear rates of balls slid against the rough microcrystalline diamond (MCD) ﬁlms. The high friction coefﬁcients of rough diamond ﬁlms can be attributed to the abrasive cutting and plowing effects of sharp asperity tips on softer counterface pins, whereas the plowing effects created by smooth NCD ﬁlms on counterface balls are minimal and thus exhibit lower friction. Previous studies also demonstrated a close correlation between higher surface roughness and greater frictional losses [155, 163, 164]. When the rough MCD ﬁlms were polished and then used in sliding-contact experiments, very low friction coefﬁcients were obtained [161, 163, 164]. Apart from physical roughness, surface chemistry and tribo-induced adhesive interactions can also occur and play a dominant role in the friction and wear performance of all diamond ﬁlms [165]. The nature and extent of interactions may be controlled by the environmental species or by ambient temperature [166]. Mechanistically, the low-friction

456

Nanocrystalline Diamond

a combination of low friction and high wear resistance under a wide range of sliding-contact conditions. NCD coatings have been applied to seals using dc-biased substrates and an oxyacetylene torch [161, 169].

1.0 Rough Diamond Smooth Diamond

Friction Coefficient

0.8 0.6

6.4. Superhard Nanocrystalline Composites

0.4 0.2 0

0

1,000

2,000

(a)

3,000

4,000 5,000

6,000

7,000

8,000

Number of Sliding Cycles 1.0 Smooth Diamond Rough Diamond

Friction Coefficient

0.8 0.6 0.4 0.2 0

IL-62 97

0 (b)

1,000

2,000 3,000 Number of Sliding Cycles

4,000

5,000

Figure 32. Friction coefﬁcients for rough MCD and smooth NCD ﬁlms against Si3 N4 balls in (a) open air and (b) in dry N2 . (Test conditions: load, 2 N; velocity, 0.05 m s−1 ; relative humidity, 37%, sliding distance, 40 m, ball diameter, 9.55 mm.) Reprinted with permission from [162], A. Erdemir et al., Surf. Coat. Technol. 120–121, 565 (1999). © 1999, Elsevier Science.

nature of smooth or cleaved diamond surfaces has long been attributed to the highly passive nature of their surfaces [167, 168]. Speciﬁcally, it has been postulated that gaseous adsorbates, such as hydrogen, oxygen, or water vapor, can effectively passivate the dangling surface bonds of diamond. When the dangling bonds become highly passive, the adhesion component of friction is diminished [162]. Thus, unlike most other engineering materials, NCD offers

Figure 33. Wear rates of Si3 N4 balls slid against smooth NCD and rough MCD in air and dry N2 (test conditions: load, 2 N; velocity, 0.05 m s−1 ; relative humidity, 37%; sliding distance, 40 m; ball diameter, 9.55 mm). Reprinted with permission from [162], A. Erdemir et al., Surf. Coat. Technol. 120–121, 565 (1999). © 1999, Elsevier Science.

Materials with Vickers hardness of ≥40 GPa are called superhard materials. The hardness of diamond depends on its quality and is typically in the range of 70 to 90 GPa. It is limited by dislocation sliding and crack propagation along the 111 crystallographic direction [170, 171]. The theoretical strength of ideal single crystalline materials can be easily estimated from the critical stress, which is needed to cause shear sliding of lattice planes [172]. However, the practically achievable strength of single crystals and engineering materials is orders of magnitude smaller because plastic deformation and fracture occur due to inherent ﬂaws, such as multiplication and propagation of dislocations and crack growth, which are present in the materials. The design of novel superhard materials is based on the formation of stable nanocomposite consisting of two or more phases, for example nanocrystals of hard transition metal nitride imbedded in an amorphous phase [173–175]. When such a nanostructure is formed by a thermodynamically driven segregation of transition metal nitride and silicon or boron nitride, it is fairly stable against recrystallization even at high temperatures of ≥1000 C [176, 177]. The resulting strong interface and compact grain boundaries avoid grain boundary sliding, which would otherwise limit the strength and hardness of nanocrystalline materials [178–181]. Beside the hardness, the nanocomposites also show a very high elastic recovery of 80% to 90% and toughness. The nanocrystalline diamond seems to be one of the promising superhard nanocomposite materials if the properties of NCD are properly tailored. The superhardness combined with a high elastic recovery and toughness will be very important for many practical applications. Hardness of a material is a measure of its ability to resist deformation upon load. Hence, the meaning of hardness depends on the exact nature of that load resulting in various kinds of engineering scales of hardness, such as scratch, static contact indentation, or dynamic ones. The most universal measure of the hardness is the energy of the plastic or pseudoplastic deformation, which is calculated from the ratio of the applied load to the area of that deformation. Depending on the shape of the indenter used, various scales of the static hardness result. For the characterization of thin ﬁlms the Vickers or alternatively Knoop hardness is most commonly used. During the measurement a diamond indenter of pyramidal shape with an angle between the opposite faces of 136 is pressed into a material with a given load L and after unloading, the area of the remaining plastic deformation A is measured. The hardness in GPa or kg/mm2 is then calculated from the equation H = ∝L/A, where the constant ∝ depends on the type of the indenter which differs by the angles between the opposite faces of the diamond pyramid. Figure 34 [175] shows examples of indentation into a NCD and into ultrahard nc-TiN/a- and nc-TiSix , x ≈ 2. It should be noted that the measurements on super- and ultrahard materials including diamond show large scattering and, therefore,

457

Nanocrystalline Diamond

indentation depth [µm]

0.25

nc - diamond Hv = 88 GPa E = 534.3 GPa HU = 1879 kg/mm2

0.20 0.15 0.10 0.05

(a) 0.00

0

5

10

15 20 load [mN]

25

30

0.40 nc - TiN/a-& nc - TiSi2 Hv = 138.9 GPa E = 607.4 GPa HU = 1914.4 kg/mm2

indentation depth [µm]

0.35 0.30 0.25 0.20 0.15 0.10 0.05 0.00

(b) 0

10

20

30 40 load [mN]

50

60

IL-6338

70

Figure 34. Example of indentation into (a) NCD and (b) nc-TiN/a- & nc-TiSix (x ∼ 2) coatings. The area between the lower (loading) and upper (unloading) curve corresponds to the energy of pseudoplastic deformation, the area between the loading curve and the y-axis corresponds to the total energy of deformation. Reprinted with permission from [175], P. Nesladek and S. Veprek, Phys. Status Solidi A 177, 53 (2000). © 2000, Wiley-VCH.

the obtained values of hardness should be considered as relative ones. Therefore, one can conclude from Figure 34 that the microhardness of the nc-TiN/a- and nc-TiSix ﬁlms is at least equal to that of the hardest diamond.

6.5. Conformal Coatings The excellent mechanical properties of diamond suggest that diamond is one of the best materials for MEMS applications. In addition, the chemical inertness of diamond makes it a suitable as a corrosion protection material. Recently, NCD ﬁlm growth has been demonstrated on various substrates of engineering interest [182]. The substrate of choice in most diamond deposition studies has been silicon while for metals, the thermal expansion mismatch between the diamond ﬁlm and substrate gives rise to thermal stress, which often results in delamination of the ﬁlm. Thus, one of the major barriers in obtaining a diamond ﬁlm that is resistant to delamination is the inherent difference between the coefﬁcient of thermal expansion between the diamond ﬁlm and a metal substrate. The desirable mechanical properties with coating containing NCD and/or amorphous carbon have been reported [183]. The combination of nanocrystalline and amorphous carbon components can result in coatings with high toughness, high hardness, and low surface roughness [183, 184]. These ﬁlms are believed to consist of NCD grains imbedded in a primarily amorphous carbon matrix, which itself

has some sp3 -bonded carbon content. These coatings have been produced with hardness up to 80% that of natural diamond and yet still exhibit appreciable toughness uncharacteristic of a pure ceramic material [182]. NCD coatings can be tailored in their coating structure and mechanical properties by appropriate changes in CVD feed gas chemistry [182]. The relative concentration of CH4 and N2 is shown to strongly inﬂuence coating structure, hardness, and adhesion. This can provide an opportunity to choose a balance between coating hardness and toughness as required for a particular application. There is general agreement that diamond ﬁlm growth occurs most readily on pure metal that supports the formation of a stable carbide layer. There has been controversy as to whether the formation of a surface carbide ﬁlm was a prerequisite for diamond growth. Ramanthan et al. [185] have studied the transition metals and found that diamond growth is supported by those metals that tend to produce stabilized sp3 carbon structures. Diamond ﬁlm adhesion is critically linked to the thermal expansion mismatch between diamond and metal. At higher deposition temperatures more stress may be induced in the ﬁlm, providing a higher driving force for ﬁlm buckling or delamination. Ager and Drory [186] measured a residual compressive stress of about 7 GPa for a diamond ﬁlm on Ti–6Al–4V that was in excellent agreement with the stress predicted from calculation. With such large residual ﬁlm stresses it would be advantageous to be able to deposit diamond at low substrate temperature. It has been well established that adding small quantities of oxygen can lower the temperature range for which diamond can grow [187, 188]. A small amount of CO and O2 diluted in H2 has also resulted in high quality diamond growth at substrate temperatures in the range of 411 to 750 C [189]. However, in high H2 -dilution conditions, the deposition rate decreases rapidly with temperature. This is because the precursor radical that is widely believed to be responsible for diamond growth under high H2 dilution is thermally activated [190, 191]. An interlayer can be used as a barrier to the diffusion of carbon into the substrate, thus providing sufﬁcient carbon concentrations at the surface to nucleate diamond. The interlayer may also be used in order to minimize interfacial stresses and to provide an intermediate layer for bonding. Diamond deposition onto WC–Co and steel substrates has been achieved using a multiplayer structure of silver and refractory metals with a resulting improvement in adhesion [192]. Such coating is typical of that required for good WC–Co cutting tool properties. Figure 35 [182] shows a thick NCD ﬁlm grown on WC–Co that was prepared with a chemical treatment prior to coating and it depicts an extremely smooth and ﬂat surface. Similarly, results of conformal coatings of NCD are very encouraging and a good degree of coating conformality has been reported wherein a hexagonal shaped silicon needle ∼5 m in diameter was successfully coated with ∼2000 Å NCD [4].

6.6. SAW Devices SAW devices have found several key applications in radio frequency and microwave electronics [193]. They offer a high degree of frequency selectivity with low insertion loss,

458

Nanocrystalline Diamond electrodes ZnO Nanocrystalline diamond

Si substrate IL-6337

Figure 36. Schematic diagram of the surface acoustic wave device multiplayer structure. Reprinted with permission from [200], B. Bi et al., Diamond Relat. Mater. 11, 677 (2002). © 2002, Elsevier Science.

making them highly suitable for use as narrow band ﬁlters. SAW devices are particularly well adapted to microwave integrated circuits since they can provide a signiﬁcant size reduction over purely electromagnetic devices. These devices are most typically implemented on piezoelectric substrates on which thin metal ﬁlm interdigitated transducers (IDTs) are fabricated using photolithography. The use of diamond as a SAW substrate offers an attractive means for relaxing the lithographic criteria [194]. With a surface wave velocity ∼ 1 × 1014 m s−1 , diamond allows SAW device operation near 2.5 GHz with nominal 1 m linewidths. Since diamond is not piezoelectric, additional complexity is introduced by a requisite overlayer of a piezoelectric thin ﬁlm, typically ZnO. Sound propagation in layered media may be highly dispersive and in general admits a multiplicity of allowed modes. Nevertheless, highly successful devices based on ZnO/polycrystalline diamond/Si layered structures have been reported [195–199]. NCD is a new form of diamond and differs from diamondlike carbon in that it contains relatively little hydrogen or sp2 -bonded carbon. The properties of NCD ﬁlms like smooth surfaces and small crystallite size are most attractive and relevant for SAW applications, since standard PCD on Si is quite rough and must be smoothed by mechanical polishing before photolithographic processing can be attempted. Furthermore, one expects acoustic scattering at large angle grain boundaries in PCD, especially if lateral grain dimensions exist on length scales between acoustic wavelength and SAW device apertures and transducer separations. The use of NCD eliminates these concerns. SAW devices based on NCD have been fabricated [200] as shown schematically in Figure 36 [200]. This device was studied using a frequency and time domain method. Phase velocities were obtained from device resonant frequencies measured with a network analyzer. Figure 37 [200] presents a compilation of experimental results and calculations of the phase velocity as a function of khZnO . It is evident that several modes are allowed for a given value of khZnO . The modes

6.7. MEMS Devices Thin ﬁlm ferroelectrics are important for MEMS because of their strong piezoelectric effect and high energy density originating from the very high dielectric constant of ferroelectric materials [201, 202]. Further, these MEMS devices are fabricated primarily in silicon because of the available surface machining technology. A major problem with Si-based

Dispersion Curves of ZnO/Diamond/Si Structure kh(diamond) = 4 12000 1st

2nd

10000

Phase Velocity (m/s)

Figure 35. Thick NCD coating made on chemically treated surface of WC–Co cutting tool insert. Edge shown is a cutting edge of the tool. Reprinted with permission from [182], R. Thompson et al., in “Proceedings of the 1st ASM International Surface Engineering and the 13th IFHTSE Congress and Exposition” (O. Popoola et al., Eds.), 7–10 October 2002, Columbus, OH. © 2002, ASM International.

are highly dispersive at small values of khZnO (i.e., the phase velocity is strongly dependent on khZnO ). As expected, the lowest order mode tends toward the phase velocity of the Rayleigh wave on (0001) ZnO at large khZnO but approaches the diamond Rayleigh wave velocity as khZnO → 0. It should be noted that the large phase velocity suggests that carbon within the grain boundaries is strongly bonded. Since surface waves are dominated by elastic shear strains, the grain boundary carbon is highly resistant to bond-bending forces. This indicates that most of the carbon is highly coordinated.

8000 0th 6000

4000 VR (ZnO) 2000

IL-6336

0

1

2 kh(ZnO)

3

4

Figure 37. Phase velocities as measured for surface waves on NCD, diamond () and on large-grain polycrystalline diamond ( ). The solid line represent calculations of phase velocities based on single crystal diamond material parameters. The label denotes the Rayleigh mode indices for the layered medium. The dashed line shows the Rayleigh wave velocity on ZnO. Reprinted with permission from [200], B. Bi et al., Diamond Relat. Mater. 11, 677 (2002). © 2002, Elsevier Science.

459

Nanocrystalline Diamond Table 2. Electrical properties of diamond thin ﬁlms. Thin ﬁlms

Figure 38. Combined lithographic patterning and selective deposition methods. Reprinted with permission from [207], A. R. Krauss et al., Diamond Relat. Mater. 10, 1952 (2001). © 2001, Elsevier Science.

MEMS technology is that Si has poor mechanical and tribological properties [203, 204]. On the other hand, diamond is an ultrahard material with high mechanical strength, exceptional chemical inertness, and outstanding thermal stability. The friction coefﬁcient of diamond is exceptionally low and the projected wear life is 10,000 times greater than that of Si, making diamond an ideal tribomaterial for MEMS components [205, 206]. Lubrication poses a major limitation on the design of MEMS devices since the usual methods of delivering lubricant to the interface between contacting parts are very difﬁcult to implement. There are a number of ways in which diamond components can be fabricated for MEMS applications using thin ﬁlm deposition methods [207]. One of the methods of obtaining the tribological beneﬁts of diamond while exploiting the availability of Si fabrication technology is to produce Si components to near-net shape and then to provide a thin, low wear, low friction diamond coating [208–213]. This approach only works if the diamond ﬁlm can be produced as a thin, continuous, conformal coating with exceptional low roughness on the Si component. Conventional diamond CVD deposition methods result in discontinuous ﬁlms with a low density of large grains and poor ability to form thin conformal coatings on Si microstructures [214]. Recently, Krauss et al. successfully developed a fabrication technique [207] for conformal UNCD coating that is comparable to the lithography, galvanoformung, Abformung (LIGA) method and permits fabrication of complex three-dimensional shapes as opposed to the two-dimensional structures produced by Si microfabrication methods. The selective deposition represents a second method that may be used for the production of UNCD microstructures but has no analog in Si microfabrication technology. All diamond ﬁlms require a nucleation layer, usually achieved by exposing the substrate to a suspension Table 1. Mechanical properties of diamond thin ﬁlms. Thin ﬁlms Grain size Speciﬁc heat (J g k−1 ) Friction coefﬁcient Surface roughness Young’s modulus Hardness

NCD

UNCD

PCD

Ref.

50–100 nm 0.68

2–5 nm —

0.5–10 m 0.54

[113] [113]

0.06

—

0.35

[168, 207]

50–100 nm

20–40 nm

400 nm–1 m

[207]

600 GPa

—

1200 GPa

[38, 207]

62 GPa

—

100 GPa

[10, 38]

Direct bandgap Indirect bandgap Field emission threshold ﬁeld Electronic bonding H content

NCD

UNCD

PCD

Ref.

4.84 eV 2.18 eV 1 V/m

— — 3.2 V/m

5.45 eV — 22 V/m

[38, 74] [38] [8, 140]

up to 50% sp2

2–5% sp2

sp3

[207]

<1%

<1%

<1%

[207]

of ﬁne diamond particles. It is possible to seed a selected portion of the substrate by (a) using photoresist to prevent exposure of selected areas to the diamond powder, (b) using diamond-loaded photoresist to produce a patterned nucleation layer, or (c) seeding the substrate uniformly and then selectively etching portions of the surface to remove diamond seeded areas. Alternatively, it is possible to use SiO2 as a mask [215–217], since diamond ﬁlms grown from CH4 –H2 plasmas are normally unable to form on SiO2 substrates. The feature resolution that can be achieved by this method is limited by the grain size [216]. The lithographic patterning is the third method to produce diamond multiplayer structures. In this method, it is necessary to deposit a thin diamond ﬁlm on a sacriﬁcial release layer such as SiO2 . Methods have been reported [216, 217] to force the growth of conventional CVD diamond on SiO2 , typically by damaging the SiO2 surface by ultrasonic abrasion in a diamond powder suspension. Continuous diamond ﬁlms have been achieved using this method for ﬁlm thickness in the range 15–20 m although there are gaps between the ﬁlm and the substrate, and the ﬁlms are extremely rough [217]. Diamond ﬁlms grown from CH4 –Ar plasmas are able to form directly on SiO2 substrates without the need for damaging the SiO2 layer since the gas-phase carbon dimer growth species in the UNCD process forms an intermediate SiC nucleation layer. The use of SiO2 as a sacriﬁcial substrate layer for fabrication of multiplayer UNCD devices has also been demonstrated [207]. Thus by combining selective deposition and lithographic etching of blanket UNCD ﬁlms, Krauss et al. fabricated multilevel diamond devices [207] using methods that are not feasible or require a large number of steps in Si. A scheme for fabricating a captive rotor using combined methods is shown in Figure 38 [207]. A blanket UNCD ﬁlm is deposited on a Si substrate, followed by deposition of photoresist to yield the structure shown in Figure 38a. A SiO2 layer is then deposited and used as a hard mask in combination with an oxygen plasma etch to produce the structure shown in Figure 38b. A conformal SiO2 layer is then deposited, and the trenches are selectively seeded as shown in Figure 38c using one of the methods described in the previous paragraph. The seeded trenches are ﬁlled with UNCD, and the SiO2 layer is then removed by HF etch, leaving the wheel free within the retaining trench of the stator as presented in Figure 38d. Further UNCD growth results in both an increase in height and a lateral spread of the stator components, which act to capture the rotor within the track as shown in Figure 38e.

460

7. SUMMARY NCD has many remarkable properties for variety of applications and these ﬁlms have been recognized as an important class of structures in the family of diamond materials. With the advent of low pressure and temperature CVD techniques, it is now possible to control the microstructure, surface roughness, and optical transmittance of the NCD ﬁlms. The grain size of the NCD ﬁlms can be varied from 4 nm to a few hundred of nanometers. The OES investigations indicate that the nucleation and growth mechanism depends on a methane fraction. It is observed that C2 dimer is the predominant growth species and plays a major role in the growth of NCD ﬁlms. The optical transmittance of NCD ﬁlms is a strong function of the surface roughness but not the grain size of the NCD ﬁlms. Ultrasmooth and highly transparent NCD ﬁlms can be grown on either nanometer- or micrometerscratched substrates with either low or high, respectively, methane fractions. These smooth diamond ﬁlms can be used in mechanical-seal applications, for protection against wear, and as abrasion resistant coatings in several demanding tribological applications. The ﬁeld emission studies of NCD ﬁlms showed a substantial reduction in emission turn-on voltage indicating their tremendous promise as a cold cathode material. In addition, these NCD ﬁlms possess unique electrochemical behavior and also exhibit a great potential for application in SAW and MEMS devices. Thus in view of many interesting properties of NCD ﬁlms that have been summarized in Tables 1 and 2, they proved to be one of the important emerging nanostructured materials with many potential applications.

Nanocrystalline Diamond

Thermal conductivity For steady-state heat ﬂow, the proportionality constant between the heat ﬂux and the temperature gradient. Also, a parameter characterizing the ability of a material to conduct heat. Transmission electron microscopy (TEM) A microscope that produces an image by using electron beams to transmit (pass through) the specimen, making examination of internal features at high magniﬁcation possible. Tribology The science and technology of two interacting surfaces in relative motion and of related subjects and practices. The popular equivalent is friction, wear, and lubrication in surfaces sliding against each other, as in gears and bearings.

ACKNOWLEDGMENTS P. D. K. is thankful to Professor L. C. Chen, National Taiwan University, Taipei, Taiwan, and Professor K. H. Chen, Institute of Atomic and Molecular Sciences, Taipei, Taiwan, for introducing him to the ﬁeld of nanocrystalline diamond. In addition, he gratefully acknowledges a PostDoctoral Fellowship from the National Science Council, Taiwan and National Science Foundation–Nanoscale Interdisciplinary Research Teams, USA.

REFERENCES 1. 2. 3. 4. 5.

GLOSSARY

6.

Chemical vapor deposition (CVD) A method for depositing some of the layers which function as dielectric, conductor, or semiconductor. A vaporized chemical containing atoms of the materials to be deposited reacts with another, liberating the desired material, which deposits on the substrate while the by-products of reaction are removed from the reaction chamber. Electron emission The liberation of electrons from an electrode into the surrounding space, usually under the inﬂuence of a high electric ﬁeld, light, or heat. Electron energy loss spectroscopy (EELS) A technique for studying atoms, molecules, or solids in which a substance is bombarded with monochromatic electrons, and the energies of scattered electrons are measured to determine the distribution of energy loss. Field emission display (FED) This is an array of a million or more electron guns, each exciting a single phosphor pixel. The electron guns are typically triodes. Hardness The measure of a material’s resistance to deformation by surface indentation or by abrasion. Speciﬁc heat capacity The heat capacity per unit mass of material. Surface roughness The ﬁner irregularities of the surface texture that usually result from the inherent action of the production process or material condition.

7. 8.

9. 10.

11. 12. 13.

14. 15. 16. 17.

18. 19.

R. C. DeVries, Annu. Rev. Mater. Sci. 17, 161 (1987). J. C. Angus and C. C. Hayman, Science 241, 913 (1988). K. E. Spear, J. Am. Ceram. Soc. 72, 171 (1989). D. M. Gruen, Annu. Rev. Mater. Sci. 29, 211 (1999). L. C. Chen, P. D. Kichambare, K. H. Chen, J.-J. Wu, J. R. Yang, and S. T. Lin, J. Appl. Phys. 89, 753 (2001). D. M. Gruen, S. Liu, A. R. Krauss, J. Luo, and X. Pan, Appl. Phys. Lett. 64, 1502 (1994). D. M. Gruen, S. Liu, A. R. Krauss, and X. Pan, J. Appl. Phys. 75, 1758 (1995). D. Zhou, A. R. Krauss, L. C. Qin, T. G. McCauley, D. M. Gruen, T. D. Corrigan, R. P. H. Chang, and H. Gnaser, J. Appl. Phys. 82, 4546 (1997). R. Erz, W. Dotter, D. Jung, and H. Ehrhardt, Diamond Relat. Mater. 2, 449 (1993). V. I. Konov, A. A. Smolin, V. G. Ralchenko, S. M. Pimenov, E. D. Obraztsova, E. N. Loubnin, S. M. Metev, and G. Sepold, Diamond Relat. Mater. 4, 1073 (1995). V. I. Konov, E. D. Obraztsova, S. M. Pimenov, V. G. Ralchenko, and A. A. Smolin, Proc. SPIE 2428, 612 (1995). L. C. Nistor, J. Van Landuyt, V. G. Ralchenko, E. D. Obraztsova, and A. A. Smolin, Diamond Relat. Mater. 6, 159 (1997). L. C. Nistor, J. Van Landuyt, V. G. Ralchenko, E. D. Obraztsova, K. G. Korotushenko, and A. A. Smolin, Mater. Sci. Forum 239–241, 115 (1997). C. Z. Gu and X. Jiang, J. Appl. Phys. 88, 1788 (2000). T. Sharda, M. Umeno, T. Soga, and T. Jimbo, Appl. Phys. Lett. 77, 4304 (2000). T. Sharda, T. Soga, T. Jimbo, and M. Umeno, Diamond Relat. Mater. 9, 1331 (2000). B. Hong, J. Lee, R. W. Collins, Y. Kuang, W. Drawl, R. Messier, T. T. Tsong, and Y. E. Strausser, Diamond Relat. Mater. 6, 55 (1997). D. V. Fedoseev, V. L. Bukhovets, I. G. Varshavskaya, and Y. N. Tolmachev, Russ. J. Phys. Chem. 70, 1594 (1996). M. Zarrabian, N. Fourches-Coulon, G. Turban, C. Marhic, and M. Lancin, Appl. Phys. Lett. 70, 2535 (1997).

Nanocrystalline Diamond 20. G. Amaratunga, A. Putnis, K. Clay, and W. Milne, Appl. Phys. Lett. 55, 634 (1989). 21. D. Zhou, T. G. McCauley, L. C. Qin, A. R. Krauss, and D. M. Gruen, J. Appl. Phys. 83, 540 (1998). 22. W. Kratschmer, L. D. Lamb, K. Fostiorpoulas, and D. R. Huffman, Nature 347, 354 (1990). 23. D. M. Gruen, Nucl. Instrum. Methods Phys. Res. 78, 118 (1993). 24. D. A. Horner, L. A. Curtiss, and D. M. Gruen, Chem. Phys. Lett. 233, 243 (1995). 25. W. Banholzer, Surf. Coat. Technol. 53, 1 (1992). 26. L. S. Pan and D. R. Kania, “Diamond: Electronic Properties and Application.” Kluwer Academic, London, 1995. 27. D. R. McKenzie, D. A. Muller, and B. A. Paithorpe, Phys. Rev. Lett. 67, 773 (1991). 28. Y. Lifshitz, G. D. Lempert, and E. Grossman, Phys. Rev. Lett. 72, 2753 (1994). 29. P. J. Fallon, V. S. Veerasamy, C. A. Davis, J. Robertson, G. A. J. Amaratunga, W. I. Milne, and J. Koskinen, Phys. Rev. B 48, 4777 (1993). 30. M. Weiler, S. Sattel, K. Jung, H. Ehrhardt, V. S. Veerasamy, and J. Robertson, Appl. Phys. Lett. 64, 2797 (1994). 31. P. K. Bachman, D. Leers, and H. Lydtin, Diamond Relat. Mater. 1, 1 (1991). 32. S. Matsumoto, Y. Saito, M. Tsutsumi, and N. Setaka, J. Mater. Sci. 17, 3106 (1981). 33. A. Erdemir, G. R. Fenske, A. R. Krauss, D. M. Gruen, T. McCauley, and R. T. Csencsits, Surf. Coat. Technol. 120, 565 (1999). 34. A. Erdemir, C. Bindal, G. R. Fenske, C. Zurker, A. R. Krauss, and D. M. Gruen, Diamond Relat. Mater. 5, 923 (1999). 35. T. Lin, G. Y. Yu, A. T. S. Wee, Z. X. Shen, and K. P. Loh, Appl. Phys. Lett. 77, 2692 (2000). 36. M. Tsuda, M. Nakajima, and S. Oikawa, J. Am. Chem. Soc. 108, 5780 (1986). 37. R. Haubner and B. Lux, Diamond Relat. Mater. 2, 1277 (1993). 38. K. Chakrabarti, R. Chakrabarti, K. K. Chattopadhyay, S. Chaudhuri, and A. K. Pal, Diamond Relat. Mater. 7, 845 (1998). 39. M. W. Geis and M. A. Tamor, “Encyclopedia of Applied Physics,” Vol. 5. VCH, New York, 1993. 40. W. A. Yarbrough and R. Messier, Science 247, 605 (1990). 41. W. Zhu, B. R. Stoner, B. E. Williams, and J. T. Glass, Proc. IEEE 79, 621 (1991). 42. W. Zhu, C. A. Randall, A. R. Badzian, and R. Messier, J. Vac. Sci. Technol. A 7, 2315 (1989). 43. B. E. Williams, J. T. Glass, R. F. Davis, K. Kobashi, and Y. Kawate, in “Materials Research Society Extended Abstracts,” EA-15, (G. H. Johnson, A. R. Badzian, and M. W. Geis, Eds.). Materials Research Society, Pittsburgh, 1988. 44. Y. Cong, R. W. Collins, R. Messier, K. Vedam, G. F. Epps, and H. Windischmann, in, “New Diamond Science and Technology” (R. Messier, J. Glass, J. Butler, and R. Roy, Eds.), p. 735. Materials Research Society, Pittsburgh, 1991. 45. J. Robertson and E. P. O’Reilly, Phys. Rev. B 35, 2946 (1987). 46. Y. Sato and M. Kamo, Surf. Coat. Technol. 39/40, 183 (1989). 47. Y. Cong, I. An, H. V. Nguyen, K. Vedam, R. Messier, and R. W. Collins, Surf. Coat. Technol. 49, 381 (1991). 48. K. A. Snail and C. M. Marks, Appl. Phys. Lett. 60, 3135 (1992). 49. L. Fayette, B. Marcus, M. Mermoux, L. Abello, and G. Lucazeau, Diamond Relat. Mater. 3, 438 (1994). 50. C. D. Zulker, D. M. Gruen, and A. R. Krauss, Mater. Res. Soc. Bull. 20, 29 (1995). 51. B. Hong, M. Wakagi, R. W. Collins, I. An, N. C. Engdahl, W. Drawl, and R. Messier, Diamond Relat. Mater. 3, 431 (1994). 52. H. Windischmann, G. F. Epps, Y. Cong, and R. W. Collins, J. Appl. Phys. 69, 2231 (1991).

461 53. B. Hong, J. Lee, R. W. Collins, Y. Kuang, W. Drawl, R. Messier, T. T. Tsong, and Y. E. Strausser, Diamond Relat. Mater. 6, 55 (1997). 54. R. W. Collins and Y. T. Kim, Anal. Chem. 62, 887A (1990). 55. R. W. Collins, Rev. Sci. Instrum. 61, 2029 (1990). 56. B. Hong, M. Wakagi, W. Drawl, R. Messier, and R. W. Collins, Phys. Rev. Lett. 75, 1122 (1995). 57. J. Lee, B. Hong, R. Messier, and R. W. Collins, Appl. Phys. Lett. 69, 1716 (1996). 58. E. Kondoh, T. Ohta, T. Mitomo, and K. Ohtsuka, Appl. Phys. Lett. 59, 488 (1991). 59. W. Zhu, R. Messier, and A. R. Badzian, in “Proceeding of the First International Symposium on Diamond and Diamond-like Films” (J. P. Dismukes, Ed.), Electrochemical Society, Pennington, NJ, 1989. 60. C. Zuiker, A. R. Krauss, D. M. Gruen, X. Pan, J. C. Li, R. Scsencsits, A. Erdemir, C. Bindal, and G. Fenske, Thin Solid Films 270, 154 (1995). 61. K. Teii, H. Ito, M. Hori, T. Takeo, and T. Goto, J. Appl. Phys. 87, 4572 (2000). 62. R. W. B. Pearse and A. G. Gaydon, “The Identiﬁcation of Molecular Spectra,” 4th ed. Wiley, New York, 1976. 63. W. L. Wiese, M. W. Smith, and B. M. Miles, “Atomic Transition Probabilities,” Vols. I and II. National Standard Reference Data Series, Washington, DC, 1966. 64. H. Ito, M. Ikeda, M. Ito, M. Hori, T. Takeo, T. Kato, and T. Goto, Jpn. J. Appl. Phys. 2 36, L880 (1997). 65. D. Robinson and P. D. Lenn, Appl. Opt. 6, 983 (1967). 66. K. Teii, J. Vac. Sci. Technol. A 17, 138 (1999). 67. S. Iijima, Y. Aikawa, and K. Baba, J. Mater. Res. 6, 1491 (1991). 68. M. Avalos-Borja, G. A. Hiraata, O. Contreras, X. G. Ning, A. Duarte-Moller, and A. Barna, Diamond Rel. Mater. 5, 1249 (1996). 69. J. Singh, J. Mater Sci. 29, 2761 (1994). 70. J. S. Luo, D. M. Gruen, A. R. Krauss, X. Z. Pan, and S. Z. Liu, in “Recent Advances in the Chemistry and Physics of Fullerenes and Related Materials” (R. S. Ruoff and K. M. Kadish, Eds.). SPIE, Bellingham, WA, 1995. 71. R. Csencsits, C. Zurker, D. M. Gruen, and A. R. Krauss, Solid State Phenom. 261, 51 (1996). 72. L. C. Qin and L. W. Hobbs, Mater. Res. Soc. Symp. Proc. 373, 341 (1995). 73. P. Keblinski, D. Wolf, S. R. Phillpot, and H. Gleiter, J. Mater. Res. 13, 2077 (1998). 74. H. Hirai, M. Terauchi, M. Tanaka, and K. Kondo, Diamond Relat. Mater. 8, 1703 (1999). 75. R. Kuzuo, M. Terauchi, M. Tanaka, Y. Saito, and H. Shinohara, Jpn. J. Appl. Phys. 30, 1817 (1991). 76. R. Kuzuo, M. Terauchi, and M. Tanaka, Jpn. J. Appl. Phys. 31, 1484 (1991). 77. R. F. Egerton and M. J. Whelan, Philos. Mag. 309, 739 (1974). 78. R. F. Egerton, “Electron Energy Loss Spectroscopy in the Electron Microscope.” Plenum, New York, 1966. 79. R. A. Roberts and W. C. Walker, Phys. Rev. 161, 7730 (1967). 80. T. Xu, S. Yang, J. Lu, Q. Xue, J. Li, W. Guo, and Y. Sun, Diamond Relat. Mater. 10, 1441 (2001). 81. B. Fabre, F. Kanouﬁ, and J. Simonet, J. Electroanal. Chem. 434, 225 (1993). 82. F. Fusallba, N. El. Mehdi, L. Breau, and D. Belanger, Chem. Mater. 11, 2743 (1999). 83. S. C. Ng, P. Fu, W. L. Yu, H. S. O. Chan, and K. L. Tan, Synth. Met. 87, 119 (1997). 84. S. P. Lawrence and R. K. Don, “Diamond: Electronic Properties and Applications.” Kluwer Academic, Boston, 1995. 85. S. E. Grillo and J. E. Field, J. Phys. D 33, 595 (2000).

462 86. G. Ilangovan and K. Chandrasekara Pillai, Langmuir 13, 566 (1997). 87. T. Jiang and K. Xu, Carbon 12, 1663 (1995). 88. S. Prawer, K. W. Nugent, D. N. Jamieson, J. O. Orwa, L. A. Bursill, and J. L. Peng, Chem. Phys. Lett. 332, 93 (2000). 89. J. Wagner, M. Ramsteiner, Ch. Wild, and P. Koidl, Phys. Rev. B 40, 1817 (1989). 90. R. J. Nemanich, J. T. Glass, G. Lucovsky, and R. E. Shroder, J. Vac. Sci. Technol. A 6, 1783 (1988). 91. V. I. Merkulov, J. S. Lannin, C. H. Munro, S. A. Asher, V. S. Veerasamy, and W. I. Milne, Phys. Rev. Lett. 78, 4869 (1997). 92. S. M. Leeds, T. J. Davis, P. W. May, C. D. O. Pickard, and M. N. R. Ashfold, Diamond Relat. Mater. 7, 233 (1998). 93. T. Sharda, T. Soga, T. Jimbo, and M. Umeno, Diamond Relat. Mater. 10, 1592 (2001). 94. V. Paillard, P. Melinon, V. Dupuis, J. P. Perez, and A. Perez, Phys. Rev. Lett. 71, 4170 (1993). 95. V. Paillard, P. Melinon, V. Dupuis, A. Perez, J. P. Perez, G. Guiraud, J. Fornazero, and G. Panczer, Phys. Rev. B 49, 11433 (1994). 96. J. Lee, R. W. Collins, R. Messier, and Y. E. Strausser, Appl. Phys. Lett. 70, 1527 (1997). 97. S. A. Catledge and Y. K. Vohra, J. Appl. Phys. 84, 6469 (1998). 98. N. Wada and S. A. Solin, Physica B 105, 353 (1981). 99. H. Windischmann, G. F. Epps, Y. Cong, and R. W. Collins, J. Appl. Phys. 69, 2231 (1991). 100. K. K. Chattopadhay and S. Matsumoto, Appl. Phys. Lett. 67, 3972 (1995). 101. H. Yoshikawa, C. Morel, and Y. Koga, Diamond Relat. Mater. 10, 1588 (2001). 102. V. F. Dorfman, Thin Solid Films 212, 267 (1992). 103. D. M. Bhusari, J. R. Yang, T. Y. Wang, S. T. Lin, K. H. Chen, and L. C. Chen, Solid State Commun. 107, 301 (1998). 104. F. J. Himpsel, J. A. Knapp, J. A. VanVechten, and D. E. Eastman, Phys. Rev. B 20, 624 (1979). 105. T. P. Ong and R. P. H. Chang, Appl. Phys. Lett. 55, 2063 (1989). 106. R. L. C. Wu, A. K. Rai, A. Garscadden, P. Lee, H. D. Desai, and K. Miyoshi, J. Appl. Phys. 72, 110 (1992). 107. C. Friedrich, G. Berg, E. Broszeit, and C. Berger, Thin Solid Films 290–291, 216 (1996). 108. K. K. Chattopadhyay, J. Dutta, S. Chaudhuri, and A. K. Pal, Diamond Relat. Mater. 4, 122 (1995). 109. A. B. Maity, M. Basu, S. Chaudhuri, and A. K. Pak, J. Phys. D 28, 2547 (1995). 110. T. Sharda, T. Soga, T. Jimbo, and M. Umeno, J. Nanosci. Nanotech. 1, 287 (2001). 111. D. R. Tallant, J. E. Parmeter, M. P. Siegal, and R. L. Simpson, Diamond Relat. Mater. 4, 191 (1995). 112. T. A. Friedmann, K. F. McCarty, J. C. Barbour, M. P. Siegal, and D. C. Dibble, Appl. Phys. Lett. 68, 1643 (1996). 113. C. Moelle, M. Werner, F. Szucs, D. Wittorf, M. Sellschopp, J. von Borany, H. J. Fecht, and C. Johnston, Diamond Relat. Mater. 7, 499 (1998). 114. I. Bahrin, in “Thermochemical Data of Pure Substances,” Part I. VCH, Weinheim, 1993. 115. C. Moelle, F. Szucs, H. J. Fecht, P. Rohrs, C. Johnston, P. R. Cholker, and M. Werner, in “Proc. of the Micromaterials Conference.” Berlin, 1997. 116. P. K. Bachmann, D. Leers, and D. U. Wiechert, Diamond Relat. Mater. 2, 683 (1993). 117. Y. Muto, T. Sugino, K. Kobashi, and J. Shirafuji, J. Appl. Phys. 59, 843 (1991). 118. S. N. Kundu, M. Basu, A. B. Maity, S. Chaudhuri, and A. K. Pal, Mater. Lett. 31, 303 (1997). 119. R. Chakrabarti, K. Chakrabarti, A. B. Maity, S. Chaudhuri, and A. K. Pal, Diamond Relat. Mater. 6, 991 (1997).

Nanocrystalline Diamond 120. J. Seth, R. Padiyath, S. V. Babu, and M. David, Thin Solid Films 212, 251 (1992). 121. E. G. Spencer, P. H. Schmidt, D. C. Joy, and F. J. Sansalane, Appl. Phys. Lett. 29, 118 (1976). 122. S. Bhattacharya, O. Auciello, J. A. Carlisle, L. A. Curtiss, A. N. Goyette, D. M. Gruen, A. R. Krauss, J. Schlueter, A. Sumant, and P. Zapol, Appl. Phys. Lett. 79, 1441 (2001). 123. J. Robertson and C. A. Davis, Diamond Relat. Mater. 4, 441 (1995). 124. T. Nishimori, K. Nakano, H. Sakamoto, Y. Takakuwa, and S. Kono, Appl. Phys. Lett. 71, 945 (1997). 125. K. Okano, S. Koizumi, S. R. P. Silva, and G. A. J. Amaratunga, Nature (London) 381, 140 (1996). 126. C. A. Spindt, I. Brodie, L. Humphrey, and E. R. Westerberg, J. Appl. Phys. 47, 5248 (1976). 127. N. S. Xu, R. V. Latham, and Y. Tzeng, Electron Lett. 29, 1596 (1993). 128. C. Wang, A. Garcia, D. C. Ingram, M. Lake, and M. E. Kordesch, Electron Lett. 27, 1459 (1991). 129. G. A. J. Amaratunga and S. R. P. Silva, Appl. Phys. Lett. 68, 2529 (1996). 130. O. Groning, O. M. Kuttel, P. Groning, and L. Schlapbach, Appl. Surf. Sci. 111, 135 (1997). 131. M. W. Geis, J. C. Twichell, J. Macaulay, and K. Okano, Philos. Appl. Phys. Lett. 67, 1328 (1995). 132. B. F. Coll, J. E. Jaskie, J. L. Markham, E. P. Menu, A. A. Talin, and P. Von Allmen, Mater. Res. Soc. Symp. Proc. 498, 185 (1998). 133. A. G. Rinzler, J. H. Hafner, P. Nikolaev, L. Lou, S. G. Kim, D. Tomanek, P. Nordlander, D. T. Colbert, and R. E. Smalley, Science 269, 1550 (1995). 134. W. Zhu, G. P. Kochanski, S. Jin, and L. Seibles, J. Appl. Phys. 78, 2707 (1995). 135. V. V. Zhirnov, E. I. Givargizov, and P. S. Plekhanov, J. Vac. Sci. Technol. B 13, 418 (1995). 136. R. Schlesser, M. T. McClure, B. L. McCarson, and Z. Sitar, J. Appl. Phys. 82, 5763 (1997). 137. F. Lacher, C. Wild, D. Behr, and P. Koidl, Diamond Relat. Mater. 6, 1111 (1997). 138. P. D. Kichambare, F. G. Tarntair, L. C. Chen, K. H. Chen, and H. C. Cheng, J. Vac. Sci. Technol. B 18, 2722 (2000). 139. P. D. Kichambare, Center for Applied Energy Research, The University of Kentucky, Lexington, KY 40511, unpublished work, 2001. 140. R. H. Fowler and L. W. Nordheim, Proc. Roy. Soc. London Ser. A 119, 173 (1928). 141. K. Wu, E. G. Wang, Z. X. Cao, Z. L. Wang, and X. Jiang, J. Appl. Phys. 88, 2967 (2000). 142. M. W. Geis, J. C. Twichell, and T. M. Lyszczarz, J. Vac. Sci. Technol. B 14, 2060 (1996). 143. K. Wu, E. G. Wang, J. Chen, and N. S. Xu, J. Vac. Sci. Technol. B 17, 1059 (1999). 144. Y. V. Pleskov, A. Y. Sakharova, M. D. Krotova, L. L. Bouilov, and B. V. Spitsyn, J. Electroanal. Chem. 228, 19 (1987). 145. M. Iwaki, S. Sato, K. Takahashi, and H. Sakairi, Nucl. Instrum. Methods 209/210, 1129 (1983). 146. J. Xu, M. C. Granger, Q. Chen, T. E. Lister, J. W. Strojek, and G. M. Swain, Anal. Chem. 69, 591A (1997). 147. R. Tenne and C. Levy-Clement, Israel J. Chem. 38, 57 (1998). 148. J. Xu and G. M. Swain, Anal. Chem. 71, 4603 (1999). 149. B. Fausett, M. C. Granger, M. L. Hupert, J. Wang, G. M. Swain, and D. M. Gruen, Electroanalysis 12, 7 (2000). 150. J. Angus, in “Synthetic Diamond: Emerging CVD Science and Technology” (K. E. Spear and J. P. Dismukes, Eds.). Wiley, New York, 1994.

Nanocrystalline Diamond 151. “Status and Applications of Diamond and Diamond-like Materials: An Emerging Technology.” National Materials Advisory Board, Washington, DC, 1990. 152. W. J. P. van Enckevort, J. Hard Mater. 1, 247 (1990). 153. S. J. Bull, Diamond Relat. Mater. 4, 827 (1995). 154. D. G. Bhat, D. G. Johnson, A. P. Malshe, H. Naseem, W. D. Brown, L. W. Schaper, and C. H. Shen, Diamond Relat. Mater. 4, 921 (1995). 155. I. P. Hayward, Surf. Coat. Technol. 49, 554 (1991). 156. A. K. Gangopadhyay and M. A. Tamor, Wear 169, 221 (1991). 157. S. K. Choi, D. Y. Jung, S. Y. Kweon, and S. K. Jung, Thin Solid Films 279, 110 (1996). 158. B. E. Williams and J. T. Glass, J. Mater. Res. 4, 373 (1989). 159. S. J. Bull and A. Mathews, Diamond Relat. Mater. 1, 1049 (1992). 160. K. V. Ravi, in “Synthetic Diamond: Emerging CVD Science and Technology,” Electrochem. Soc. Monograph. Wiley, New York, 1994. 161. S. Hogmark, O. Hollman, A. Alahelisten, and P. Hedenqvist, Wear 200, 235 (1996). 162. A. Erdemir, G. R. Fenske, A. R. Krauss, D. M. Gruen, T. McCauley, and R. T. Csencsits, Surf. Coat. Technol. 120–121, 565 (1999). 163. A. Erdemir, M. Halter, G. R. Fenske, A. Krauss, D. M. Gruen, S. M. Pimenov, and V. I. Konov, Surf. Coat. Technol. 94–96, 537 (1997). 164. I. P. Hayward, I. L. Singer, and L. E. Seitzman, Wear 157, 215 (1992). 165. M. N. Gardos and K. V. Ravi, Diamond Films Technol. 4, 139 (1994). 166. M. N. Gardos and K. V. Ravi, J. Mater. Res. 5, 2599 (1990). 167. S. V. Pepper, J. Vac. Sci. Technol. 20, 643 (1982). 168. S. Chandrasekar and B. Bhushan, Wear 153, 79 (1992). 169. P. Hollman, O. Wanstrand, and S. Hogmark, Diamond Rel. Mater. 7, 1471 (1998). 170. S. Veprek, J. Vac. Sci. Technol. A 17, 2401 (1999). 171. A. Kelly and N. H. Macmillan, “Strong Solids,” 3rd ed. Clarendon, Oxford, 1986. 172. R. W. Hertzberg, “Deformation and Fracture Mechanics of Engineering Materials,” 3rd ed. Wiley, New York, 1989. 173. S. Veprek, S. Reiprich, and Li Shizhi, Appl. Phys. Lett. 66, 2640 (1995). 174. S. Veprek, Thin Solid Films 317, 449 (1998). 175. P. Nesladek and S. Veprek, Phys. Status Solidi A 177, 53 (2000). 176. S. Veprek, P. Nesladek, A. Niederhofer, H. Mannling, and M. Jilek, TMS Annu. Meeting, San Diego, 28 February–4 March, 1999, in “Surf. Eng. Science and Technology I” (A. Kumar, Y. W. Chung, J. J. Moore, and J. E. Smugeresky, Eds.) (1999). 177. A. Niederhofer, P. Nesladek, H. D. Mannling, K. Moto, S. Veprek, and M. Jilek, Surf. Coat. Technol. 120/121, 173 (1999). 178. J. E. Carsley, J. Ning, W. W. Milligan, S. A. Hackney, and E. C. Aifantis, Nanostruct. Mater. 5, 441 (1995). 179. R. W. Siegel and G. E. Fougere, Nanostruct. Mater. 6, 205 (1995). 180. E. Arzt, Acta Mater. 46, 5611 (1998). 181. J. Schiotz, F. D. Di Tolla, and K. W. Jacobsen, Nature 391, 561 (1998). 182. R. Thompson, Y. Vohra and M. Koopman, Surface engineering: Coatings and heat treatments in “Proc. of the 1st ASM International Surface Engineering and the 13th IFHTSE Congress and Exposition” (O. O. Popoola, N. B. Dahotre, S. J. Midea, and H. Kopech, Eds.), 7–10 October 2002. Columbus, OH, 2003. 183. N. Toprani, S. A. Catledge, Y. K. Vohra, and R. Thompson, J. Mater. Res. 15, 1052 (2000). 184. S. A. Catledge and Y. K. Vohra, J. Appl. Phys. 86, 698 (1999).

463 185. T. Ramanthan, K. Jagannadhan, K. Mallika, and R. Komanduri, Philos. Mag. B (1999). 186. J. W. Ager III and M. D. Drory, Phys. Rev. B 48, 2601 (1993). 187. Y. Muranaka, H. Yamashita, and H. Miyadera, J. Vac. Sci. Technol. A 9, 76 (1991). 188. Y. Liou, A. Inspektor, R. Weimer, D. Knight, and R. Messier, J. Mater. Res. 5, 2305 (1990). 189. Y. Muranaka, H. Yamashita, and H. Miyadera, J. Appl. Phys. 69, 8145 (1991). 190. C. E. Johnson, W. A. Weimer, and F. M. Cerio, J. Mater. Res. 7, 1427 (1992). 191. C. J. Chu, R. H. Hauge, J. L. Margrave, and M. P. D’Evelyn, Appl. Phys. Lett. 61, 1393 (1992). 192. C. Nesladek, J. Spinnewin, R. Lebout, and R. Lorent, Diamond Relat. Mater. 3, 912 (1994). 193. C. Campbell, “Surface Acoustic Wave Devices and their Signal Processing Applications.” Academic Press, New York, 1989. 194. K. Yamanouchi, N. Sakumi, and T. Satoh, Proc. IEEE 351 (1989). 195. H. Nakahata, A. Hachigo, S. Shikata, and N. Fujimori, Proc. IEEE 361 (1995). 196. S. Fujii, Y. Seki, and K. Yoshida, Proc. IEEE 183 (1997). 197. K. L. Dreifus, J. J. Higgins, R. B. Henard, R. Almar, and L. P. Solie, Proc. IEEE 191 (1997). 198. H. Nakahata, H. Kitabayashi, and A. Hachigo, Proc. IEEE 319 (1998). 199. A. Hachigo, H. Nakahata, K. Itakura, S. Fujii, and S. Shikata, Proc. IEEE 325 (1995). 200. B. Bi, W. S. Huang, J. Asmussen, and B. Golding, Diamond Relat. Mater. 11, 677 (2002). 201. L. E. Cross and S. Trolier-McKinstry, Encl. Appl. Phys. 21, 429 (1997). 202. D. L. Polla and L. F. Francis, Mater. Res. Soc. Bull. 21, 59 (1996). 203. J. J. Sniegowski, in “Tribology Issues and Opportunities in MEMS” (B. Bhushan, Ed.), Kluwer Scientiﬁc, Dordrecht, 1998. 204. A. P. Lee, A. P. Pisano, and M. G. Lin, Mater. Res. Soc. Symp. Proc. 276, 67 (1992). 205. M. N. Gardos, in “Tribology Issues and Opportunities in MEMS” (B. Bushan, Ed.). Kluwer Scientiﬁc, Dordrecht, 1998. 206. M. N. Gardos, Surf. Coat. Technol. 113, 183 (1999). 207. A. R. Krauss, O. Auciello, D. M. Gruen, A. Jayatissa, A. Sumant, J. Tucek, D. C. Mancini, N. Moldovan, A. Erdemir, D. Ersoy, M. N. Gardos, H. G. Busmann, E. M. Meyer, and M. Q. Ding, Diamond Relat. Mater. 10, 1952 (2001). 208. J. L. Davidson, R. Ramesham, and C. Ellis, J. Electrochem. Soc. 137, 3206 (1990). 209. M. Aslam, G. S. Yang, and A. Masood, Sensors Actuators A 45, 131 (1994). 210. D. R. Wur, J. L. Davidson, W. P. Kang, and D. L. Kinser, J. Micromech. Syst. 4, 34 (1994). 211. O. Dorsch, K. Holzner, and M. Werner, Diamond Relat. Mater. 2, 1096 (1993). 212. G. Zaho, E. M. Charlson, and E. J. Charson, J. Appl. Phys. 73, 1832 (1993). 213. S. Moller, E. Obermeir, and J. Lin, Sensors Actuators B 25, 343 (1995). 214. E. I. Givargizov, V. V. Zhirnov, A. V. Kuznetsov, and P. S. Plekhanov, Mater. Lett. 18, 61 (1993). 215. P. K. Bachmann and H. J. Hagemann, Diamond Relat. Mater. 4, 820 (1995). 216. C. F. Chen, S. H. Chen, T. M. Hong, and M. H. Tsai, J. Appl. Phys. 77, 941 (1995). 217. M. D. Irwin, C. G. Pantano, P. Gluche, and E. Kohn, Appl. Phys. Lett. 71, 716 (1997).

Plasma Chemical Vapor Deposition of Nanocrystalline Diamond

Nanocrystalline Phosphors

Nanocrystalline Materials

Nanocrystalline apatite-based biomaterials

Nanocrystalline TiO2 for Photocatalysis

Diamond

Nanocrystalline Silicon Superlattices

Nanocrystalline Aluminum Alloys

Diamond

Diamond

Diamond

Diamond

Diamond

Diamond Cut Diamond

Diamond

Diamond

Diamond

Diamond

Diamond

Diamond

Diamond

Diamond

Nanocrystalline Ceramics by Mechanical Activation

Nanocrystalline and Amorphous Magnetic Microwires

Nanocrystalline Silicon: Electron Spin Resonance

Diamond Mask

Diamond Solitaire

Diamond Dove

Diamond Dogs

The Diamond

Nanocrystalline Diamond

Plasma Chemical Vapor Deposition of Nanocrystalline Diamond

Nanocrystalline Phosphors

Nanocrystalline Materials

Nanocrystalline apatite-based biomaterials

Nanocrystalline TiO2 for Photocatalysis

Diamond

Nanocrystalline Silicon Superlattices

Nanocrystalline Aluminum Alloys

Diamond

Diamond

Diamond

Diamond

Diamond

Diamond Cut Diamond

Diamond

Diamond

Diamond

Diamond

Diamond

Diamond

Diamond

Diamond

Nanocrystalline Ceramics by Mechanical Activation

Nanocrystalline and Amorphous Magnetic Microwires

Nanocrystalline Silicon: Electron Spin Resonance

Diamond Mask

Diamond Solitaire

Diamond Dove

Diamond Dogs

The Diamond

Recommend Documents